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By Roland W. Rammler

With oil consumption increasing more rapidly than known reserves in industrial countries, there has been a growing interest in the production of synthetic crude oil. The main starting materials for such production are coal, oil shale and tar sand, all of which occur in enormous quantities in North America. A process sequence suitable for all three of these starting materials incorporates a preliminary retorting step to produce oils for subsequent hydrotreating. For this retorting step, a technique known as the Lurgi-Ruhrgas process has proven highly efficient. The process has been commercially used for the flash carbonization of coal and has also been tested in pilot- plant experiments with oil shale and tar sand.

Jan 1, 1970

By A. E. Dwight, Paul A. Beck

Known CsCl-type phases in transition metal alloys are shown to be formed preferentially at an electron concentration of approximately 6; a tetragonally distorted version of the CsCl structure is stable at -e/a = 6.5. The relationship of the distortion to electron concentration is verified by experimental data for the quasibinary systems Time ,Os) , Ti(Re,Ir), and Ti(Re,Pt) and the ternary system Ti-Ru-Rh. It was noted some years ago1j2 that the occurrence of CsC1-type ordered structures in equiatomic binary alloys, AB, of transition metals correlates with the location of the component elements in the periodic table. One of the components, A, must be to the left and the other component. B, must be to the right of the chromium group. It was pointed out2 that the CsCl structure does not occur in CrFe. is metastable in VFe,3 and is stable in TiFe up to high temperatures, and that this progressively increasing stability may correspond to an increase in electron transfer between the two components and to a correlated decrease of the magnetic moment associated with the iron atoms. Such a decrease was in fact later demonstrated by Nevitt. 4 The binary CsC1-type phases that are formed by a transition metal with either another transition metal or a lanthanide are listed in Table I. Thirty-three of these phases occur at an average electron concentration (average number of electrons outside of closed shells) of e/a = 6. The five phases that are formed at an electron concentration of e/a = 5.5 appear to be borderline cases. This is suggested for instance by the fact that TiRe is bcc. without any long-range order, while TiTc (at the same electron concentration) has a CsCl ordered structure. Six of the undistorted CsCl phases formed by transition metals with one another have an average electron concentration of e/a = 6. 5. Table I. With the exception of Zr4,Rhs4 and Hf48Rh52. the B component in these phases is a first long-period transition metal, i.e.. iron or cobalt. TiNi. with an electron concentration of e/n = 7, appears to undergo transformation at around room temperature to a more complex ordered struct~re.~ The five other CsCl structures known to occur at e/a = 6.5 have either scandium or a lanthanide metal as the A component; the difference in the elec-tronegativity and in the e/a ratio between the component elements is in these phases very large, A(e/a) - 7, and they may be expected to have a rather strong ionic bond component. Presumably the ionic bond stabilizes the CsCl-type structure at the equiatomic composition and it overrides the electron concentration requirement. When the B element is not from the first long period and the A component is not a rare earth metal. equiatomic (AB) alloys of transition metals with an average electron concentration of e/a = 6.5 in most cases have a crystal structure closely related to the CsC1-type, but slightly distorted tetragonally, Table I. This was first found6 for VRu, NbRu. and TaRu. It was noted6 that the undistorted CsCl structure is also stable in the three systems. but at a lower electron concentration, near e/a = 6. In these systems this requires a deviation from equiatomic stoichiometry. which is otherwise ideally suited for the CsC1-type structure. A similar situation was later found in the Ti-Rh7 and the Ti-Ir systems. 8 The Ti-Rh alloys undergo a second distortion and become orthorhombic when the rhodium content becomes somewhat larger than 50 pct. As mentioned previously. TiRe does not have the CsC1-type ordered structure. In order to investigate how close TiRe is to the borderline of stability of the

Jan 1, 1970

By Norman Weiss

There was a time in the early history of man's use of metals when ore could be directly reduced to usable form by the use of fire. In those days the term "milling," or its many synonyms signifying the concentration of an ore's metallic con- tents before smelting, was little known. Centuries later, even before Agricola, or even before the Greek and Roman civilizations, simple means were found to concentrate metallic minerals that possessed special characteristics such as high specific gravity, hardness, magnetic response, or simple appearance. In this way, the art of milling became valuable to civilization because it eliminated most of the cost and labor of transporting and melting the worthless nonmetallic part of an ore. As civilization became more dependent upon metals and solid fuels, the art of ore dressing (one of our synonyms) increased in importance and complexity, until today it has become a gigantic industry firmly based upon a mature but rapidly advancing technology. Technically trained people are needed in growing numbers to serve this industry.

Jan 1, 1970

By C. R. Whitney, G. N. Maniar, D. R. Muzyka

Previous results have shown that Pyromet 860, an Fe-Ni-base heat-resistant alloy, is stable at temperatures as high as 1500°F for aging times as long as 100 hr. This Paper describes the results of long-time creep-rupture testing at 1050" to 1400°F at various stress levels. Times as long as 37,660 hr were employed. The effects of time, temperature, and stress on the precipitates and their morphologies were studied by optical and electron microscopy, X-ray and electron diffraction, and microprobe techniques. phase, containing cobalt, nickel, and molybdenum, was detected after extended exposures from 1200" to 1400°F and careful study was performed to describe the kinetics of its formation in this alloy. µ phase formation apparently has little effect on the elevated-tem-perature properties of Pyromet 860. For times as long as 500 hr at 1300°F and below, with µ phase present, m significant effects on ambient temperature properties were noted. For longer times at 1300°F and after 1400°F exposure, the effects of u phase on ambient temperature tensile strength properties are not clear due to y' effects and grain boundary reactions. Electron-vacancy, N,, numbers were calculated using different methods described in literature and correlated with the present findings. In the selection of alloys for use in gas turbine applications, structural stability ranks as a primary criterion. High-temperature strength and cost are also of major concern. With these factors in mind, Pyromet 860 alloy, an Fe-Ni-base superalloy was designed. This alloy combines the cost advantages of Fe-Ni-base alloys such as A-286, 901, and V-57 with improved strength and structural stability'1,2 and no tendency to form the embrittling cellular 77 phase. A previous study3 reported on the stability of Pyro-met 860 at temperatures from 1375" to 157 5°F and times up to 100 hr. That study showed that the y' precipitates increased in size and separation and decreased in number with an increase in time or aging temperature. No deleterious phases were found to occur. In the present work, samples from four production heats were subjected to long-time creep-rupture testing at 1050" to 1400°F at various stress levels. Various heat treatments were used on the starting samples and tests were run up to 37,660 hr. The effects of time, temperature, and stress on the precipitates and their morphologies were studied by optical and electron microscopy, X-ray and electron diffrac- tion, and microprobe techniques. Electron vacancy numbers, Nv , calculations were made by TRW.4 Experimental results are correlated with the Nv data used to predict occurrence of intermetallic phases such as a phase. EXPERIMENTAL PROCEDURE Mechanical Tests. Material for the present study came from four production size heats of Pyromet 860 alloy, weighing from about 3000 to about 10,000 lb. All of these heats were made by vacuum induction melting plus consumable electrode vacuum remelting. The nominal analysis for this alloy is compared with the actual analysis of the four heats in Table I. Sections of these heats were forged to 9/16-in. round bar,3/4-in. square bar, 3-in. round bar, 4-in. square bar, and a gas turbine blade forging about 16 in, long, about 6 in. wide, and weighing about 20 lb. In general, all forging of this alloy is done from a 2050°F furnace temperature. Longitudinal test blanks were cut from the centers of the smaller bars, from mid-radius positions for the 3- and 4-in. bars, and from the air foil of the gas turbine blade and heat-treated according to the procedures outlined in Table 11. Heat treatment A is the "standard treatment" recommended for this alloy for best all-around strength and ductility. Heat treatment B is a modification of treatment A for improved tensile strength at moderate temperatures. The treatment coded C was designed for treating large sections according to a procedure previously described.' Heat treatment D was developed to yield optimum stress relaxation characteristics at 1050°F for a steam turbine bolting application. After heat treatment, the test blanks were machined either to plain bar creep specimens with a gage diameter of 0.252 in., to combination smooth-notched stress-rupture bars with a plain bar diameter of 0.178 in. and a concentration factor of Kt 3.8' at the notched section, or to notch-only specimens. All specimens conformed to ASTM requirements. Metallography. Most of the creep-rupture tests were continued to failure. A few bars were fractured as smooth or notch tensiles after creep-rupture exposures. After fracturing, ordinary metallographic sections were made primarily in gage areas adjacent to fractures to represent a "high-stress" region and through specimen threads to represent a "low-stress" region. All metallographic sections were made in a longitudinal direction with respect to the test specimen axes. For optical microscopy, the samples were etched in glyceregia (15 ml HC1, 5 ml HNO,, 10 ml glycerol). For XRD analysis, the phases were extracted electrolytically in two media: 20 pct &Po4 in H20 for selective extraction of y' and 10 pct HC1 in methanol for carbides and other phases.

Jan 1, 1970

By J. R. Cahoon, H. W. Paxton

The mechanical properties of unidirectionally solidified A1(rich)-Mg and A1(rich)-Cu castings containing up to 15 wt pct solute have been determined with re -spect to the volume fraction of interdendritic eutectic. Pioperties were determined in the directions pumllel and Perpendicular to that of solidification; the volume fraction of eutectic was varied between the "as-cast" and equilibrizcm amounts by approperiate heat treatment following solidification. The principles of fiber strengthened composites and dispersion strengthened materials are adapted to explain the mechanical properties of these castings. It is generally accepted that castings often have inferior mechanical properties when con~pared to wrought products. However, there is little quantitative data available concerning the factors which make apparently sound castings weak and/or brittle. The relative ease and inexpensiveness of the casting process have always been attractive and, therefore, an understanding of the factors which contribute to the mechanical properties of castings would seem desirable. Such an understanding may lead to an improvement in the mechanical properties to an extent where castings would become competitive in applications where presently only wrought products are considered to have the requisite properties. Such an understanding could also improve the reliability of present cast products. Much of the recent research on castings has centered about determining the extent of segregation in cast alloys. Macrosegregation, particularly inverse segregation, has been studied in some detail 1-8 and a considerable understanding of microsegregation has been obtained.9'10 The effect of solidification rate on dendrite spacing and on the amount of interdendritic eutectic in binary alloys has been established, particularly for Al(rich)-Cu alloys.""0 However, the extension of these ideas to relate the amount of interdendritic eutectic, concentration gradients, micro-segregation, dendrite spacings, and so forth, to the rnechanical properties has been limited. Dean and spear" have related the mechanical properties of an Al-Si-Mg alloy, A356-T62, to the dendrite spacing and have shown that the mechanical properties improve with decreasing dendrite spacing. Passmore et al.12 have shown that annealing at high temperature improves the mechanical properties of Al(rich)-Cu al- loys and Archer and Kempf 13 have shown that an Al-1 pct Mg-1.75 pct Si alloy behaves in a similar manner. Ahearn and Quigley 14 have shown that high temperature homogenization also enhances the mechanical properties of an SAE 4330 steel. However, in the above investigations, no underlying reasons were suggested for the improvement in mechanical properties. The purpose of the present investigation is to relate the mechanical properties of castings to some of the solichfication variables and to derive some equations by which calculations of the mechanical properties may be attempted. In particular, the effect of the amount of interdendritic eutectic and the effect of stress direction with respect to that of solidification on the mechanical properties will be considered. The Al(rich)-Mg and Al(rich)-Cu binary alloy systems were chosen for study. The A1-Mg system was chosen because its constitutional relationships are such that large volunles of eutectic (up to 24 vol pct) may be obtained in the as-cast condition and then be completely dissolved by subsequent heat treatment at about 440°C. This allows a comprehensive study relating the mechanical properties of castings to the amount of interdendritic eutectic. Also the Al(rich)-Mg eutectic is almost a single phase 15 which should make the experimental results more amenable to theoretical interpretation and calculation. The A1-Cu system was chosen for study because of the large amount of related information available concerning segregation, dendrite spacing, and so forth. Unidirectionally solidified castings were used throughout the investigation so that the effect of solidification direction with respect to the direction of applied stress could be determined. THEORETICAL It is well known that upon solidification of binary alloy castings, the nonequilibrium amount of eutectic which forms is given by 10 where fe o is the weight fraction of eutectic, Cs is the solid solubility of solute at the eutectic temperature, k is the equilibrium partition coefficient, and C, is the average composition of the alloy. In the development of Eq. [I], it is assumed that the effects of inverse segregation and diffusion in the solid are negligible, and that no porosity is present. If the casting is homogenized at a high temperature for a long period of time, some (or all) of the eutectic is dissolved and the amount of eutectic for this "equilibrium" condition may be calculated directly from the constitutional diagram. By appropriate intermediate annealing, the

Jan 1, 1970

By Willard L. Hunter

Four solvent extraction systems were studied to determine their efficiency jor extraction and separation of tantalum and columbium. Aqueous feed solutions of varying HF-HCl concentrations and metal content were contacted with equal volumes of cyclohexanone, 3-methyl-2-butanone, and 2-pentanone and solutions of varying HF-H2S04 concentrations were contacted with equal volumes of 2-pentanone. One multistage continuous test was made in a polyethylene pulse column using cy clohexanone as the organic phase. In each system studied, columbium and tantalum purities in excess of 95 pct with respect to each other were obtained in single-stage tests at low acidities in the feed solution. Separation factors ranging from 1700 to 2400 were obtained when rising HF-HCl mixtures in the aqueous phase. Best results were obtained when a solution of HF-H2S04 was used as the aqueous phase and 2-pentanone as the organic phase. A separation factor in excess of 6000 was obtained in one stage with aqueous solution concentrations of 2 _N HF and 2N H2S0,. When acid concentrations were increaszd to 52 HF and 10 _N H2S0,, 99.9 pct of the tantalum and 98.2 pct of the columbium initially present in the feed solution were transferred to the organic phase. The separation of columnbium and tantalum obtainable by means of the solvent extraction systems presented in this paper was found to corn -pare favorably with other systems, including the HF-H2SO4-methyl isobutyl ketone system currently used by most producers for the extraction and separation of these metals. TANTALUM and columbium are always found together in minerals of commercial significance, although the proportion of the two metals in ores varies within broad limits. Columbium is estimated to be 13 times more abundant than tantalum. Five methods generally employed for the separat:ion of these metals are: 1) fractional crystallization (the Marignac process),2 2) solvent separation, 3) fractional distillation of their chlorides, 4) ion exchange, and 5) selective reduction. Of these methods, the one currently used by industry to the greatest extent is that of solvent separation. One of the early technical developments in solvent separation of tantalum from columbium was reported by the Bureau of Mines: the HF-HC1-methyl isobutyl ketone system; data were presented for both laboratory and pilot-plant experimentation.3 Of twenty-eight organic solvents tested for their ability to extract tantalum from an HF-HC1 solution of columbium and tantalum, 3-pentanone (diethyl ke-tone), cyclohexanone, 2-pentanone, and 3-methyl-2-butanone were chosen for further study. Data on the HF-HC1-diethyl ketone system has been published4 and data describing the use of cy clohexanone, 2-pentanone, and 3-methyl-2-butanone as the organic phase are included in this report. RAW MATERIAL The source of tantalum and columbium oxides for this study was ('Geomines" tin slag from the Manono Smelter, Cie Geomines, Gelges, S.A., Congo. In order to extract the valuable Ta-Cb content, the slags were carbided, chlorinated, and the sublimate from chlo-rination was hydrolized and washed free of chloride with water. The washed material was air-dried and stored in a stoppered container. Throughout the paper, "feed material" refers to this mixture of hydrated oxides which was employed because of its high solubility in aqueous solutions. Typical analysis of the hydrated oxides is shown in Table I. I) HF-HC1-CYCLOHEXANONE SYSTEM Batch Separation. Effect of Acid Concentration. To determine the effect of varying the acid concentration upon the transfer of tantalum and columbium, a series of tests was made in which approximately 2.5 g of feed material was added to 25 ml solutions of 2, 4, 6, 8, and 10 N HF and 0 through 5 N HC1. Tantalum pentoxide concentration of the solu%ons was approximately 21 g per liter and columbium pentoxide was 14 g per liter. These starting solutions were shaken with equal volumes of cyclohexanone in 100 ml polyethylene bottles for 30 min. The phases were carefully separated in 125 ml glass separatory funnels. The time of contact of the solutions with the separatory funnels was kept at a minimum to reduce silica contamination. The measured phases were separated into 400 ml polyethylene beakers and the metal contents of each were precipitated by addition of an excess of ammonium hydroxide. Precipitate from each phase was filtered on ashless filter paper, ignited at 800" to 1000°C for 45 min, weighed, and analyzed by X-ray fluorescence.5 Data tabulated in Table I1 and illustrated in Fig. 1, show that maximum separation of tantalum from columbium for each HF concentration was obtained with no HCl present. The purest tantalum product was obtained with some HCl present. The highest separation factor was obtained at 2 N HF and

Jan 1, 1970

By M. M. Wong, D. E. Couch, G. M. Martinez

The Bureau of Mines electrowon hafnium metal with an average oxygen content of' 150 ppm at 700°C from an electrolyte containing 27 wt pct LiCl, 62 wt pct RbCl, and 11 wt pct HfC14. The average anode and cathode current efficiencies were 90 pct at anode and initial cathode current densities of 86 amp per sq ft. Haf-nium metal with an average oxygen content of 440 ppm was electrowon at 800oC from an electrolyte containing 90 wt pct KC1 and 10 wt pct HfCl4. The average anode and cathode current efficiencies were similar to those obtained in the LiCL-RbCl-HfCl, electrolyte. The chlorine gas given off at the graphite anode was vented through either a silica or a graphite tube to prevent cell corrosion. THE current method for the commercial production of high-purity hafnium is the thermal decomposition of Hfl4.1 The iodide method is not adaptable to continuous process techniques. Nettle, Hiegel, and Baker2 studied the electrorefining of hafnium from hafnium sponge containing 800 ppm oxygen. They failed to obtain hafnium with 600 ppm oxygen in their initial deposits and obtained AEC specification for oxygen only after 75 pct of the soluble hafnium had been removed from the electrolyte. Calculations using their data indicated this was approximately 4 lb of hafnium. The electrolyte was then used to produce approximately 3 lb of hafnium with a low oxygen content. However, no data are shown concerning the amount of anode material initially used or what percent of it was dissolved, therefore, results are not suitable for evaluation of a continuous operation. In general, it was not possible to consistently obtain low oxygen content metal with the electrolytes described by Nettle, Hiegel, and Baker. Wong, Hiegel, and Martinez3 investigated the electrorefining process for hafnium and showed that even by strict control of electrolyte composition only relatively low oxygen reduction could be obtained. The oxygen contained in the hafnium anode material tended to transfer to the cathode deposit and only a limited purification was possible. Both the "iodide" and the "electrorefining" processes depend upon hafnium sponge as a starting material. The sponge is normally produced by magnesium reduction of HfC14 ' and does not meet AEC specifications for hafnium metal. Since only 30 pct of the anode feed could be utilized3 in the electrorefining cells, the Bureau of Mines developed an electrowinning process. HfC14 was used as the feed material for the electro-winning process described in this report. Many of the electrolytes used in the electrorefining studies3 ap- peared to be suitable carrier-electrolytes for HfC14. However, in the initial studies on electrowinning, it was desirable to use electrolytes that had low solidus temperatures and could be operated over a wide temperature range to investigate parameters of the process. Therefore, electrolytes containing LiC1, NaC1, KC1, RbC1, CsC1, and HfC14, in various combinations were explored. EQUIPMENT Chlorinator. Hafnium carbide was chlorinated to produce HfC14 in the batch-type chlorination shown in Fig. 1. Chlorination temperatures were measured with a thermocouple placed in the center of the HfC charge. A flow meter was used to monitor the helium and chlorine. The exhaust side of the silica chlorina-tor tube was equipped with a flask for collecting organic material released during the initial heating of the HfC. The temperature of an internal heater, which extended from the HfC14 condensing flask to the hot end of the chlorinator, was adjusted to prevent the HfC14 from condensing before entering the collection flask. Helium and excess chlorine were exhausted through the lid of the collection flask to an aqueous NaOH solution. Sublimer. Initial studies were conducted using a sublimer, Fig. 2, made by placing a 13/8-in. OD nickel thimble 11 in. long, inside a 11/2-in. ID nickel bell 12 in. long, and locking it in place. This unit was loaded with HfC14 and partially immersed in the molten electrolyte for sublimation directly into the electrolyte. In another sublimer shown in Fig. 3, the HfC14 was contained in a "resin reaction flask". Quartz wool, previously heated to 600aC, secured between two nickel wire screens, was placed just above the HfC14 powder. The lid contained a vacuum outlet, a gage, an argon inlet, and an air-cooled pipe for condensing the HfC14. This sublimer was evacuated and heated. The sublimation temperature was not critical and the sublimer operated satisfactorily at all temperatures between 250" and 350°C. Electrolytic Cell. The electrolyte chamber, Fig. 4, was made of mild steel 8-in. schedule 20 pipe, 30 in. long. The exterior was metallized with a Ni-Cr alloy. The electrolytes were contained in a 16 gage nickel or iron liner with a nickel heat shield on top. The cell was heated by a resistance furnace. A 21/2-in. ID by 25 in. long air lock was connected to one port of a two-port cell cover assembly through a slide valve. The cover assembly of the air lock was electrically insulated from the cell and was equipped with a rubber sleeve that provided for the passage of the cathode lead. This allowed the cathode deposits to be removed and a new nickel cathode to be introduced without allowing air to enter the cell. A tube-rod assembly was bolted to the other port on the cell cover assembly and was sealed by a packing seal. The tube-rod assembly consists of a graphite

Jan 1, 1970

By J. M. Glass, C. D. Sholtess

We at Hughes Tool Co. are extremely proud of the quality of the hardware and techniques introduced through our efforts in tunnel-machine development and of the ready acceptance of them by manufacturers and contractors throughout the world. All of us who attended the conference and who are represented in the book, we believe, anticipate an increasing need for improved methods of underground excavation. The development of such methods will permit the further exploitation of our underground resources plus providing underground space for transportation and habitation at a rate compatible with demand. We hope that our comments, derived from the application of Hughes Tool Co.'s rotary boring experience, may provide a better understanding of the boring method as a possible avenue through which the desired improvement in underground excavation may be attained. We hope you may also become aware of some of the problems Hughes, as a pioneer developer of raise, shaft, and tunnel-boring equipment, must face prior to further major undertakings in this very interesting field. Our task is to acquaint you with the present status of the Hughes Tool Co. development program as it relates specifically to tunnel-boring machines. We feel this task can best be accomplished by a brief review of the general sequence of events leading to and constituting an equipment development program and more specifically the manner in which these events have occurred to either influence or further our boring-machine development program. The development of a machine to accomplish a certain function or series of functions is usually a very long and tedious process to which any number of individuals must contribute, perhaps over a period of centuries. Of particular irony is the fact that often those who contribute most generously do not necessarily reap the greatest immediate benefit. A marketable ma- chine will, however, result from a sequence of events such as these: 1) A problem or need is defined. 2) An idea is born to eliminate the problem.

Jan 1, 1970

By Frank A. Crossley

The work described has one serious shortcoming: grain size was not monitored during heat up to the temperature (1675°F) of isothermal anneal. The chief obstacles to p grain growth in titanium a-ß structures are the a grains. Heating a structure such as that shown in Fig. 1 to a temperature close to ß transus causes dissolution of a grains in order to adjust the a/ß ratio, compare Figs. 1 and 2. During this process there is a concomitant growth of the ß grains—otherwise ß grains would have to be nucleated in order to adjust the a/ß ratio. When the annealing temperature is reached the a/ß ratio stabilizes and this initial phase of ß grain growth ceases. After an incubation period growth resumes as described. Some 0 grain boundaries can be clearly seen in Fig. 1. (Skilled metallography commonly produces photomicrographs which show more structural detail than direct observation.) These boundaries suggest that the average ß grain size is approximately 0.002

Jan 1, 1970

Curious petroleum, mining and metallurgical engineers together with oceanographers, geologists, geophysicists and those in related disciplines, some 11,500 of them, poured into the Albert Thomas Convention Center in Houston, Tex. to see, and learn. This year's successful Conference, held April 22-24 featured a technical offering of 145 papers presented at 28 sessions. In addition, a total of 500 exhibits were provided by 260 companies.

Jan 1, 1970

By Eugene F. Broome, Hugh A. Walls

Self-diffusion coefficients were measured using the modified shear cell technique over a temperature range from 31" to 401°C. These data agree within experimental error with those of Petit and Nachtrieb, which were determined from 30°to 99°C. LIQUID self-diffusion behavior in gallium has been studied previously by Petit and Nachtrieb&apos; using their shear cell technique.&apos; Data were obtained in that investigation over temperatures ranging from 30.6 to 98.7"C, and these comprise the only published data for liquid gallium. In the work reported here, diffusion data were obtained from 30.8 to 400-9.Cusing a shear cell technique different from that of Petit and Nachtrieb. This modified shear cell method and the mathematical analysis for diffusion in a finite capillary geometry have been previously described.3 PROCEDURES AND ANALYSIS The segments of the modified shear cell were fabricated from types 304 and 316 stainless steel. No corrosion problems with gallium were encountered until attempts were made to obtain data at 500°C. The diffusion capillary diameter was 1.59 mm and capillary lengths for the four sets of cell segments used are given in Table I. Gallium with a purity of 99.999 wt pct was obtained from Var-Lac-Oid Chemical Co. The gallium-72 isotope used to measure the self-diffusion coefficients was produced from this inert gallium by irradiation in The University of Texas Reactor Facility, and was checked for radiochemical purity. Following a period of isothermal diffusion controlled to within 0.0l°C, the contents of each capillary section were emptied into a counting planchet, weighed to assure that gallium had not been lost during handling, and the radioactivity determined by windowless gas-flow proportional counting. Great care was taken to obtain a reproducible counting geometry for the samples. Diffusion coefficients were obtained from the solu-tion3 to the diffusion equation for the modified shear cell problem, viz.: in which B2= C2(t)Li + L2)- CSL1-COL2 P* (C0-Cs)(L1+ L2) and L, and L2 are the capillary segment lengths. The step-function initial concentration distribution was expressed as C(x,O) = C, for -L1 = x < 0 and C(x,O) = Co for 0 < x 5 L?. Also, C2(0 = ^- fcix,t)dx Computation of the diffusion coefficients was done numerically as described in our earlier paper.3 Because of the short half-life of gallium-72 (14.1 hr) and the relatively long counting times taken to give a precise value of the diffusion coefficient, it was necessary to account for the radioactive decay both during the actual count and between counts for different samples (before and after diffusion). Thus, the average count rate Cz(t) for a second sample counted from times t3 to t4 can be made consistent with the average rate Co determined for a first sample counted from times t1 to t2 by the relationship: N(e-»i-e-M») , , C2(f) " &-«,)(*-«.-«-»«) [21 in which N is the total number of counts taken and is the decay constant. RESULTS AND DISCUSSION The self-diffusion coefficients determined in this investigation are listed in Table 11. Values of the parameters used to calculate the coefficients are given in the table, as are values of the random counting error &apos; for the concentration ratio pz and the corresponding expected random variation AD in the diffusion coefficient. The quantity is calculable from: &apos;Cg — C2(t)} , „ ,2 .-„¦, M co-csl(ACs)J W in which AC is the error in the individual count rate. The mean values of the diffusion coefficients within each temperature group are also given in Table 11. Such temperature groupings of data are reasonable because of the small variations between the temperature~.~

Jan 1, 1970

By P. S. Kotval, O. H. Nestor

Stacking fault energy measurements of nickel-base alloys have been mainly confined to binary and ternary systems. In this paper, the stacking fault energy has been measured by the rolling texture method in a series of ten alloys which comprise successive additions of Cr, Mo, Fe, and C to pure nickel, eventually resulting in an alloy of the composition of Hastelloy alloy X. The alloys studied here are single-phase, solid solutions with the exception of two alloys in which some undissolved particles of "primary" carbide have been retained. It is found that successive additions of chromium, molybdenum, and iron all lower the stacking fault energy, with iron having only a minor effect. The stacking fault energy is found to increase when carbon is added in solid solution. The results from the rolling texture measurements are correlated with thin foil observations of dislocation substructures in these alloys. In a recent paper&apos; it was pointed out that the dislocation substructure of various superalloy matrices could be classified into three broad categories based on &apos;high&apos;, &apos;medium&apos;, and &apos;low&apos; stacking fault energy. It has also been demonstrated2 that the dislocation substructure in each of these categories has a well defined role in the nucleation of strengthening precipitates which is different from the role played by the dislocation substructure in other categories. Thus, it becomes desirable to understand the influence of various solute elements on the stacking fault energy and hence on the dislocation substructure of the matrix, before any further development of superalloys by mi-crostructural predesign can be undertaken. Recently, Beeston and France have studied the influence of increasing solute additions on the stacking fault energy of a series of binary nickel-base alloys relevant to the Nimonic series using the rolling texture method, and have then estimated the effect of a given alloy addition in five commercial Nimonic alloys. However, comparison with stacking fault energy data from other investigations&apos;&apos;5 suggests that the influence of a given solute element in a nickel-base binary system is not necessarily the same in a ternary or more complex superalloy system. Accordingly, the present work was undertaken to study the effect of successive addition of solute elements to pure nickel, the final composition being the nominal composition of Hastelloy X. The rolling texture method of stacking fault energy measurement was used since it can be used for the whole range of stacking fault energy values and does not have the disadvantage of, say, the Node method which is only applicable to low values of stacking fault energy. In addition, the rolling texture method provides a means of determining the stacking fault energy which is statistically more significant than that provided by other methods. EXPERIMENTAL TECHNIQUES Button heats of alloys of the compositions shown in Table I were prepared. Each button was remelted not less than four times. After a slight deformation (approximately 5 pct) all alloys were homogenized at 2200°F except alloys, H . I, and J. Alloys H and I were solution heat treated at 2150°F and alloy J at 2282OF. The buttons were cold worked by rolling, using "end-to-end" passes and intermediate anneals at the homogenization temperatures mentioned above. After each annealing treatment the samples were rapidly water quenched to avoid any precipitation. In alloys F and I, however, a few particles of "primary" carbides were retained even after the homogenization treatments at the temperatures mentioned above. Part of the solution heat treated material was cold worked to 0.04-in.-thick sheet and the penultimate reduction was -50 pct of deformation as recommended by Dillamore et al. All annealing was carried out in vacuo within sealed quartz capsules. Some of the material from each alloy was rolled down further to 0.004 in. strip for thin foil transmission electron microscopy specimens. Specimens of this strip were annealed at the homogenization temperature for 1 hr and then strained 7 pct by rolling at room temperature. Thin foils were prepared from the strip specimens by the &apos;window" technique using an Ethanol-Perchloric acid electrolyte at 32°F and a voltage of 22 v. Stainless steel cathodes were employed. All transmission electron microscopy was performed in a JEM-7 electron microscope using an accelerating voltage of 100 kv. Specimens from the 0.04 in. sheet which had been rolled -60 pct in the final pass were electropolished to remove the surface layers to a depth of approximately 0.002 in. Rolling texture pole figures for all the alloys were determined using a Schulz ring and nickel filtered CuKa radiation at 50 kv and 20 ma. The texture parameter Io/(lo + I,,) (where Io is the

Jan 1, 1970

By Freeman Bishop

The Nation&apos;s capital has one over- riding minerals-metals question: How long will prices and wages continue to move upward at the inexorable pace of the last year? Is the wage pattern to settle this year around 10%? The big industrial unions are coming into line, some of them later this year. Steel is standing in the wings for next year. Copper, steel, nickel and most other industrial metals have taken big hikes in their stride. Aluminum, which is making inroads into areas of industrial use formerly sacred to copper and some other metals, seems to be fighting to hold down its prices to attract customers from the higher-priced metals.

Jan 1, 1970

By Giovanni Barla

In engineering analyses, rock is frequently treated as a homogeneous, isotropic, and linearly elastic medium. However, rock material exhibits, in most cases, physical nonlinearity, time-dependency, and anisotropy. In some instances, anelasticity due to nonmechanical effects, such as density and temperature changes, must also be taken into account. The influence of these different anomalies in the physical behavior of a rock material is often neglected. An estimate of the discrepancies which are introduced when classical theory of elasticity is used is consequently required. The easiest theoretical path by which these discrepancies can be analyzed is to vary the constitutive equations of the rock material. The constitutive equations, representing various physical experiences with different types of media, define a class of ideal substances. Assuming different constitutive equations for the description of the rock physical behavior, one can establish a sensitivity analysis to determine how much information is contained in elastic theory for reliable design purposes. The method may be applied also using numerical techniques (e.g., finite element or finite differences) which, because of the nature of the formulation, require one to assume in each case the rock constitutive equations. Theoretical and experimental work is needed to be able to perform such analyses with confidence. Adequate and practical procedures for the determination of the material functions which enter the various constitutive equations must be established. This paper presents such an attempt with respect to the simulation of some behavior patterns of rock material. It is felt worthwhile to stimulate activity in this field and it is hoped to contribute to this objective.

Jan 1, 1970

By Sidney G. Parker, Jack E. Pinnell

Cubic ZnSe crystals have been grown with HCl, HBr, and I, as chemical transport reagents. The growth of large, well faceted crystals is in the order HCl < HBr < I, with some produced by I, transport being 1 cm on an edge. ZnSe:I crystals showed no strain in cross polarized light and were homogeneous as determined by electron microprobe. These ZnSe:I crystals had a mobility at 25°C of about 300 sq cm per v-sec and contained 103 dislocations per sq cm; further, these crystals were pure cubic and showed both photolumi-nescence and photo conductance. Doping was attempted during crystal growth. ZnS and Zn (S,Se) crystals were also grown by I, transport. ZnSe crystals have been grown by most of the common techniques.1"6 Of these, chemical vapor transport as developed by Kaldis2 showed promise of producing high perfection, cubic ZnSe. With chemical transport, crystals can be grown far below the transition temperature for conversion of cubic to hexagonal structure. A small temperature gradient between source and growth region tends to allow near equilibrium growth, which should lead to good crystalline perfection. Because it is necessary to form volatile compounds, halogens or their compounds are extensively used as transport agents. HCl, HBr, and I, were compared as transport agents for growth of ZnSe crystals. Doping during crystal growth was attempted. ZnS and ZnSxSe1-x crystals were also grown by chemical transport. Electrical and optical properties of the crystals were examined. EXPERIMENTAL Fig. 1 is a sketch of the growth apparatus, which was a modification of Kaldis&apos;. A quartz rod which butted against the seed crystal served as a heat sink to promote growth only on the seed. Initially, crystal growth was attempted with ZnSe prepared by wet chemical methods. Because such a material contained oxides and selenates, HC1 reacted with these oxygen compounds to form H2O and no ZnSe transport occurred. Further, this ZnSe contained about 10 ppm Cu plus other impurities. High purity ZnSe, containing small amounts or no oxygen compounds, was prepared by reaction of H2Se with zinc vapors at 600° to 800°C in an atmosphere of palladium diffused H2.7 The ZnSe thus produced was usually copper free and contained only 1 ppm by weight of magnesium as determined by emission spectrographic analysis. G.E.&apos;s electronic grade ZnS, which had about 10 ppm metallic impurities, was used without further treatment. ZnS-ZnSe charges were prepared by mixing the two compounds together, then sintering 24 hr at 1100°C in an H2S ambient. The gases, HBr and HC1, were commercial grade, of 99.9 pct purity, while the I2 was 99.999 pct pure. The single crystal seeds with (111) faces for growth were mechanically held in place by sealing the outer quartz tube to the growth tube. After the charge was introduced, the system was outgassed at l0-5 torr while heated to 500°C. The desired quantities of transport reagents were then sealed off in the system, which had been cooled to 25°C. The growth tubes were inserted into a furnace with the desired temperature and thermal gradient. The system was allowed to cool for 1 to 12 hr from growth temperatures of about 850" to 25°C. The tubes remained stationary during crystal growth. HCl CHEMICAL TRANSPORT GROWTH Preliminary runs with HC1 had a temperature gradient of 100° to 200°C and no seed or heat sink. Only small crystals were obtained in these experiments as

Jan 1, 1970

By R. S. Hendricks

During the past two years, Hecla Mining Co. has used shotcrete for ground support purposes with substantial cost savings. Hecla&apos;s experience indicates that shotcrete is an extremely effective ground support material which can be successfully employed in problem areas where various conventional systems would fail. To date, Hecla&apos;s use of shotcrete has been limited to openings where long-term high-quality support was required, to specific areas where exceptionally difficult ground conditions were encountered [ ] and, occasionally, to stope support. It is anticipated that when the necessary engineering and procurement of the components of an efficient materials handling system have been accomplished, shotcrete will be used on a routine basis in stoping and development situations. The real future of shotcrete is based, however, on the fact that this form of support offers greater opportunity for mechanization than rock bolt, timber or steel. Unlike all other support operations, which are cyclic, shotcrete is a continuous process, and therein lies its potential for incorporation into an efficient rapid excavation system. Hecla has obtained and operated both wet- and dry-process shotcrete application equipment. The operational characteristics of the two processes are presented and discussed below.

Jan 1, 1970

By T. S. Lundy, R. E. Pawel

Effects of short-circuiting paths on observed diffusion behavior in real crystalline systems are considered. It is concluded that experimentally measured diffusion coefficients may vary widely from values associated with true lattice diffusion to those appropriate for short-circuiting paths alone depending upon the conditions of the experiment and the sensitivity and completeness of the data. IN 1961 Harrison1 classified experimentally observed diffusion behavior in systems involving both lattice and dislocation-enhanced atomic transport into three categories. His types A and C were limiting cases where the measured diffusion coefficients using the usual serial sectioning techniques were representative of lattice diffusion (with only slight effects of dislocations) and dislocation-network diffusion (with only slight effects from lattice diffusion), respectively. In both cases A and C for a sectioning experiment in which the diffusion couple has the semi-infinite geometry and unidirectional diffusion takes place from a plane source located at the origin of the coordinate system, he pointed out that Gaussian penetration behavior should be observed. For case A the lattice coefficient should be slightly enhanced by dislocation contributions, and the model of Hart2 can be used to predict the amount of this enhancement. For case C the observed coefficient should be that for the dislocation network and have little or nothing to do with lattice diffusion. Hart&apos;s model in its general form is still applicable, but the apparent diffusion coefficient for this case is virtually independent of dislocation density. Harrison&apos;s case B is that for which non-negligible contributions exist for transport by both mechanisms and "the concentration distribution does not approximate to any simple form." With the possible exception of very low temperatures where diffusion distances approach atomic dimensions, any real diffusion specimen will in general show non-Gaussian behavior provided the data are complete enough to uncover it. Thus, a "complete" set of data for any real system cannot be expected to be totally amenable to treatment by a simple mathematical model which does not take into account the relative distribution of transport by the various modes or submodes. For a particular sensitivity of specimen analysis and at a given temperature of annealing if the concentration profile could be observed as a function of time, "the behavior will initially be type C, and will develop into type B. and ultimately type A".1 Thus, it appears that the type of diffusion behavior which is experimentally observed in a given system depends very much upon the particular experimental conditions and the particular tools available to obtain the data. Interpreted in this fashion, Harrison&apos;s classification refers not to the diffusion behavior as a whole but to the relative sensitivity of the method of examination and completeness of the data. In the present paper we will: 1) assume that the measured penetration behavior follows the same general classification as that proposed by Harrison; 2) give, for diffusion in tungsten, examples of the range of diffusion coefficients that it is possible to obtain in a single experiment; and 3) speculate that some arguments which consider the possibility of short-circuiting effects being responsible for "anomalous" diffusion in several bcc metals may have been discarded prematurely.3 Both monocrystalline and polycrystalline tungsten were used in this work. The former was obtained as triple-pass, electron-beam zone-refined 6.0-mm-diam rod from Materials Research Corp.. Orange-burg. N.Y. The latter was swaged, high-purity bar stock which was annealed for 2 hr at 2200°C and 10-9 torr to yield an average grain diameter of about 25 /i. The tracer Nb95 was received in oxalate solution from the Isotopes Division, Oak Ridge National Laboratory. Experimental procedures of specimen preparation, isothermal diffusion annealing, serial sectioning by anodizing-and-stripping, and y counting of the individual sections to determine the tracer penetration profile after a given diffusion anneal have been described elsewhere.4 However, it should be noted that the anodizing-stripping sectioning technique has the particular advantage for this research of permitting sectioning to be accomplished over a wide range of

Jan 1, 1970

By J. W. Edington

THE Hopkinson bar has become a popular technique for the measurement of the mechanical properties of materials deformed at high strain rate. Maximum use of the equipment is made in the arrangement first used by Kolskyl in which a short compression specimen is sandwiched between two pressure bars and is loaded by a single pulse travelling through the system. The pressure bars are used both to apply the load to the specimen and as transducers to obtain continuous strain-time histories of three pulses, incident on, reflected, and transmitted by the specimen. The data measured by the pressure bars can be analyzed in terms of the stress/strain behavior of the specimen.2&apos;3 However, one of the assumptions of the analysis of the observed pulses is that the total stress and total strain do not vary significantly from point to point within the specimen at any given instant during the deformation process. Although this assumption is generally justified for very short disc-like specimens2 the situation is uncertain for larger specimens. For example, at small plastic strains (-0.01) Hauser et al.2 have some evidence of small flucations in the total stress within the crystal during deformation, even in relatively short aluminum specimens. In addition, Karnes4 has shown that the plastic strain, and by inference strain rate, is different at each end of a compression specimen tested in a Hopkinson bar, although the length of the specimen was not specified. Recently, the mechanical properties and the dislocation substructure have been investigated in single crystals of columbium5 (length 0.25 in., diam 0.19 in.), and copper6 (length 0.5 in., diam 0.5 in.) deformed at high strain rates. As part of this research program the assumption that the plastic strain is constant throughout the specimen has been checked by measuring the total dislocation density as a function of position in the specimen. Compression specimens of the same orientations and dimensions were tested as described previously5,6 sing a split Hopkinson bar. Since any discontinuity in strain distribution is most likely to arise during the initial stages of deformation the investigation was performed on specimens deformed to plastic shear strains of 0.054 (copper) at a strain rate 1.2 x l03 sec-1, and 0.06 (columbium) at a strain rate 1.5 X l03 sec-1. The orientation of the single crystals is shown in Figs. 1 and 3. Thin foils were taken parallel to the most highly stressed slip plane, i.e., (111) in copper and (011) in columbium, using conventional disc techniques. The dislocation densities were measured using first order reflections with compensation for invisible dislocations.5&apos;6 In the copper single crystals the discs were randomly distributed throughout the cross section of the specimen. However, the dislocation density obtained from each disc was plotted vs the disc positions relative to the ends of the specimen. The results for the copper specimens are shown in Fig. 1. Clearly the dislocation density is constant throughout the main portion of specimen within the experimental error. The error bars on the dislocation densities correspond to a shear strain variation of 0.015 on the basis of previous measurements% ± of the rate of increase of dislocation density with strain in copper single crystals of the same geometry. Thus within this experimental error the plastic strain can be concluded to be constant within the specimen and the assumptions used in the analysis of the stress/time curves are therefore reasonably valid. The higher measured dslocation density near the impact end and the lower dislocation density at the bar end of the copper specimen is in agreement with the results of Karnes4 who showed that this strain/time curve rose to a maximum more rapidly at the impact end compared with the bar end. Hauser et al.2 have also pointed out that at small plastic strains (-0.01) the strain at the impact end of the specimen may be greater than that at the bar end. Thin foils taken from different points within the columbium single crystals demonstrated that the dislocation density could vary significantly within the specimen, see Fig. 2. Large areas of some thin foils up to 30 µ sq contained very few dislocations, see Fig. 2(a). However, in other parts of the compression specimen dislocation configurations like those shown in Fig. 2(b) existed over large areas (-30 µ sq). As a result, when the average dislocation density in a thin foil is plotted as a function of the position of the thin foil relative to the ends of the specimen, considerable scatter is observed, see Fig. 3. In this material then, the local dislocation density, and consequently the

Jan 1, 1970

By John W. Wagner, Robert K. Willardson

Single crystals of Pbl-xSnxTe have been grown from nonstoichiometric, cation-rich melts with the objective of producing as-grown, bulk material containing carrier concentrations ranging from 1016 per cu cm to 1018 per cu cm. Three specific crystal composi-tions were investigated in detail; x = 0.00, 0.10, and 0.17. Pull rates of from I to 3 mm per hr were used. Single crystals were successfully pulled from melts containing as little as 30 at. pct Te. Hall coefficients, resistivities, and carrier mobilities of these materials were determined. The relationship between the composition of the melt and the carrier concentration in the as-grown crystal has been studied for the three crystal compositions of interest. BULK single crystals of Pb,-,Sn,Te have previously been grown from stoichiometric melts.1,2 Such crystals are p-type and have relatively high carrier concentrations ranging from -9 x 1018 per cu cm for PbTe to -8 x l020 per cu cm for SnTe at 77°K. These carrier concentrations result from deviations from stoichiometry (lead vacancies) in the as-grown crystals. Since lower carrier concentrations are desirable for electrooptic device applications, these crystals are usually subjected to long-term, isothermal anneals.&apos; This paper reports on the growth of Pbl-xSnxTe single crystals from nonstoichiometric melts with the primary objective of producing as-grown material containing relatively low (1016 to 1018 per cu cm) carrier concentrations and also reports on the general characteristics of these crystals. The phase relationships in the Pbl-xSnxTe systems are such that materials solidifying from nonstoichiometric, cation-rich melts will have smaller deviations from stoichiometry than materials grown from stoichiometric melts. Fig. 1 is the T-x phase diagram for PbTe in the vicinity of the stoichiometric composition.3 A crystal grown from a stoichiometric melt will solidify at a melting point maximum at which the solid will contain -0.5002 atom fraction of tellurium. PbTe single crystals grown in our laboratories from stoichiometric melts have carrier concentrations of 9 x 10" per cu cm, indicating that the excess tellurium in the crystals is as expected from this phase diagram. However, growth of PbTe from a lead-rich melt will result in material having a more nearly stoichiometric composition. Although the addition of Sn to the melt shifts the solidus curve further toward the tellurium-rich side,4 the general discussion given for PbTe applies to the Pbl-xSnxTe systems as well. EXPERIMENTAL Single crystals of Pb1-xSnxTe have been grown in our laboratories from nonstoichiometric, cation-rich melts using the Czochralski technique and boric oxide liquid encapsulation. The details of the growth apparatus and growth technique have been reported in a previous paper on growth of these alloys from stoichiometric melts.2 In the present study, three specific crystal compositions were investigated in detail; x = 0.00, 0.10, and 0.17. Growth of alloy crystals from nonstoichiometric melts requires considerable care, and good quality single crystals were obtained in this study only by optimizing the mechanical and thermal stability of the growth system and by using pull rates of from 1 to 3 mm per hr (Pbl-xSnxTe crystals are easily pulled from stoichiometric melts at rates of from 5 to 10 mm per hr). The liquid encapsulation technique was found to yield near-ideal conditions, since the B2O3 layer increased the thermal stability at the growth interface, permitted easy attainment of near-ideal thermal gradients, and dampened vibrations at the melt surface. The hygroscopic character of the B2O3 was a slight problem and vacuum heat treating was necessary to completely remove the water from the boric oxide. During initial growth, the seed diameter was reduced and a narrow neck (1 to 2 mm diam) of several millimeters length was grown. The latter steps were found to be necessary for the growth of single crystals, i.e., if either of these two requirements were

Jan 1, 1970

By F. J. Smit

The Tintic Division of Kennecott Copper Corporation operates the Burgin Mine in the East Tintic District near Eureka, Utah, approximately 75 miles south of Salt Lake City. The Division was organized in 1963 to develop concealed replacement orebodies in the area located by Bear Creek Mining Company, Kennecott&apos;s domestic exploration subsidiary. Several accounts of the discovery of these orebodies and the general geology of the area have been given (1, 2). The ore is very complex and of varying degrees of oxidation and alteration. The principal metal values are lead, zinc, silver, and cadmium. The Burgin No. 2 shaft was sunk in 1963 to provide access for mining the orebodies. Much of the ore mined after this shaft was sunk was high grade and shipped directly to smelters, but concurrently it was recognized that concentration would be required to upgrade the lower grade ore in the deposits. Accordingly, in January 1967, authorization was received to construct a 500 tpd concentrator. From the beginning it was expected that the characteristics of the mill feed would vary widely and unpredictably as mining progressed. This paper summarizes the characteristics of the ore, the test work upon which the plant design was based, the engineering and construction, and the present operation of the facility eighteen months after startup. ORE CHARACTERISTICS The Burgin ore is a complex assemblage of minerals from replacement of limestone. The replacement and alteration occurred in at

Jan 1, 1970