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By S. Dixon

The Peñasquito deposit in Mexico is one of the largest grassroots precious metal discoveries developed in recent years. Goldcorp Inc. acquired the gold-silver-lead-zinc property with the acquisition of Glamis in 2006. The Peñasquito property is located in the northeast corner of Zacatecas state, Mexico. The elevation of the property averages 1,900 m (6,200 ft) above sea level. It is situated in a wide east-west valley. Minera Peñasquito, S.A. de C.V., wholly owned subsidiary of Goldcorp Inc. has operated twin SABC-flotation lines since June 2010. Each line is capable of processing 50,000 tonne per day using a SABC grinding circuit followed by flotation to produce lead and zinc concentrates. The addition of a Polysius Polycom HPGR 2417 with dual 3,000 kW drives to reduce the crushed pebbles produced by the SAG mills was planned to increase production to 65,000 tonne per day per line. The production has been limited to 100,000 tpd due to operating issues in the grinding circuit. The operating issue has been the amount of pebbles generated by the SAG mills is significantly less than originally planned. This paper describes the planned installation of a crushed augmented feed to maximize utilization of the HPGR capacity. This coupled with full utilization of the SAG and Ball Mill available energies there is the potential to achieve a 50% total capacity increase from the grinding circuit.

Jan 1, 2012

By Dinabandhu Ghosh

The stability diagram of the ternary system Ti-C-O is drawn in the present work for four temperatures, 1473 K, 1573 K, 1673 K and 1873 K, showing the stability fields of TiC, Ti, TiO, Ti2O3 and TiO2. An analysis of the system reveals that it is not possible to produce Ti in a carbon-saturated system; instead, the stable product is TiC. However, Ti can be an equilibrium product in a carbon-unsaturated system at prohibitively low values of pCO and pCO2; for example, at 1873 K, these limiting values are: pCO = 10-6.76 atm and pCO2 = 10-13.36 atm, which, when exceeded, produce TiC and TiO respectively. The magnesiothermic reduction is examined through the stability diagram (log pO2 vs. 1/T) of the system Ti-Mg-O, drawn for a number of pMg values, presenting the pairwise stability fields such as Mg(l) + Ti(s), Mg(v) + Ti(s), MgO(s) + Ti(s), etc. The diagram shows that magnesium reduction of TiO to Ti is possible only at T < 1641 K with pMg = 1 atm and only at T < 1569 K with pMg = 0.5 atm. An experimental scheme for Mg vapor reduction of TiO2 is proposed in which an argon stream is passed over a Mg(liquid) evaporator at temperature T1 (say, 1234 K); the resulting Ar-Mg mixture at 1 atm is subsequently flown over a bed of TiO2 at temperature T2. A thermodynamic analysis establishes that complete reduction to Ti is possible and that the best results are obtained when T1 = T2.

Jan 1, 2006

By J. C. Branner

(New York Meeting, February, 1012.) DURING the past two or three years I have seen statements regarding the first discovery of fullers&apos; earth in this country that seem to require correction or modification. One appears again in the Mineral Resources of the United States for 1909, part II., p. 738, where it is said that ? fullers&apos; earth was first discovered in the United States at Quincy, Fla., in 1893." In this connection the following bit of unpublished history concerning fullers&apos; earth in the State of Arkansas is worthy of attention. In April, 1891, when I was State Geologist of Arkansas, a Swede named John Olsen, formerly a copper-miner in Michigan, but then living south of Little Rock, brought me a sample of what he thought was fullers&apos; earth. He reported that it was from near the town of Alexander, the locality being 1 S. 13 W., section 8, SW. quarter of the SE. quarter. My chemist, Dr. R. N. Brackett, analyzed the material, and reported: Per Cent. Silica, . . SiO2., . 63.19 Alumina, . AL2O3, . 18.76 Iron, . Fe203 7.05 Lime, . Ca0 0.78 Magnesia, . . MgO 1.68 Potash, . K2O 0.21 Soda, . . Na2O 1.50 Loss on ignition, . 7.57 100.74 Water at 110° to 115° C. = 7.84. Another sample from about the same locality, also brought by John Olsen, on analysis by Dr. Brackett gave

Jul 1, 1912

By Umetsu Y, Su Qing-quan

The electrical conductivity of electrolytes is very important in metal electrowinning and electrorefining processes as it has a very large influence on the energy consumption. There are few measurements of the electrical conductivity of zinc electrowinning electrolytes available. The electrical conductivity of zinc electrowinning electrolytes were measured over the following range of concentrations, 100 to 200 g/L H2SO4, 0 to 80 g/L Zn and 0 to 20 g/L Mg, at temperatures from 25 to 60 C. The specific conductivity of zinc electrowinning electrolytes may be expressed as a function of concentrations of sulphuric acid, zinc and magnesium and temperature by the following equation: k(S/cm) = 0.004 + 0.00115 t + 0.00282 [H2SO4] t - 0.00114 [M] t + 0.3442 [H2SO4] - 0.0451 [1-12SO4]2 - 0.1058 [M]ò[H2SO4] - 0.0224 (M] + 0.0286 [M] 2 where t is temperature (C), [H2S04](moI/L) is sulphuric acid concentration and [M](mol/L) is a sum of zinc and magnesium concentrations, ([Zn]+[Mg]). The specific conductivity of zinc electrolytes is a function of the concentrations of the hydrogen ion and free water molecule, which are calculated by taking into account the second ionisation constant of sulphuric acid and the stability- constants of zinc-sulphate and magnesium-sulphate complexes at 25 C.

Jan 1, 1993

By J. C. Branner

DuriKg the past two or three years I have seen statements regarding the first discovery of fullers&apos; earth in this country that seem to require correction or modification. One appears again in the Mineral Resources of the United States for 1909, part II., p. 738, where it is said that &apos;fullers&apos; earth was firat discovered in the United States at Quincy, Fla., in 1893." In this connection the following bit of unpublished history concerning fullers&apos; earth in the State of Arkansas is worthy of attention. In April, 1891, when I was State Geologist of Arkansas, a Swede named John Olsen, formerly a copper-miner in Michigan, but then living south of Little Rock, brought me a sample of what he thought mas fullers&apos; earth. He reported that it was from near the town of Alexander, the locality being 1 S. 13 W., section 8, SW. quarter of the SE. quarter. My chemist, Dr. R. N. Brackett, analyzed the material, and reported: Per Cent. Silica,.....SiO2 .... 63.19 Alumina,.... A12O3 .... 18.76 Iron,.....Fe2O3 .... 7.05 Lime,.,...CaO .... 0.78 Magnesia,.... MgO .... 1.68 Potash,.... K2O ,... 0.21 Soda,.....Na2O .... 1.50 Loss on ignition,........7.57 100.74 Water at 110° to 115° C. = 7.84. Another sample from about the same locality, also brought by John Olsen, on analysis by Dr. Brackett gave :

Jan 1, 1913

By Wm. Bleloch

SYNOPSIS The feasibility of rejection of nitrogen from stack gas by recycling of top-gas with cold-blast 95 to 98 per cent oxygen is assessed. It is concluded that nitrogen could be removed to a level below 3 per cent of the stack-gas using recycled cleaned cold top-gas and 95 to 98 per cent oxygen produced by a single-column air separation plant with liquid-air reflux instead of liquid nigroten reflux. The partial pressure of carbon monoxide in the stack-gas could be doubled without top pressure and the deposition of carbon and reduction of iron oxides by the stack gas could be markedly increased. It is concluded that rate of stock descent could be doubled compared with present-day practice and an iron make of 4400 kg/m 2 of hearth area/hour may be attainable with a working volume of 129 m2/m hearth diameter by this method. Oxygen consumption is estimated at possibly 280 m2 (0.4 t)/t iron made and carbon consumption at possibly 280 kg/t iron. Power consumption for oxygen including auxiliaries is estimated at 304 kWH/t 100 per cent oxygen as 95 to 98 per cent gas. The air separation cycle due to Schuftan is discussed in some detail, and the components of a 1 600 metric ton/day separation plant are briefly discussed. The paper is a theoretical essay.

By Bratt G. C

When zinc is electrowon from concentrates or residues of high manganese content, the zinc sulphate electrolyte may contain high levels of manganese. The manganese level is usually controlled by removal of manganese at the anode as the dioxide during the zinc electrodeposition stage.By increasing the temperature of an electrolyte containing 8 gll manganese from 35&#176;C to 90&#176;C, decreasing the acidity from 100 gll to 30 gll, reducing the current density from 432.A.m-2 t6 286 A.m-2 and using new lead anodes containing over 0&#183;75 per cent silver, the current efficiency for manganese dioxide removal was .increased from less than 1 per cent to over 20 per cent.At the same time the current efficiency for zinc deposition was maintained above 90 per cent by additions of high molecular weight cationic polyacrylamides to the cell. By utilization of these conditions in a small number of cells, the manganese level in the circuit could be economically concrolled to any desired level.With these conditions the formation of the dioxide changes from an indirect route, via Mn04-, to a direct oxidation of Mn2+ to Mn02.The incorporation of impurities such as Pb, K, NH4, Cu, Ni, or Fe makes the deposit unsatisfactory for direct use in Zn - Mn02 dry cells. The deposit structure changes from cryptomelane at low temperature to y3-Mn02 at high temperature.

Jan 1, 1975

By Barry Welch, Kai Grjotheim

"The capital intensive nature of the present process for aluminium smelting, the low productivity per unit reactor, and pressure to reduce the electdcal energy demand have motivated the search for an alternative process. Optional routes include carbo-thennal reduction of alumina (using an alloy phase immediate), chlorination followed by electrolysis of aluminium chloride, and electrolytic decomposition of alumina using inert electrodes. Materials performance and reactor design constraints have been a major limitation for alternatives. However parallel research has led to considerable advances in the process efficiencies and scale of both the Bayer process and the Hall-Heroult cells.Future development of alternatives is dependent on breakthroughs in Materials Science such as the development of stable conducting anodes and stable wettable cathode materials. However if successful it is probable that some spin-offs will be retrofitted to existing cells. Therefore the basic Hall-Heroult technology will continue as the dominant process for at least the next half century.IntroductionThe present Bayer/Hall-Heroult process for production of aluminium is just over 100 years old. Frequently, during the century of its technological development questions have been raised as to whether it is the ideal route for metal production. The questioning was particularly intense in the period between the mid-1950s and the mid-1060s when the limitations of the Hall-Heroult electrolytic stage were evident. Then the cells were only 30 to 35 % efficient with respect to the utilisation of the hydro-electricity (requiring 16 18 kWh/kg, thus making it strongly dependent on cheap electrical energy), while the size of cells being installed were 80 to 130 kA. Hence the process was capital intensive, the productivity per reactor volume low, and the process was labour intensive (requiring about 15 man hours per tonne of aluminium). In addition, the use of pitches for Soderberg anodes and the fluoride electrolyte were creating problems through emission. At that time material science had paid little attention to the manufacture and use of carbonaceous materials or the development of alternatives that were corrosion resistant to cryolite or liquid aluminium. Hence cell lives were short - typically 600 to 1000 days - thus adding to operating costs. Thus the grass over the fence, as foreseen in 1960s vintage crystal balls appeared much greener for alternative processes."

Jan 1, 1988

By M. D. Banus, M. B. Bever, A. K. Jena

Measurements of the heat effects on addition of various solid phases of InSb to a solution calorimeter have confirmed the existence of a new high-pressure plzase InSbQII), which was jormed at 37 kbar and 623°K. The heats of formation of InSbQ), InSbfl), and InSb(111) have been measured and the heats of transjormation of InSb(1) to InSb(1l) and InSb(Iii) and of &apos;InSb(ll) to InSb(iiI) at atmospheric pressure and 78°K have been determined independently. Other thermodynamic properties of these phases, the phase boundary InSb(i)/InSb(ll), and the relative positions of the triple points are calculated. The pressure-Dolume term is shown to be more critical than the internal energy difference in making the phase InSbQ) less stable than InSb(II) at high pressures. Observations on partially transformed specimens suggest that the transformation of&apos; metastable InSb (III) takes place in the two steps InSb(lll) — InSbQI) and InSbQI) — InSb(l). The compound InSb can exist as several phases depending on the temperature and pressure. The phase InSb(I), which is a semiconductor having the zinc-blende structure, transforms into the phase InSb(11) at a pressure of about 23 kbar and room temperature.&apos;-3 The phase InSb(II) has the 6 tin structure,4-6 is metallic, and becomes a superconductor below 2.lok7 An orthorhombic phase of InSb has been reported to occur above a pressure of approximately 30 kbar at room temperature.&apos; This phase is also a superconductor below about 3.4k9 The phase InSb(I1) is retained at atmospheric pressure if it is quenched to 78°K under pressure,6,7 but tends to transform at higher temperatures and does so explosively above approximately 210ok,7 A value of the heat of transformation of InSb(1) to InSb(II) at 210°K and atmospheric pressure has been reported.&apos; A concurrent publication&apos;0 reports evidence for a new solid phase InSb(III), which was formed by annealing InSb above 58 1°K at 37 kbar, above 560°K at 52 kbar, and above 473°K at 65 kbar. Specimens of InSb(II1) had a superconducting transition temperature of 4.lo ± 0.1oK. X-ray diffraction patterns also indicated the presence of a new phase. In this communication, thermodynamic properties of the phases InSb(I), InSb(II), and InSb(II1) are reported and discussed. The phase boundary InSb(I)/InSb(II) and the relative positions of the triple points are calculated. 1) EXPERIMENTAL 1.1) Preparation of Specimens. Specimens of the high-pressure phases of InSb were prepared in a belt-type high-pressure chamber by a tichnique described in detail elsewhere.&apos;&apos; Charges of 0.6 to 0.8 g of Cominco semiconductor grade InSb were used. The temperatures were measured with a chromel-alumel thermocouple, the hot junction of which was within 0.5 mm of the specimen. The specimens were heated under a pressure of 37 kbar for 1 hr at 723°K (which is well above the liquidus) and quenched to room temperature. One set was annealed at 523°K and another set at 623°K for 20 to 24 hr. These annealing temperatures were chosen because 523°K was below and 623°K was above the temperature of the transformation of InSb(II) to InSb(1Tl) at 37 kbar.10 The specimens were then cooled to 125°K by passing liquid nitrogen through the jacket of the belt die before the

Jan 1, 1968

By B. A. Kulp, R. R. Reeber

Lattice parameters for the wurtzite form of&apos; CdS mere measured by powder X-ray diffraction techniques over the temperature range 26° to 1000 K&apos;. A negative thermal -expansion coefficient was determined parallel and perpendicular to the c axis below 13.5°K. Near 808°K, the curves of the a, parameter, the differential thermal expansion, and the resistivitv vs absolute temperature exhibited a reversible discontinuity. j\. direct determination of the lattice parameters of CdS over an extended temperature range is of particular use in interpretation of the effects of high pressure on the band structure of semiconductor materials. For example, Langer&apos; has studied the effects of pressure on the absorption edge of CdS from 77°K to room temperature. In order to properly interpret his results it is necessary to know the change in lattice constants as a function of temperature for the wurtzite, zincblende, and rocksalt2 modifications of this material. Thermal-expansion measurements for CdS, wurtzite phase, have previously been determined over a limited temperature range.s74&apos;5 The zincblende phases, of other compounds with similar bonding, have been investigated by Novikova et a! .&apos;&apos; Previous thermal-expansion measurements for metals with hexagonal symmetry have for the most part used bulk sample techniques for low-temperature result. Most X-ray measurements at low temperature have been on materials with cubic symmetry.&apos;jQ The use of X-ray powder-diffraction techniques offers a convenient and accurate method of measurement over a wide temperature range. The X-ray method, besides measuring unit cell size, also provides information about change in structure type and the nature of some of the stacking defects present. I) EXPERIMENTAL PROCEDURE The thermal expansion was measured by X-ray powder-diffraction techniques. Lattice parameters for the range 300" to 1000°K were determined in a 19-cm Unicam camera. From 26" to 300°K a 34-cm back-reflection camera with the Van Arkel" film arrangement was used. The sample was bonded to a copper cold finger below a helium, hydrogen, nitrogen, oxygen, dry ice, or ice brine bath in a metal cryistat. che inner bath was surrounded by another one for radiation shielding. Nitrogen was used in this outer bath for the very low-temperature runs while dry ice and acetone was used for the 195°K point. CuKa X-rays, filtered with a nickel foil, entered the cryostat through a 0.001-in.-thick aluminum window. During a typical 9-hr exposure the sample was continuously rotated through a 30-deg arc. Heat losses from the sample holder limited the minimum temperature attained to 26°K with helium in the inner dewar. Temperature calibration between 25" and 125°K was determined with an Allen Bradley 300-ohm carbon resistor. The resistor was attached to the back of the copper finger with a thin layer of silicone grease. Liquid He, HZ, N2, and O2 baths were used to calibrate the resistor. Temperature variation is estimated to be *2"K with a slightly higher uncertainty at room temperature and 264°K. The thermal expansion of high-purity germanium from 373" to 900°K was measured as a calibration for the Unicam camera. This value was compared with the known value1&apos; for the thermal expansion. A discontinuity in the thermal-expansion coefficient of the CdS a, parameter was observed. Similar discontinuities in the resistivity and differential thermal analysis at the same absolute temperature indicate that the temperature error is *5°K in the temperature range above 500°K. Temperature errors in the Unicam camera below 500°K are appreciably higher due to large gradients in the resistance furnace. II) SAMPLE PREPARATION The CdS was Eagle Picher ultrapure grade. The crystallite size range of the as-received powder

Jan 1, 1965

By Alden B. Carpenter

Brines (a solution containing more than 100,000 mg/l dissolved solids) in sedimentary basins are genetically related to evaporites in one of three ways: (1) dissolution of evaporite minerals, generally halite; (2) interstitial fluids in evaporites which are expelled from their source rock during compaction; (3) incongruent alteration of hydrous evaporite minerals (carnallite). Most brines of commercial interest originate as interstitial fluids in halite-rich rocks. These fluids are initially rich in magnesium, sulfate, and potassium but evolve in composition to an Na-Ca-Cl brine during their migration to their present host rock. A particularly important reaction in this process is: 2CaCO3 + Mgt+ = CaMg(CO3)2 + Cat (1a) This reaction reduces the magnesium content of the brine and ultimately converts it to a calcium-rich brine. The loss of sulfate occurs by two mechanisms. One of these is by direct precipitation as calcium sulfate (either as gypsum or anhydrite): Ca 2+ + SO4 2- = CaSO4 (2a) Additional sulfate may be lost from the brine by sulfate reduction: S042 + 1.33CH2 + 214- = H2S + 1.33CO2 + 1.33H20 (3a) The loss of carbonate species occurs by precipitation as carbonates: H2O + C02 + Cat = CaCO3 + 2H+ (4a) The net result of reactions 3 and 4 is: Ca 2+ + SO4 2- + 1.33CH2 = H2S + CaCO3 + 0.33H2O + 0.33CO2 (5a) The loss of potassium from brines is caused by the formation of potassium aluminosilicates: 2K+ + Al2Si2O5 (OH)4 + 4SiO2 = 2H + 2KAlSi2O8 + H2O (6a) The hydrogen ions produced by this reaction dissolve carbonates: 2H+ + CaCO3 = Ca 2+ + CO2 + H2O (7a) The net result of reactions 6 and 7 is: 2K+ + CaCO3 + Al2Si205 OH)4 + 4SiO2 = Ca2 + 2KA1Si3O8 + H2O + C02 (8a) A study of brines from central Mississippi has shown that the composition of these brines is related to the interaction of pore fluids from evaporites with minerals in the enclosing sedimentary rocks.

Jan 1, 1978

By B. B. Rath, Hsun Hu, R. S. Cline

FOLLOWING the techniques used by Beck and co-worker~ "~ the characteristics of selective growth in the recrystallization of a zone-refined aluminum single crystal, rolled 80 pct at -78°C in the (110)[112] orientation, have been studied. The material was 99.999+ pct pure. The as-grown crystal had a resistance ratio, R(298°K) /R(4,2°K), of approximately 5000. The method of preparation of the single crystals and the necessary precautions in handling them during rolling in order to avoid recrystallization in locally distorted regions were described in an earlier paper.3 The rolling texture of this zone-refined aluminum crystal is practically identical with that of less pure aluminum crystals rolle! in the same orientation, i.e., a very sharp (110)[112]. Random nuclei were provided by artificial nuclea-tion by abrading one face of the rolled crystal on 600-grit emery paper in a random fashion. The thin layer of disturbed metal recrystallized at room temperature into fine grains almost immediately. An X-ray diffraction pattern was taken from the abraded surface of the crystal, and the relative intensities of the various reflection peaks were compared with those obtained from a powder specimen. As shown in Fig. 1, the relative intensities of the various peaks from the recrystallized layer varied somewhat from experiment to experiment, depending presumably on the abrading technique. However, the texture of the recrystallized layer produced by this method is nearly as random as that of a powder specimen. Fig. 2 shows the texture determined from the "growth side" (face opposite the abraded surface) of the crystal after it was kept at room temperature for 100 hr. The specimen was only partially recrystallized, as shown by the remaining deformation texture. The new orientations are derived from the recrystallized grains grown through the thickness of the crystal from the "nucleation side" (the abraded surface). Fig. 3 shows the texture after complete recrystallization produced by annealing at 60°C for 1 hr. The arrows indicate 40-deg rotation around the (111) axes of the deformation texture. The microstructure of a completely recrystallized specimen is shown in Fig. 4. As shown in Figs. 2 and 3, the texture of the re-crystallized grains grown from the nucleation side to the growth side is essentially the same as that observed in less pure aluminum&apos; (99.994 pct) and dilute alloyz crystals rolled in the same orientation. How-

Jan 1, 1968

By Alla D. Mah

Relatively pure aluminum nitride (AIN) recently has been prepared in sufficiently large amounts to permit a variety of experimental studies.4/ Several mechanical and chemical properties have been investigated, which show that this substance has good potentialities as a container for liquid aluminum, gallium, and boric oxide. The present work by the Bureau of Mines was undertaken to obtain thermo-dynamic data pertinent to the calculations regarding thermal stability and reactivity of aluminum nitride. The following measurements were made: (1) determination of the heat of formation from the elements at 298.15° K.; (2) determinations of the heat capacity over the temperature range 51° to 298.15° K. and the evaluation of the entropy at 298.15° K. therefrom; and (3) determinations of the heat content over the temperature range 298.15° to 1,800° K. These data are combined to obtain free energies of formation from the elements between 298.15° and 2,000° K.

Jan 1, 1961

By B. D. Cullity

GOLD and beryllium form an intermediate phase composed of the two metals in equal atomic proportions and having the formula AuBe. According to Winkler,1 this phase probably exists in two modifications: beta, stable below about 600°C, and alpha, stable between 600°C and the melting point, about 740°C. Only the low- temperature beta form will be considered in this paper. EXPERIMENTAL PROCEDURE The gold used had a fineness of 999.75 to 999.85, balance mainly silver, and the beryllium purity was better than 99.8 pct. AuBe was prepared by melting the constituents in a beryllia crucible in an evacuated induction furnace. After melting and thorough mixing, the alloy was allowed to solidify in the crucible. No difficulty was encountered in obtaining a homogeneous ingot. Both chemical analysis and microscopic examination of the ingot showed that the alloy was essentially one phase, containing 50 at. pct gold. The alloy was quite brittle and powder samples could easily be pre- pared by grinding in an agate mortar. The X-ray diffraction patterns were made with CuK-alpha radiation in a Debye-Scherrer powder camera 57.3 mm in diameter. The film was placed unsymmetrically in the camera in order to afford a basis for am-shrinkage corrections. A sample homogenized at 5o0°C gave the same diffraction pattern as a sample taken from the as-melted ingot. GENERAL RESULTS The diffraction pattern of AuBe was readily indexed on the basis of a simple cubic lattice with ao = 4.659 ± 0.001kX. If the unit cell is assumed to contain four AuBe "molecules," the calculated X-ray density is 13.46 grams per cc as compared with the measured density of 13.55 gm per cc. The only missing reflections in the dsraction pattern are (100) and (220) (see Table I). Since these missing reflections are not those characteristic of body or face centering, the lattice of AuBe must be simple cubic. To locate the positions of the atoms within the unit cell, two possibilities had to be considered: I. The structure was disordered and there were eight positions to be filled by an "average" gold-beryllium atom. That the structure was disordered seemed unlikely since the lattice parameter did not vary with composition in alloys containing some- what more and somewhat less than 50 at. pct gold, indicating a fixed composition for AuBe.

Jan 1, 1947

By B. D. Cullity

Gold and beryllium form an intermediate phase composed of the two metals in equal atomic proportions and having the formula AuBe. According to Winkler,&apos; this phase probably exists in two modifications: beta, stable below about 600°C, and alpha, stable between 600°C and the melting point, about 740°C. Only the low-temperature beta form will be considered in this paper. Experimental Procedure The gold used had a fineness of 999.75 to 999.85, balance mainly silver, and the beryllium purity was better than 99.8 pct. AuBe was prepared by melting the constituents in a beryllia crucible in an evacuated induction furnace. After melting and thorough mixing, the alloy was allowed to solidify in the crucible. No difficulty was encountered in obtaining a homogeneous ingot. Both chemical analysis and microscopic examination of the ingot showed that the alloy was essentially one phase, containing 50 at. pct gold. The alloy was quite brittle and powder samples could easily be prepared by grinding in an agate mortar. The X-ray diffraction patterns were made with CuK-alpha radiation in a Debye-Scherrer powder camera 57.3 mm in diameter. The film was placed unsym-metrically in the camera in order to afford a basis for film-shrinkage corrections. A sample homogenized at 500°C gave the same diffraction pattern as a sample taken from the as-melted ingot. General Results The diffraction pattern of AuBe was readily indexed on the basis of a simple cubic lattice with ao = 4.659 + 0.001kX. If the unit cell is assumed to contain four AuBe "molecules," the calculated X-ray density is 13.46 grams per cc as compared with the measured density of 13.55 gm per cc. The only missing reflections in the diffraction pattern are (100) and (220) (see Table I). Since these missing reflections are not those characteristic of body or face centering, the lattice of AuBe must be simple cubic. To locate the positions of the atoms within the unit cell, two possibilities had to be considered: I. The structure was disordered and there were eight positions to be filled by an LCaverage" gold-beryllium atom. That the structure was disordered seemed unlikely since the lattice parameter did not vary with composition in alloys containing sornewhat more and somewhat less than so at&apos; pet gold, indicating a fixed composition for AuBe.

Jan 1, 1947

By B. D. Cullity

Gold and beryllium form an intermediate phase composed of the two metals in equal atomic proportions and having the formula AuBe. According to Winkler,&apos; this phase probably exists in two modifications: beta, stable below about 600°C, and alpha, stable between 600°C and the melting point, about 740°C. Only the low-temperature beta form will be considered in this paper. Experimental Procedure The gold used had a fineness of 999.75 to 999.85, balance mainly silver, and the beryllium purity was better than 99.8 pct. AuBe was prepared by melting the constituents in a beryllia crucible in an evacuated induction furnace. After melting and thorough mixing, the alloy was allowed to solidify in the crucible. No difficulty was encountered in obtaining a homogeneous ingot. Both chemical analysis and microscopic examination of the ingot showed that the alloy was essentially one phase, containing 50 at. pct gold. The alloy was quite brittle and powder samples could easily be prepared by grinding in an agate mortar. The X-ray diffraction patterns were made with CuK-alpha radiation in a Debye-Scherrer powder camera 57.3 mm in diameter. The film was placed unsym-metrically in the camera in order to afford a basis for film-shrinkage corrections. A sample homogenized at 500°C gave the same diffraction pattern as a sample taken from the as-melted ingot. General Results The diffraction pattern of AuBe was readily indexed on the basis of a simple cubic lattice with ao = 4.659 + 0.001kX. If the unit cell is assumed to contain four AuBe "molecules," the calculated X-ray density is 13.46 grams per cc as compared with the measured density of 13.55 gm per cc. The only missing reflections in the diffraction pattern are (100) and (220) (see Table I). Since these missing reflections are not those characteristic of body or face centering, the lattice of AuBe must be simple cubic. To locate the positions of the atoms within the unit cell, two possibilities had to be considered: I. The structure was disordered and there were eight positions to be filled by an LCaverage" gold-beryllium atom. That the structure was disordered seemed unlikely since the lattice parameter did not vary with composition in alloys containing sornewhat more and somewhat less than so at&apos; pet gold, indicating a fixed composition for AuBe.

Jan 1, 1947

By D. P. Lindroth

The thermal diffusivity above 298 K, up to 400 K and at pressures from ambient down to 10-10 torr was measured for the following six igneous rock types: tholeiitic basalt, obsidian, rhyolite, granodiorite, gabbro, and dunite. The thermal diffusivity data were obtained by the laser flash method. Extreme values range from a low of 5.3X10-3 cmF/sec for rhyolite at 383 K and 2X10-7 torr to a high of 13.4x10-3 crrf/sec for dunite at 298 K and 760 torr. Data are also included on a commercial fused silica, which was used as a standard.

Jan 1, 1974

By R. W. Gould, B. G. LeFevre, A. G. Guy

The resistivity anomaly known as the K state was studied in Ni-Mo alloys containing 10.5 and 14.0 at. pct Mo. Both these alloys exhibit a large K effect which depends on the mechanical and thermal treatment. On the basis of X-ray diffuse scattering studies which were correlated with resistivity measurements, it appears that the K state in dilute Ni-Mo alloys can be associated with changes in the degree of short-range order within the a phase. An interesting phenomenon that has received much attention in recent years is the K state. The K state is marked by anomalous changes in some of the physical properties of certain alloys without the occurrence of observable microscopic structural changes. One of the early pieces of work in this area was by Thomas&apos; who studied alloys of Ni-Cr, Ni-Cu, Ni-Cu-Zn, Fe-A1, Fe-Si, and Ni-A1. Upon annealing specimens which had been previously cold-worked or quenched from an elevated temperature he found an anomalous increase in resistivity over a certain temperature range. He also found that specimens which had been appropriately annealed to develop the K state showed a decrease in resistivity upon subsequent cold working. These effects are opposite to those found in normal alloys. Although the resistivity anomaly has been rather arbitrarily taken as the "definition" of the K state, there are several other interesting effects which accompany the resistivity increase. In Ni-Cr alloys,2, 3 for example, it was found that the hardness increases with increasing resistivity. It was also found that specimens which have been treated to develop the K state can be cold-worked for as much as an 80 pct reduction of area without an increase in the hardness. In Fe-A1 and Fe-Si alloys4 the K state is accompanied by an increase in flow stress and by a lattice contraction. In Ni-A1 alloys,5 specimens which have been treated to develop the K state also show an increase in elastic modulus. In Ag-Pd alloys6 the increased resistivity observed on annealing a cold-worked specimen is accompanied by an increase in the thermoelectric power and an increase in the Hall coefficient. The explanations of the K state phenomena are varied and depend upon the particular alloy in question. Several theories have been advanced to explain the increased conductivity with cold work on the basis of changes in the electronic configuration of the alloy as a result of local lattice distortions.7"9 Most investigators, however, believe that some type of local order in the solid solution, either short-range order (SRO) or clustering, is responsible for this effect. Theories concerning the relationship between ordering and the K state have for the most part been speculative, since there is little direct X-ray evidence that can be correlated with the above property changes. Much of the previous work on the K state was done in the Ni-Cr system where the small difference in the X-ray atomic scattering factors of the components nickel and chromium makes it very difficult to use X-ray diffuse-scattering measurements to determine the role of local order. In the Ni-A1 system, however, Starke et al.10 succeeded in detecting a connection between local order and the K state. It was found that a small but measurable K effect correlated with increasing SRO in the nickel-rich a phase. The manner in which local order might increase the resistivity of K state alloys is not completely clear. Since most of the known K state alloys contain at least one transition element, significance has frequently been attached to the presence of an unfilled d shell. It has been suggested that during the formation of the K state the number of conduction electrons decreases as a result of the transfer of s electrons to the d shell where they are more tightly bound.1&apos;11&apos;12 Koster and Rocholl13 have proposed that SRO can cause an increase in resistivity for alloy systems in which the number and mobility of charge carriers are reduced when the percent solute is increased. According to this hypothesis, the local environment of a given solvent atom changes in the same manner with increasing percent solute as it does with an increasing degree of SRO; hence the change in physical properties should tend in the same direction. In this hypothesis, SRO is considered only in a statistical sense, and the increased resistivity of the K state is attributed to a change in the mean distribution of electrons and holes in the s and d states as a result of SRO. From the work of Chen and Nicholson on Ag-Pd alloys,6 it appears that the K state can occur in systems for which the d shell is completely filled. These investigators explained the increased resistivity by picturing the SRO as small domains of some form of long-range order (LRO). According to ~ibson,&apos;~ the number of effective electrons can be reduced by the creation of a new Brillouin zone boundary near the Fermi surface of an alloy as a result of the changing crystallographic symmetry that accompanies the formation of a superlattice. This idea may be expressed in terms of the superzone concept.15 In the present work the role of local order in the formation of the K state in Ni-Mo alloys was investigated. The principal tools used in this study were X-ray diffuse scattering and electrical resistivity measurements; however, these data were supplemented by electron microscopic and field-ion microscopic data. The purpose of the work was to determine whether or not the K state in Ni-Mo alloys can indeed be attributed to the formation of SRO as has been proposed by previous investigators.

Jan 1, 1969

By A. Niedbalska

A technique was developed for low pressure diamond synthesis using a specially prepared glassy carbon. The glassy carbon is produced in the form of solid plates by slow carbonization process of phenol formaldehyde resin interrupted at temperatures of 1273° K, 1073° K, and 923° K. The diamond synthesis is performed using cobalt plates as a catalyst-solvent. Synthesis was carried out at applied pressures of 6.0 and 3.8 GPa and a temperature of 1970° K and 2170° K. The products were analyzed by the x-ray powder dyffraction method and by microprobe to identify the diamond crystallites. The samples were taken from the place of contact between the glassy carbon and the cobalt catalyst. They were separated from the charge by a chemical method. Diamonds of the best form and the largest size were obtained at pressures that were much lower than those used in the traditional graphite method of diamond synthesis.

Jan 1, 1989

By W. W. Weller

Heat capacities of ferrous molybdate (FeMoO4) and ferrous tungstate (FeWO4) were measured over the temperature range 51° to 298.15° K. Ferrous tungstate has a lambda type heat capacity anomaly at 74° K. The 298.15° K entropies are FeMoO4) 30.9 cal/deg mole and FeW04) 31.5 cal/deg mole.

Jan 1, 1966