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By L. W. Saperstein, Y. J. Wang

The development of computer programs to solve complex ventilation networks has reached a point of refinement where these programs become a necessary tool of the ventilating engineer. Such a program is presented. This program has an increased efficiency over its predecessors; it also includes the important facility of fixed quantity branches. This is in addition to the earlier capabilities of free-splitting, internal or external fans, and natural ventilation pressures. The program is written in Fortran IV with double-precision arithmetic and is listed here. Sample problems with input and output are also given. The recent, and presumably continuing, official interest in mine safety is bound to result in more stringent control requirements for ventilation. Anticipating this requirement, the Dept. of Mining, The Pennsylvania State University, has developed a computer program for the solution of complex mine ventilation networks. It is the opinion of this department that the program has sufficient merit to warrant its use in normal industrial operating conditions. Although the undergraduate mining engineering students of this department are routinely trained to use the program, it was felt that this training did not present a fast enough means of dissemination; consequently, this paper was written. The conception of the program is not new;1; however, the present version incorporates features that changes it from a scientific endeavor to a useful operating tool. These features include those of earlier programs, such as free-splitting, external or internal fans, and natural ventilation processes plus the important ability of handling fixed quantity branches, and of increased efficiency giving reduced cost. The facility for fixed quantities will be discussed in more detail. The solution of any network depends upon satisfying Kirchchoff's laws. These state that the sum of flow at any junction is zero, and that the sum of the pressure drop in branches totals the pressure drop in the mesh constituted by those branches, that is that the pressure drop around any mesh is zero. The terminology is derived from Synge," and is, in fact, that used by Topologists. A branch is any airway of uniform characteristics; a junction is where two or more airways meet (the minimum number of two airways is useful for describing the point where a low resistance airway changes into a high resistance one due to some physical constriction); a mesh is a path, along some branches, that returns on to itself, but does not need to traverse the same branch twice in order to do so; and a tree describes all those branches in a network which, while connected through all the junctions, do not form any meshes. Obviously the branches-in-tree are open-ended and can be shown to be one less than the number of junction (J-1). The branches-out-of-tree, also called basic branches, are equal to the number of basic meshes and are M = N — J + 1, where M is the number of meshes and N is the number of branches. The major difference between this network and an electrical one is that the law attributed to Atkinson for pressure drop is utilized rather than Ohm's law. Atkinson stated that the pressure drop (potential) is equal to a resistance factor multiplied by the quantity squared. This is written H = RlQlQ where H is the head loss. Q? is written in this factored fashion so that H will always have the same sign as Q. A negative sign is used to indicate that flow in a branch is opposite to that of its containing mesh. A mesh takes the direction of its basic branch. Utilizing Atkinson's law in Kirchchoff's second law and bearing in mind that the first law must remain satisfied, the program will determine all quantities, or head losses as they are related, for the branches of the network. The program uses a Geuss-Seidel form of in-teration, starting from an arbitrary Q and using a correction factor of the type first suggested by Professor Hardy Cross' that will handle the nonlinear equation of Atkinson's law. A fuller description of the interative process is included in Wang and Hartman.' Iteration continues until the correction factor is less than or equal to a preset error (E) or until a maximum number of iterations (MAXIT) have been reached. Utilizing an E of 50 cfm, sample problems have had rapid solutions. Rapidity of solution is also ensured by the method of selecting meshes and basic branches. This is done by the computer, which chooses the highest resistance branches, fixed quantity branches, and those containing fans, as basic branches. Natural ventilation pressures are handled as fans of constant pressure and may be assigned to any branch. The computer will generate a fan characteristic by polynominal fitting if a few operating points from the desired fan are input. A fan may be placed in any branch of the network, except those with fixed quantities. The strength of the present program over its earlier versions is its ability to handle fixed quantity branches. This means that certain branches can have their quantity determined, or preset, by the ventilating engineer. Operating under this constraint, the program will analyze the network and output the pressure change necessary to achieve this desired quantity. A call for positive head loss would require an auxiliary fan; negative head loss would require a regulator. This represents a powerful tool for the ventilating engineer who must insure that certain quantities of air pass across the operating faces or through working stopes. McPherson6 suggested that a ventilation program would be useful for determining in advance which branches were receiving inadequate air, and that subsequent calculations would then indicate the necessary corrections. The present program makes these corrections. Thus the mine manager can know immediately the consequences of changes to his system;

Jan 1, 1971

By C. B. Manula

This report is volume 5 of an eleven-volume final report for a Grant completed by The Pennsylvania State University for the Bureau of Mines to develop a dynamic general purpose computer simulation model for underground coal mining. The simulator developed under the Grant has a wide range of possible applications in the design and evaluation of underground coal mining systems from the standpoints of health and safety, productivity, and mine environment. The Master Design Simulator consists of nine free-standing function-oriented subsystem models linked together into an integrated, comprehensive model of an entire mining system. The subsystems comprising the Master Design Simulator include geology and reserves, methane generation, roof support, subsidence, production, rail haulage, cost, water generation, and ventilation. This volume of the final report describes the Subsidence subsystem of the Master Design Simulator. This subsystem uses the finite element method to simulate ground movement caused by an advancing mining face in an underground coal mine. The model is structured around the concept that the ground mass is a layered media of different thicknesses and material properties. Gob areas are considered to be unfilled with broken rock at the beginning of mining. After mining starts, the gob areas are filled with broken material which provides resistance to the lowering of the overlying strata. The model is tied to the Geology and Reserves subsystem, but it can also operate as a free standing routine. A case study is presented illustrating the application of the model. A detailed user's manual, including a program listing, is included in the appendices.

Jan 1, 1974

By Simon D. Strauss

This article is part of a presentation Strauss made Dec. 3, 1988 at the sixth annual American Mining Hall of Fame (AMHofF) banquet in Tucson, AZ. At the banquet, the Mining Club of the Southwest honored Strauss by naming him the 1988 recipient to the AMHofF. (Industry newswatch, page 6). Introduction In the last two decades, there have been two previous periods of sharply rising copper prices. One was in 1973-74; the other in 1979-80. In both cases, a short period of high copper prices was followed by steep and prolonged periods of low prices. In the language of the market analyst, the high prices were called spikes. Is it likely, then, that three or four years from now the present price of copper will show up as a spike in the charts? There are substantial differences between the current position and the markets of 1973-74 and 1979-80. In 1973-74, US copper prices were subject to government-imposed price ceilings of $0.60 per lb through early December, 1973. These were increased to $0.68 per lb through May, 1974. Meanwhile, mines outside the United States were free of restraint. Rising demand in London pushed the price above $1 per lb in the third quarter of 1973. In early April 1974, it reached a peak of $1.50 a lb. In May of that year, the US ceilings were removed and the domestic price began to rise. What were the causes of this price strength? While ceilings held the domestic price below the world market, domestic consumers acquired all the tonnage offered by US mines. Meanwhile, in the world market, buyers feared prolonged shortages. The Club of Rome had warned that world economic growth would be limited by impending depletion of mineral resources. There was al so concern about a copper cartel that might increase copper prices even more. In early 1974, at a special United Nations session, developing-country spokesmen warned the industrialized world to expect high commodity prices. Also adding to the concerns of copper buyers was the fear that when US wage contracts expired in mid-1974, protracted strikes might ensue. The combined effect of this atmosphere was that consumers accumulated large amounts of copper, far above their normal carry. By midsummer, the position changed radically. While some strikes did occur in the United States, they proved short-lived. The most significant development was a massive economic recession. The energy crisis was causing enormous dislocations in the world economy. Construction activities were sharply reduced and automobile sales plummeted. Public confidence was shaken. Copper sales fell precipitately as consumers decided to draw down their swollen inventories. The price came down too. In April 1974, the London price- as noted-was $1.50 per lb. By the end of the year, it was down to $0.57 per lb. In 1973-74, the copper price rise was fed by fears of shortages, cartels, and strikes. In 1979-80, the market rose on the strength of good demand plus political and military instability. The price rise was also fed by inflationary trends, exemplified by speculation in precious metals. Prior to this second spike in copper prices, the market had been depressed despite economic recovery. Despite reduced US output, copper production in the rest of the world was increasing. Though consumption was rising, surplus stocks reached unprecedented levels. At the end of 1977, the aggregate of stocks at the refineries and in the exchange warehouses was 1.5 Mt (1.67 million st). That was more than three months' shipments at then rates of consumption. In 1978, the statistical position improved. Shipments of refined copper that year were 8% greater than in 1977. The result was reduced visible stocks but -just as was to be the case in 1984-86 - the price remained depressed. In 1979, an even sharper drop in visible stocks occurred as shipments rose by another 4%. This time, though, the price responded. In December 1978, the price in current dollars had been less than $0.70 per lb. By the end of 1979, it was above $1 per lb. And in the first quarter of 1980, it averaged more than $1.20 per lb. Military and political developments were a factor. The hostage crisis in Iran was followed by the Soviet invasion of Afghanistan. To many, it appeared that war might result if the Soviets pushed on into Pakistan. At the same time, speculation drove gold and silver prices to incredible levels. As silver prices rose above $40 per oz), copper above $1 per lb did not seem unreasonable. Then, in March 1980, the precious metals price boom collapsed. Copper prices also retreated, though not as much. In February 1980, the copperprice averaged $1.32 per lb. By December, it had dropped to $0.85 per lb. In 1981 and 1982, the decline continued. The present copper situation Why is the current situation different from these two previous episodes of high prices? First of all, today's inventories are far lower than they were in either 1974 or 1980, when the price declines began. The stocks at refineries and in exchange warehouses are substantially less than they were then. More significantly, consumers today have very limited holdings compared with stocks they carried in 1974 and 1980. Even if faced with a downturn in their own business, copper consumers will not be able to stay out of the market. To the extent they have any business at all, they will have to buy copper. And on the supply side, there is also a significant difference. The year 1988 has been unusual. US production is up sharply, responding to higher prices. But production is little changed in the rest of the Free World. Indeed, several of the major producing countries have mined less copper this year than in 1987 despite the prices increases. Peru has had labor problems that have cut its output. Zambia and Zaire suffer from shortages of expatriate technicians, not to mention spare parts. They also suffer from the failure to maintain equipment or perform adequate development. In Chile, the two largest mines, Chuquicamata and El Teniente, have not met their production targets. Chuquicamata is struggling with high arsenic content in its ore. And, in late 1987, severe rockbursts hampered operations at El Teniente. In addition, depletion of some mines in Australia and South Africa has offset rising output at other properties. Indeed, when final mine output figures for 1988 become available, it may well be

Jan 1, 1989

By A. E. Bibb, F. W. Kunz

A platelet form of zirconium hydride was found in zirconium and ZY-1 wt pct U single crystals containing hydvogen in the range of 50 to 100 ppm. The habit planes for the hydride plateletg in the zirconium crystals were found to be {1012}, {1121}, and (1122) while the habit planes in the ZY-1 wt pct U crystals were {1121), {1123}, and (1.231). With the exception of the {1231} plane, these have been identified as the twin planes in zirconium. The different effect of hydridc on the mechanical puoperlies of zirconium at high and low strain rates has been explained in terms of the defornzation mechanism and the habit planes of the hydride. ZIRCONIUM and zirconium-base alloys are considered for many nuclear reactor applications, where the pick-up of hydrogen is a distinct possibility, if not an unavoidable occurrence. An excellent example of this is the absorption of hydrogen by zirconium and its alloys during aqueous corrosion, where as much as 30 pct of the hydrogen produced by the corrosion reaction is picked up by the metal. Under the proper conditions, the absorbed hydrogen would precipitate as zirconium hydride and consequently alter the mechanical and physical properties of the zirconium-base alloys. The knowledge obtained from a study of the crystallographic aspects of the hydride precipitation is necessary for the interpretation of the observed changes in the mechanical and physical properties. The deleterious effect of hydrogen on zirconium-base materials is not always observed in normal slow tensile loadings at or above room temperature, but becomes immediately obvious under impact loading, where the energy absorbed decreases with increasing hydrogen content,' and in notched specimens where triaxial loadings are attained.2 This effect is of serious consequence in the applications of Zr where shock loads or high strain rates are experienced. The purpose of this investigation was to determine the crystallographic planes of the Zr and Zr-1 wt pct U matrices upon which zirconium hydride will precipitate and to correlate these with the observed mechanical property data. PROCEDURE Crystal Growth—Single crystals of Zr and Zr-1 wt pct U alloy were used in this investigation. The Zr crystals were made from sponge that had a typical analysis as shown in Table I. Large grains were grown in this material by rapid cooling an arc-melted charge in a button furnace. This operation undoubtedly lowered the volatile impurity concentration below those listed in Table I. Back-reflection Laue patterns of these large grains showed the presence of considerable substructure and lattice distortion within each grain. It was found that an 18-day anneal in an argon atmosphere at 840°C (a high a heat-treatment) followed by furnace cooling, resulted in the further growth of the large grains and removed the detectable substructure ana lattice distortion. Equiaxed crystals with a cube edge of approximately 1/4 in. were produced in this manner. Large crystals of Zircaloy-2 (1.5 wt pct Sn, 0.15 wt pct Fe, 0.1 wt pct Cr, 0.05 wt pct Ni) also were produced by this technique but the large amount of substructure formed could not be removed by annealing. The Zr-1 wt pct U solid solution crystals were produced from an alloy rod consisting of Westing-house-Foote hafnium-free crystal bar zirconium and depleted uranium. The chemical analysis of the alloy is also given in Table I. Three-inch sections were cut from an 0.120-in.-diameter alloy rod and tapered to a fine point on each end. The samples were sealed in evacuated Vycor tubes, homogenized for 5 hr at 1000°C in the ß-phase field, and slowly lowered in a gradient furnace into the a phase (800°C), annealed for 60 hr at this temperature to produce grain growth and furnace cooled. This technique produced large grains approximately 1/2 to 3/4 in. long over the entire cross section of the rod, which were free of observable substructure. Metallographic examination up to X1000 did not show the existence of any second phase in these grains. Crystal Orientation—All crystals were abraded on fine emery paper to produce two flat surfaces 90 deg to each other. The crystals were analyzed by the standard Back-Reflection Laue technique to obtain the crystal orientation with respect to the abraded surfaces. The Back-Reflection Laue Photograms of all crystals used were clear and sharp indicating that the crystals were free from any substructure that could be detected by these means. Hydriding—Only samples containing four to five grains were used for hydriding. Samples that exhibited fine grained patches were discarded. The samples were hydrided using the following technique. After etching in 48 vol. pct HNO3, 48 vol. pct H2O, and 4 vol. pct HF (48 pct) solution, the specimens

Jan 1, 1961

By W. C. Ellis, E. S. Greiner, M. E. Fine

Discontinuous changes of Young's modulus, internal friction, coefficient of expansion, electrical resistivity, and thermoelectric power are evidence for a transition in chromium near 37OC. Although the X-ray diffraction pattern gives no clue, a difference between the thermal expansivity and the temperature dependence of the lattice parameter suggests a crystal-lographic change. Young's modulus data disclosed another transition near THE thermal dependence of a number of proper- ties of chromium indicates a transition occurring over a temperature range near room temperature. Bridgman' noted this first from a minimum in the electrical resistivity near 12°C. In a sample of greater purity, Sochtig' observed the minimum at 41 °C. Erfling3 reported an inflection in the thermal expansivity curve at 36°C. These temperature-dependence curves are reversible with no hysteresis being detected. No discontinuity or inflection has been observed in the heat capacity' or the paramagnetic susceptibility.' Likewise, no one has noted a change in crystal symmetry. In the present investigation Young's modulus, internal friction, coefficient of expansion, electrical resistivity, illustrated in Fig. 1, lattice constant, Fig. 2, thermal electromotive force, Fig. 3, and paramagnetic susceptibility, Fig. 4, were measured over an extended temperature range. The samples were prepared in two ways: (1) By cold pressing a sintered electrolytic powder compact,' and (2) by electroforming from an aqueous solution. The electroformed samples were prepared by R. A. Ehrhardt and G. Bittrich by plating on copper or nickel tubes from an aqueous solution according to the method of Brenner, Burkhead, and Jennings.' The pressed powder samples (method 1) were finally annealed at 1400°C in purified helium: the electroformed samples, packed in powdered chromium, were vacuum-annealed at 1000°C. From the composition of the original powder' the purity of the pressed powder samples is estimated to be 99.8 pct Cr. Spectrochemical analysis furnished by E. K. Jaycox, revealed only slight traces of impurities in the electroformed sample (less than 0.001 pct), neither iron nor nickel being detected. Chromium deposited by this method is reported to contain approximately 0.05 pct 0.- Methods and Data Young's Modulus: For measurement of Young's modulus, an annealed, pressed powder sample, 0.114 x0.237x1.845 in., and an electroformed sample, 0.10 in. od, 0.01 in. id, and 1.86 in. long, were prepared. The resonant frequencies of the rod samples in forced longitudinal vibration were measured at a series of temperatures from —192" to 200°C by a method previously de~cribed.6,7 Because the samples were nonferromagnetic, iron- silicon or molybdenum permalloy tips (0.013 in. thick) were soldered to the ends. The softening temperature of the solder limited the temperature of measurement. Young's modulus, E,, of chromium at temperature, T, may be calculated from the resonant frequency of the composite rod, f,; the thickness of the tips, t; the length of the chromium sample at 25°C, l,.; the density of chromium at 25°C, pc (7.20 g per cc);5 and the thermal expansivity, Al/l25. Modulus differences for two temperatures, ET and E, are accurate to +0.002x10" dynes per cm2. ET = 4PAlc+2tyf Young's modulus at 25°C, Fig. la, is 28.2~10" dynes per cm' (40.8xlO" si) in wrought chromium (upper curve). The modulus of the electroformed sample is apparently lower due to cracks. From the modulus measurements two transitions were observed: one with a critical temperature at 37°C, the other at —152°C. Internal Friction: The internal friction, 1/Q, was determined from the width, Af, of the strain amplitude-frequency curve at 0.707 times the strain amplitude at resonance;' the internal friction, 1/Q, then equals Af/f. Fig. lb shows a sharp peak in internal friction at 38°C. The internal friction of the electroformed sample had a similar maximum. Thermal Expansion: The expansivity measurements, shown in Fig. 2, covering the temperature range —195" to +400°C were made by D. MacNair in an interferometric dilatometer8 sing a sample consisting of three pyramids 0.25 in. high prepared from wrought electrolytic chromium. Below —120°C the experimental points deviated up to ±2x10 from the drawn expansivity curve in Fig. 2 because of decreased precision of the quartz wedge thermometer at low temperatures. Near 38°C the thermal expansivity curve, Fig. 2, goes through an inflection corresponding to a minimum in the coefficient of expansion, Fig. ld, and a relative volume decrease on heating. Electrical Resistivity: The variation of electrical resistivity with temperature measured by the po-tentiometric method is also shown in Fig. l. The resistivity of the wrought chromium (upper curve) at 20°C is 13.6 microhm-cm; of electroformed chromium (lower curve) 12.8 microhm-cm. The lower value reflects higher purity and agrees closely with a published value." A minimum at 40°C occurs in the resistivity curves of both samples. No conclusive evidence for a transition near — 150°C was observed, but the points, Fig. 1, appear to deviate from a smooth curve between —120" and —160°C.

Jan 1, 1952

By R. P. Abendroth

The hydrogen-water vapor ratio at which Fe-Cr alloy, chromium oxide, andiron chromite coexist in equilibrium was determined between 900" and 1200°C. A thermogravimetric method was used to determine equilibrium conditions. The results fit a straight-line relationship in the temperature region studied, and are given by Reduction experiments were also performed to confirm the results of the equilibrium investigation. ThE oxygen pressures at which Fe-Cr alloy, chromium oxide, and iron chromite coexist in equilibrium have been previously determined by Boericke and angert,' Morozov and Novokharski,' and Katsura and uan. Only one determination (at 1300°C) was made by Katsura and Muan, but it agrees with the results of orozov and Novokharski. The results of Boericke and Bangert, however, differ appreciably from the results of these investigators. Previous studies have assumed that the equilibrium metallic phase is pure iron, but Dahl and Van vlack have shown that the iron contains from about 1 wt pct Cr at 1000°C to over 2 wt pct above 1300°C. The chromium oxide also contains a small amount of iron in solid solution. In the present study, hydrogen-water vapor mixtures were equilibrated with the condensed phases, using a therrnogravimetric method to determine equilibrium conditions. The reaction can be written EXPERIMENTAL General Procedure. The starting material was a sintered pellet of Fe2O3-Cr2O3 solid solution with a hole in the center, and was placed on a fused silica hook. This assembly was raised into the preheated hot zone of the furnace in a helium atmosphere, hooked onto a fused silica hangdown suspended from one arm of an Ainsworth Model RV-AU-1 recording balance, and the starting weight determined. A flowing hydrogen-water vapor atmosphere was then exchanged for the helium by evacuation, and the sample reduced until the weight loss indicated the sample composition to be in the alloy-Cr2O3-chromite field. The tem- perature was adjusted incrementally until constant sample weight was achieved for several hours, to within 0.02 mg. A hydrogen-water vapor atmosphere of different composition was then admitted, and the same procedure carried out. At the end of a series of determinations, the sample was examined by X-ray diffraction to verify the presence of the desired phases. Microscopic examination of the silica hook showed no interaction with the sample, nor did it lose any weight. Several criteria were used to insure equilibrium besides constancy of weight. For a given hydrogen-water vapor composition, equilibrium was approached from both oxidizing and reducing sides by varying the furnace temperature slightly. The resulting slow weight loss or gain was observed for several hours. Constant weight could be re-established by returning to the original furnace temperature. The last criterion used was varying the relative amounts of the phases by further reduction or oxidation, and observing any changes in temperature required for constant weight for a given hydrogen-water vapor atmosphere. None were observed. This procedure was essentially the same as approaching the equilibrium from oxidizing and reducing sides, but larger weight excursions were carried out. Sample Preparation. Reagent-grade Fe2O3 and Cr83 powders were mixed in the desired proportions and heated in air at 1250°C for 2 hr. The mixture was re-ground and heated in air overnight at 1250°C. X-ray diffraction showed complete solid-solution formation as a result of this procedure. The solid solution was then pressed into l/2-in.-diam pellets using Carbowax 4000 as a binder. The hole was drilled in the center, and the pellets were sintered 24 hr at 1250°C in air on a bed of Fe2O3-Cr2O3 of the same composition, contained in an alundum boat. After cooling, the pellet surfaces were abraded with 310 paper to remove any surface compositional differences, such as loss of Cr2O3. Chemical analysis of the sintered pellets was 67.16 wt pct CrP3 and 33.02 wt pct Fe203. Atmosphere Generation and Control. The hydrogen-water vapor atmospheres were generated by passing Matheson ultrahigh-purity hydrogen, with no further purification, through two water bubblers contained in a constant-temperature water bath. Since the water-vapor dew points required in this study were below room temperature, the bath was insulated, and was cooled by thermoelectric-immersion devices. The bath temperature was controlled to 0.0l0C. Since rather high flow rates of about 900 ml per min were used through the furnace tube, an independent check of the dew point was made to insure saturation of the hydrogen by the water vapor. Although the dew point could only be determined to within 1/2"C, the determined dew points agreed with the water-bath temper-

Jan 1, 1967

By W. A. Backofen, R. L. Fleischer

Single crystal, bicrystal, and polycrystal tensile tests of aluminum at 4.2°K, 77°K, and 300°K have been used to examine the role of grain boundaries in the deformation process. Results indicate that a grain boundary may affect the extent and slope of easy glide. The stage II hardening rate, on the other hand, is independent of the presence or absence of grain boundaries. This conclusion allows the size of the region of multiple slip caused by an incompatible grain boundary to be determined. For the size of bicyystal sample used in this study, multiple slip occurs in about half of the cross section. PREVIOUS studies of the stress-strain characteristics of bicrystals of face-centered-cubic metals have been limited to aluminuml-5 at room temperature. Recent results, however, indicate that the stress-strain curves of single crystals of such metals may be separated into at least three stages6 in which different deformation processes are occurring7 provided testing is done at sufficiently low temperatures.' Since for aluminum a well-defined stage II develops only below room temperature, previous studies have not been able to relate effects of grain boundaries to all of the three stages of deformation. It is therefore to be expected that low-temperature deformation of aluminum single crystals and bicrystals should clarify the effects of grain boundaries on the different processes of deformation. EXPERIMENTAL PROCEDURE Single crystals and bicrystals were grown from the melt by the standard techniqueg with aluminum reported by Alcoa to be 99.993 pct pure. Ridges in the boat were used to guide the grain boundary during growth, assuring that the boundary would bisect the sample.10 The rate of furnace motion during growth was 1.0 cm per hr. During growth zone purification resulted, as evidenced by the ability of the first material to freeze to recrystallize at room temperature following severe deformation. Samples were approximately 4.4 X 6.6 mm in cross section and 103.5 mm in length between grips. Samples were annealed at 635" i 5°C for 40 hr and furnace cooled over a 7-hr period. They were then electropolished in a solution of 5 parts methanol to 1 part perchloric acid at a current density of 15 amp per sq dm for about 30 min at temperatures below 0°C. Tensile testing was performed at 295" (room temperature), 77" (sample in liquid nitrogen), and 4.2"K (sample in liquid helium) on the hard-type machine indicated schematically in Fig. 1. The machine con- sists basically of a tube surrounding a rod; one end of the sample is attached to each member, and the rod is pulled up the tube to extend the sample. The rod is rigidly mounted and is moved vertically by a system described by asinski." The pulling force is measured continuously by an electrical strain gage load cell, and the relative displacement of the tube and rod is also recorded continuously by a soft cantilever beam with electrical strain gages. Maximum stress and strain sensitivities were ±2g per sq mm and * 3-10-5. In all tests the strain rate was approximately 5.10-5 per sec. The thin wires in the tensile apparatus introduce softness, which may be corrected for, however, by measuring load vs displacement with the sample replaced by an elastic member. For loads greater than 15 kg the spring constant is 1.875.106 g per cm. The flexible wires also served to reduce substantially the large shearing forces which may arise in the case of grips having horizontal rigidity.'' As in any gripping system, however, bending moments will arise in the course of deformation by single slip. Engineering stress, s = (load)/(original cross-sectional area), and strain, E = (increase in length)/ (original length), are used for stress-strain curves unless otherwise indicated. Tables list resolved shear stress, T=mo and shear strain ? = dm, where m is the usual Schmid resolved shear stress factor for the primary slip system at the start of deformation. The first group of samples to be described forms an isoaxial set, all of the crystals making up the single crystals or bicrystals having the same tensile axis, the orientation of which is indicated by the cross in Fig. 2. For this orientation the primary slip plane and slip direction make angles of 45 deg with the tensile axis and the Schmid factor m has its maximum possible value of 0.5. Rotations about the tensile axis are indexed by means of an angle 0 between the small-area surface of the samples and the projection of the primary slip direction onto the cross section, as defined in Fig. 3. In single crystals, values of 0 were 0 and 90 deg, while in bicrystals 0 values were (0 deg, 180 deg), (90 deg, 270 deg), and (0 deg, 90 deg) as indicated in Fig. 4.

Jan 1, 1961

By G. Bhat, J. F. Libsch

lsoembrittlement curves depicting the influence of time and temperature in the range 800' to 1260°F (425' to 680°C) on the development of embrittlement in a commercial chromium alloy steel and a commercial molybdenum alloy steel are presented. Two distinct regions of embrittlement occur in the chromium alloy steel: I—at 800' to 1000°F (425' to 540°C) and 2—in the region just below the lower critical temperature. Embrittlement is most pronounced at 800' to 1000°F, decreasing very rapidly with increasing temperature above this region, only to increase again as the lower critical temperature is approached. The data suggest two distinct modes of embrittlement with possible superposition of the two modes at extended embrittling times in the temperature range 1100° to 1150°F (590' to 620°C). While the molybdenum alloy steel shows little susceptibility to embrittlement at 800' to 1000°F (425' to 540°C), considerable embrittlement may occur just below the lower critical temperature. THE subject of temper embrittlement in alloy steels has received considerable attention in the last few years. Points of view on the mechanism of embrittlement differ, however, resulting in part from the incompleteness of the data developed and in part from the speculation regarding the susceptibility of plain carbon steel to temper embrittlement. Libsch, Powers, and Bhat1 carried out short-time embrittling treatments on an AISI 1050 steel and demonstrated that hardened plain carbon steels are quite susceptible to embrittlement when tempered in the range from 850°F (455°C) to the lower critical temperature. The isoembrittlement diagram,' representing the embrittling characteristics of this steel, is reproduced in Fig. 1. It is evident from the shape of the curves shown that embrittlement in plain carbon steel increases progressively with both temperature and time in the embrittling range. A comparison of the isoembrittlement diagram for AISI 1050 steel with that presented by Jaffe and Buffum' for an SAE 3140 steel shows that up to 930°F (500°C) the isoembrittlement characteristics of the plain carbon steel are similar to those of SAE 3140 steel, although the embrittlement is much more severe in the latter steel. Above 930°F (500°C), the rate of embrittlement in the plain carbon steel increases continuously with increasing temperature; whereas, in the SAE 3140 steel, the embrittlement rapidly decreases. The influence of alloying elements upon embrittlement during tempering thus appears to cause a decrease in embrittlement above the region of maximum embrittlement, i.e., 850" to 1000°F. The question naturally arises as to what effect individual alloying elements have upon the embrittling characteristics of the plain carbon steel. Current knowledge on the influence of alloying elements on temper brittleness may be found in the review papers of Hollomon" and Woodfine. Hollo-mon," from the results of other investigators, has shown that, in general, the amount of embrittlement increases with increasing alloy content (except for molybdenum and possibly tungsten and columbium). Jaffe and Buffum," by a comparison of the embrittlement in a plain carbon steel with that of a SAE 3140 steel postulated that the presence of alloying elements in moderate amounts tends to retard the development of temper brittleness. It is difficult to determine what effect chromium has upon temper brittleness, since most of the information available has been based on the combined effect of other elements with chromium, particularly nickel and manganese. However, Wilten, and recently Jolivet and Vidal,' Vida1, and Woodfine have reported that chromium steels are temper brittle, that the embrittlement is reversible with a maximum rate of embrittlement at approximately 975°F (525"C)," and that the susceptibility increases with increasing amounts of chromium. Taber, Thorlin, and Wallacel" have found a large embrittling effect with increasing chromium content in a medium C-Mn-Ni steel. But Hultgren and Chang," from their experiments conducted on synthetically prepared ternary Fe-C-Cr alloys, could not conclude that these alloys are susceptible to temper embrittlement. However, on addition of manganese or phosphorus, these Fe-C-Cr alloys became susceptible, from which fact they concluded that the embrittlement developed in chromium-bearing Fe-C alloys is due chiefly to the presence of these elements. Considerable data are available to show that molybdenum decreases the susceptibility of steel to temper embrittlement. However, its effectiveness in preventing or decreasing embrittlement appears limited to its presence in small amounts. Vidal" has shown that a plain 2 pct Mo steel was susceptible. Hultgren and Chang" also have shown that molybdenum additions in excess of 2 pct to synthetically prepared Ni-Cr steels did not prevent embrittlement. Jolivet and Vidal' and Lea and Arnold found that molybdenum reduced temper brittleness. Lea and Arnold further stated that molybdenum decreased the rate of embrittlement rather than the total amount of embrittlement, whereas Preece and Carter" have shown that the presence of molybdenum greatly reduces the equilibrium extent of the change at a given temperature but does not appear to influence the rate of embrittlement. There appears to be very little information as to how molybdenum by itself affects the temper brittleness susceptibility of a plain carbon steel.

Jan 1, 1956

By J. D. Miller, J. E. Sepulveda

Tar sand deposits in the state of Utah contain more than 25 billion bbl of in-place bitumen. Although 30 times smaller than the well-known Athabasca tar sands, Utah tar sands do represent a significant domestic energy resource comparable to the national crude oil reserves (31.3 billion bbl). Based upon a detailed analysis of the physical and chemical properties of both the bitumen and the sand, a hot-water separation process for Utah tar sands is currently being developed in our laboratories at the University of Utah. This process involves intense agitation of the tar sand in a hot caustic solution and subsequent separation of the bitumen by a modified froth flotation technique. Experimental results with an Asphalt Ridge, Utah, tar sand sample indicated that percent solids and caustic concentration were the two most important variables controlling the performance of the digestion stage. These variables were identified by means of an experimental factorial design, in which coefficients of separation greater than 0.90 were realized. Although preliminary in nature, the experimental evidence' gathered in this investigation seems to indicate that a hot-water separation process for Utah tar sands would allow for the efficient utilization of this important energy resource. The projected increase in the ever-widening gap between the domestic energy demand and the domestic energy supply for the next few years has motivated renewed interest in energy sources other than petroleum, such as tar sands, oil shale and coal. Although a number of research programs on the exploitation of national coal and oil shale resources have already been completed, very few programs have been initiated on the processing of tar sand resources in the United States. In recognition of their significance as a domestic energy resource, investigators at the University of Utah have designed an extensive research program on Utah tar sands. An important phase of this program, and the main subject of this publication, is the development of a hot-water process for the recovery of bitumen from Utah tar sands, as a preliminary step toward the production of synthetic fuels and petrochemicals. The term "tar sand" refers to a consolidated mixture of bitumen (tar) and sand. The sand in tar sand is mostly a-quartz as determined from X-ray diffraction patterns. Alternate names for "tar sands" are "oil sands" and "bituminous sands." The latter is technically correct and in that sense provides an adequate description. Tar sand deposits occur throughout the world, often in the same geographical areas as petroleum deposits. Significantly large tar sand deposits have been identified and mapped in Canada, Venezuela and, the United States. By far, the largest deposit is the Athabasca tar sands in the Province of Alberta, Canada. According to the Alberta Energy Resources Conservation Board (AERCB),2,3 proved reserves of crude in-place bitumen in the Athabasca region amount to almost 900 billion bbl. To date, this is the only tar sand deposit in the world being mined and processed for the recovery of petroleum products. Great Canadian Oil Sands, Ltd. (GCOS) produces 20 million bbl of synthetic crude oil per year. Another plant being constructed by Syncrude Canada, Ltd. is expected to produce in excess of 40 million bbl of synthetic crude oil per year. According to the Utah Geological and Mineral Survey (UGMS), tar sand deposits in the state of Utah contain more than 25 billion bbl of bitumen in place, which represent almost 95% of the total mapped resources in the United States.4 The extent of Utah tar sand reserves seems small compared to the enormous potential of Canadian tar sands. Nevertheless, Utah tar sand reserves do represent a significant energy resource comparable to the United States crude oil proved reserves of 31.3 billion bbl in 1976.5 Tar sands in Utah occur in 51 deposits along the eastern side of the state.4 However, only six out of these 51 deposits are worthy of any practical consideration (Fig. 1). As indicated in Table 1, Tar Sand Triangle is the largest deposit in the state and contains about half of the total mapped resources. Information regarding the grade or bitumen content of Utah deposits is still very limited. The bitumen content varies significantly from deposit to deposit, as well as within a given deposit. In any event, the information available6-8 seems to indicate that Utah deposits are not as rich in bitumen as the vast Canadian deposits which average 12 to 13% by weight.9 Although many occurrences of bitumen saturation up to 17% by weight have been detected in the northeastern part of the state (Asphalt Ridge and P. R. Spring), the average for reserves in Utah may well be less than 10% by weight. Separation Technology As in any other mining problem, there are two basic approaches to the recovery of bitumen from tar sands. In one

Jan 1, 1979

By P. M. Kelly, E. P. Butler

Four Mn-CLL alloys, containing 60, 70, 80, and 90 pct Mn, respectively, have been examined in the quenched and the quenched and aged conditions using electron microscopy and electron, neutron, and X-ray diffraction. After certain heat treatments the alloys transform from fee to fct and in the tetraom1 condition show a domain structure parallel to {101} planes. Neutron diffraction indicates that the domains are antiferrornagnetically ordered. The domain boundary contrast has been examined using bright- and dark-field microscopy, and the contrast effects observed under favorable conditions have been used to deduce the c axis orientation in each domain. The domains are extremely mobile and can be nucleated at precipitate particles and screw dislocations. The domain mobility is responsible for the high damping capacity. In the aged material a Mn precipitates in the Kurdjumov-Sachs orientation and results of both electron microscopy and neutron diffraction indicate that the matrix separates into two components—one rich in manganese and the other rich in copper. ALLOYS of manganese and copper have the unusual combination of a high damping capacity and good mechanical properties and have been the subject of a number of investigations as part of a general interest in high damping capacity alloys for engineering purposes.&apos;,&apos; SO far, however, there has been no reported electron metallographic study of these alloys. The Mn-Cu system has an extensive range of solid solubility at high temperatures, and the equilibrium phases are expected to be y (fee) and a Mn. The high damping capacity is associated with a metastable tetragonal structure of variable c/a ratio, which forms from the high-temperature y phase. This latter phase becomes more difficult to retain as the manganese content increases, and alloys containing more than 82 wt pct Mn undergo a reversible martensitic fcc — fct transformation on quenching. The X-ray work of Basinski and christian3 showed that the Ms temperature for the transformation was below room temperature for alloys in the range 70 to 82 pct Mn and increased linearly with manganese content. When quenched from the y region, alloys in the range 50 to 82 pct Mn are cubic at room temperature, but become tetragonal if aged at temperatures between 400" and 600°C. The martensite transformation occurs on cooling from the aging temperature. Tetragonal alloys have a banded microstructure and the bands analyze to be traces of (110) planes. Similar microstructures have been observed in In-Tl4 and in other manganese-base systems, such as Mn-Au5 and Mn-Ni.6 The mobility of the bands in Mn-Cu alloys can be demonstrated by optical examination of a polished specimen surface subjected to a cyclic stress.7 The bands appear and disappear as the stress is varied, and X-ray measurements of the (200,020) and (002) peak intensities confirm that a reversible reorientation of the tetragonal structure occurs. Meneghetti and sidhu8 investigated the magnetic structure of Mn-Cu alloys and found antiferromagnetic ordering in furnace-cooled alloys of composition >69 at. pct Mn. Magnetic super lattice reflections occurred at the (110) and (201) positions and the proposed structure was fct with the spins along the c axis. A more complete investigation by Bacon et al.9 confirmed this structure. The magnetic ordering temperature Tn was found to increase linearly with manganese content in the same way as the Ms temperature, and at any composition, Tn > Ms. This relationship suggested that the magnetic ordering was responsible for the cubic — tetragonal transformation in the manganese-rich alloys. The purpose of this investigation was to study the mechanism of high damping and the structural changes that occur on aging. The main technique used was transmission electron microscopy, but X-ray and neutron diffraction experiments were also carried out. EXPERIMENTAL Materials and Heat Treatment. The four alloys, provided by the Admiralty Materials Laboratory. were of nominai composition 60, 70, 80, and 90 Mn and all had low impurity levels, <0.05 pct C, <0.2 pct Fe. This material was cold-rolled to 200-µ strip with intermediate annealing and then given a final heat treatment of 24 hr in the range 800° to 900°C followed by water quenching. An identical heat treatment was given a length of 3/4-in.-diam bar of the 70/30 alloy from which the neutron diffraction specimens were machined. It was suspected that the tetragonal structures would be metastable at room temperature, and so the alloys were not aged until required for experiments. After aging in a salt bath the alloys were water-quenched. Thin Foil Preparation. Initial thinning to 50 to 75 µ was possible in a solution consisting of: 50 ml nitric acid 25 ml acetic acid 25 ml water The surface deposit and grain boundary etching was removed by a final electropolish at around 20 V in an electrolyte consisting of:

Jan 1, 1969

By D. T. Goettge, E. L. Robinson

The phase changes occurring in low carbon steel during hot strip mill rolling are shown to be metallurgically significant when related to commonly used temperature control points, particularly finishing and coiling temperatures. In combination, these temperatures are shown to have an important influence on the level and uniformity of hardness, grain size, and carbide characteristics of the finished hot and cold rolled sheets. PRODUCTION of wide flat-rolled products ordinarily requires a number of operations in sequence to prepare the material for shipment to the customer. Most products are tailor-made for specific end uses, with each operation contributing certain properties to the finished material. Since the characteristics imparted to the semifinished product by a given step in processing carry through to the finished product in varying degrees, it is important that the intermediate stages of production of flat-rolled strip be carried out with the same care which characterizes the last or finishing operations. The step of hot strip mill rolling is common to the production of all of the various types of flat-rolled product; therefore, the hot strip rolling is an especially important point at which to recognize and control those variables which have an effect on the surface characteristics and metallurgical properties of the finished product and which influence the ease of conducting subsequent operations. Orders entered at a producing mill usually show an end use or describe an article or part into which the ordered product is to be fabricated. Applying his experience as to the properties necessary in a finished sheet to suit the end use and to perform successfully in the fabrication involved, the metallurgist selects a steel of suitable composition and deoxidation practice, and slabs of appropriate dimensions are produced for rolling on the hot strip mill. At this stage of processing, the metallurgist faces the problem of controlling hot strip mill practice in the light of his diagnosis of the properties necessary to meet the end use, paying due attention to the accompanying problem of producing a strip which can meet processing requirements on subsequent units in the mill. It is the purpose of this paper to describe some of the factors which he must consider in solving these problems and to indicate some of the principles which guide him. Equipment, Physical Requirements of the Strip, and Temperature Measurement The metallurgist must, of course, be familiar with the physical layout of the mill, the temperature-measuring equipment available, and the physical requirements of the hot strip product before he can apply his metallurgical knowledge to the problem; hence, the first section will consist of a brief discussion of these matters. The usual hot strip mill consists of reheating furnaces, five or six roughing stands including a scale-breaker, holding table, and second scalebreaker, six-stand finishing mill, runout table with spray cooling facilities, and coilers. A schematic diagram of a typical layout is shown in Fig. 1. Slab temperatures are primarily a function of heating time and furnace temperatures, while mill speeds, spray practice, drafting practice, available water pressure, temperature of the cooling water, cross sectional dimensions of the strip, coil size, and equipment limitations, either singly or in combination, determine what rolling temperatures are practical on a given hot strip mill unit. Thus, it is possible that a set of temperatures which can be utilized successfully on one mill cannot be used on another. However, adjustments in temperatures and rolling practice can usually be made to develop the desired metallurgical properties. In addition to the metallurgical properties developed through proper temperature control, the hot strip mill must also provide strip with certain physical attributes which may be summarized as follows: Strip Cross Section—The strip contour should conform to a section which will give the best results in the cold reduction operation. This is generally recognized as a strip with 0.001 to 0.003 in. crown or shoulder-to-shoulder convexity depending on width, and freedom from concave, flat, or wedge-shaped cross sections which cause metal buildup in cold reduction. Excessive drop off in thickness at the edges can also be very detrimental in cold reducing to light gages. Gage, Width, and Camber—All of these must be controlled. For example, rundown or increasing thickness from the front to the back of the coil results in nonuniformity in the thickness of hot-rolled sheet product and in added difficulty with gage and welds in cold reduction. Similarly, excessive width variation is the cause of guide trouble and excessive edge scrap at later stages of processing, while excessive camber is the source of a variety of processing troubles. Type of Oxide—Product intended for pickling should have a predominance of the type of oxide most easily removable in sulfuric acid. It is generally recognized that this type is obtained by use of maximum table cooling water and cold coiling

Jan 1, 1957

By C. S. Giggins, F. S. Pettit

The oxidation of Ni-Cr alloys in 0.1 atm of oxygen has been studied at temperatures between 800" and 1200°C. For alloys with 30 wt pct or more Cr, continuous layers of Cr2O3 are formed during oxidation. In the case of alloys with chromium concentrations between approximately 5 to 30 wt pct, external scales of Cr203 are formed over grain boundaries whereas internal precipitates of Cr2O3 and external layers of NiO are formed at other areas on the alloy surface. When such conditions are present on the alloy surface, chromium diffuses laterally from those areas covered with a continuous layer of Cr2O3 to areas where a Cr2O3 sub scale exists and it is possible for the sub-scale zone to become separated from the alloy by a continuous layer of Cr2O3. Whether such a state will be attained depends upon the initial grain size of the alloy and the oxidation time. When the concentration of chromium in the alloy is less than 5 pct, Cr2O3 is formed internally both at grain boundaries and within the interior of grains and the alloy is covered with an external layer of NiO. MECHANISMS which describe the growth of oxide scales on nickel-base superalloys are complex and the effects produced by the various elements in these alloys on the oxidation behavior of superalloys are not clearly understood. In order to determine the influence of the different elements on the oxidation behavior of superalloys, it is first necessary to examine the oxidation properties of binary nickel-base systems which contain the principal elements present in the superalloys and then progressively more complex systems until compositions typical of the superalloys are attained. Chromium is present in virtually all nickel-base superalloys and the purpose of the present studies was to examine the selective oxidation of chromium in Ni-Cr alloys. The oxidation characteristics of Ni-Cr alloys have been extensively studied1-" to date principally as a result of the high oxidation resistance exhibited by some of these alloys. Ni-20Cr* has long been known *All compositions are given as wcight percent unless specified otherwise. to be oxidation resistant and is commonly used as resistance heating elements for service temperatures up to 1100°C. This alloy cannot be used for extended periods of time at higher temperatures because of the apparent reaction of the external scale with oxygen to form gaseous CrO3. In spite of the considerable work cited above some important aspects of Ni-Cr oxidation still remain unresolved. Virtually all of the previous studies agree that small additions of chromium to nickel, e.g., <10 wt pct Cr, result in increased oxidation rates as compared to that of pure nickel, whereas larger additions, e.g., 20 to 30 wt pct Cr, form alloys with substantially lower oxidation rates. The controversial aspects of the oxidation mechanisms for these alloys that still remain unresolved are as follows: 1) A description of the oxidation mechanism for the low chromium alloys. 2) A description of the oxidation mechanism for the high chromium alloys, particularly with respect to the composition of the external scale which results in the lower oxidation rates. 3) The specific alloy compositions at which the oxidation mechanism changes from that obtained for low chromium contents to that of the high chromium alloys and the reason for this transition. EXPERIMENTAL The Ni-Cr alloys listed in Table I were prepared from high purity metals by nonconsumably arc melting and casting as buttons. These alloys were then given a preliminary annealing treatment in argon at 815°C for 100 hr to promote homogeneity. Each button was cut into 0.250 in. thick sections that were subsequently cold-rolled to 0.050 in. thicknesses and annealed in argon at 815°C for 48 hr to provide a twinned, equi-axed grain structure. The grain size for these alloys was not uniform and the limits, within which the average grain size lies, are given in Table I for the single-phase alloys. All the alloys were single phase with the exception of the Ni4OCr alloy in agreement with the Ni-Cr phase diagram.&apos;&apos; Rectangular specimens were cut from the sheet to provide surface areas of approximately 2.5 sq cm. Exact areas were determined with a micrometer after surface preparation was completed. All of the specimens except the Ni-40Cr alloy and pure chromium were polished through 600-grit Sic abrasive paper, ultrasonically agitated in ethylene trichloride, rinsed with ethyl alcohol, and electro-polished. The specimens were electropolished in a 10 vol pct H2SO4 (conc), 6 vol pct lactic acid, methyl alcohol solution at 70" to 80°C for 2 min at a current density of 0.8 to 1.2 amp per sq cm. This electro-polishing procedure did not produce acceptable surfaces on the Ni-40Cr alloy nor on pure chromium and the oxidation properties of these materials were obtained for specimens polished through 600-grit Sic

Jan 1, 1970

By Eugene S. Machlin

A classical theory of solid solutions involving neavest-nergkbor intevactions with both central and linked-central forces between atoms has been developed. It has been found that the theory, where it can be checked quantitatively, is in ageement with experiment. The theory encompasses a description of many diverse pkenomena, such as antiphase shift structures, size effect, relative stabilities of various solutions, lattice para,neters, and order-disorder transitions. In particulav. a quantitative prediction not involving adjustable pavameters is made concevning the deviation of the Au-Cu interatonlic distance in long-range ordered (Ll,) Cu-Au I fronl the average distance based on the distances in pure gold and copper. This prediction, which is in agreement with expel-intent, has not been encompassed by any preuious theory. The theory of order-disorder is fragmentary. That is, no one theory exists that can explain the variety of qualitative phenomena observed. Further, many theories are not in good quantitative agreement with experiment. This subject has been reviewed by Muto and Takagi, Tuttman, and Oriani.3 There exists no doubt that the quasi-chemical approximation is not a complete description and that the inclusion of strain energy using macroscopic elasticity theory concepts leads to results in disagreement with experimenL4 The observation of antiphase domains and ordering systems such as Cu-Pt has led to Brillouin zone treat-ment of the order-disorder transition as opposed to the classical Ising model. The objective of this paper is to demonstrate that it is possible to develop a pairwise approximation model that can explain many of the observed order-disorder phenomena that have puzzled investigators in the past. This theory is based upon an empirical model due to ergmman&apos; for the elastic constants of metals. This model is generalized for multicomponent systems. As will be shown, the theory yields a short-range ordering energy for the disordered solution which differs from the ordering energy calculated from the differences in energy of disordered and long-range ordered solutions. It will be demonstrated that there is no necessary correlation between heats of formation and the tendency to order or between size effect and the tendency to order. Also, the existence of antiphase domains and iso-short-range-order systems that form superlattices (Cu-Pt) is predicted on the basis of the theory. Further, the relative stability of competing superlattices is calculable from the theory. If single-crystal elastic-moduli data are available for the pure components and one superlattice then there exists but one adjustable parameter in the calculation of lattice parameters for both the disordered and ordered solid solutions. In one special case, no adjustable parameters are required and a quantitative prediction is made. For the calculation of energies and partial order, there exists but one additional adjustable parameter, the pair-exchange energy V used in the quasi-chemical approximation (or the Ising model.) However, in these calculations, much more precise values are required for the single-crystal elastic moduli than available if the quantitative uncertainties in the predicted values of the energies are to be sufficiently small. THEORY ~er~man&apos; has developed a model with which he was able to obtain fair agreement with experiment for the relations between the elastic constants for metals. This model which we shall call Bergman&apos;s model is a linear combination of his models I and 11. In effect, Bergman, in this model, considers that each interatomic distortion is composed of two components: a classical central force distortion with an associated central force constant and what we shall call a linked-central force distortion with an associated linked central-force constant. The linked-central force distortion component obeys the constraint that the sum of such distortions over all the bonds equals zero. No constraint is imposed on the classical central force distortion component. Bergman&apos; derives the constraint on the linked-central force distortion on the basis of application of Pauling&apos;s relation between bond distortion and bond number to metals.ga This assumption is not logically necessary, however, and the Bergman model may be taken as a mathematical model for elastic constants, e.g., a purely empirical model without a physical basis. In the present work, the method of Bergman has been applied to two-component systems (solid solutions). In place of an external strain—which would allow a calculation of the elastic constants for the two-component system—it is considered that internal interatomic distortions exist as a consequence of having three potentially unequal distortion-free interatomic distances and but one "average" interatomic distance. It is assumed that the distortion-free interatomic distances between atoms of the same element are those found in the pure element having the same undistorted crystal structure as the solid solution. The distortion-free interatomic distance between unlike atoms is in general not measurable except in the probably nonexistent

Jan 1, 1967

A calibration of the Frantz isodynamic magnetic separator relates the current in the electromagnet and the transverse slope of the chute to the specific susceptibility at which separation is being effected. In addition to providing a rapid measure of the specific susceptibility of minerals and composite grains, the calibration provides a guide to operation of the Frantz isodynamic separator for mineral separation, and to the nature of composite grains. The most sensitive magnetic separation is obtained by using the highest practicable current and transverse slope, and the lowest practicable longitudinal slope.INTRODUCTIONProbably all minerals have a magnetic susceptibility measurable by sensitive techniques. However, magnetic separation between minerals which have low but different specific susceptibilities, and which for many purposes are considered essentially non-magnetic, has only become practicable with the development of the Frantz isodynamic separator.The construction and principle of this highly sensitive separator has been described by S. G. Frantz as an appendix to a description by &apos;Gaudin and Spedden (1943) of its application to the separation of sulphides constituting mill products.Various publications (e.g. Hess, 1949, Hutton, 1952) refer to its use in separating the mineral constituents of rocks, ores, and unconsolidated sands.As used in this study the Frantz isodynamic separator consists of a vibrating chute mounted centrally between the pole pieces of an electromagnet (Fig. 1). This unit may be inclined in any direction by the universal mounting, and the electromagnet current is continuously adjustable from 0 to 1&#183;5 amperes.

Jan 1, 1957

By C. F. Ellis

Interest in the composition of automobile exhaust gases has been greatly accelerated by recent studies that have indicated that these gases contribute to smog formation in urban areas. Several analytical methods have been published for the analysis of the hydrocarbons in the exhaust gases, but similar information has not been published on the analysis of the oxygenated compounds. This report presents procedures for chemical analyses of certain oxygenated compounds, procedures for sampling the exhaust gases for these analyses, and the analytical results of some experimental work in which these procedures were used. A stationary 1956 model, 170 horsepower, V-8, dynamometer-mounted engine, operated to simulate vehicular cruise conditions, was used for the study. Exhaust gases produced from a regular-grade Midcontinent gasoline during 15-40- , and 60-mile-per-hour cruise operations were analyzed by chemical methods for formaldehyde, total aldehydes, total carbonyls, total alcohols, carbon associated with the oxygenated compounds, and the oxides of nitrogen. Experimental results indicate: 1. The range of concentrations of the total of the functional groups of oxygenated compounds, expressed as moles per million moles of water-free exhaust gases, is about 90-100; the concentrations increase with engine speed. 2. Aldehydes comprise about 80-90 mole- percent of the total of the oxygenated groups, and formaldehyde accounts for about 50 mole-percent of the total aldehydes. 3. The order of magnitude of the concentration of the carbonyls is the same as that for the aldehydes, indicating that ketones are not a major group. A colorimetric test indicates that methyl ketones Constitute about 2 mole-percent of the oxygenated compounds.

Jan 1, 1961

The U.S. economy continued to expand during 1986, although at the slowest pace since 1982 when the current business expansion began. Real gross national product (GNP) grew about 2.6%, slightly below 1985 growth but significantly below that of 1984. The continuing expansion was fueled predominantly by rising personal consumption expenditures as well as growing government purchases of goods and services. Fixed investment, although positive, failed to grow at the rate of the overall economy. The trade deficit, which increased more than 30% over the record 1985 level, continued to depress economic growth despite significant declines in the price of petroleum and the value of the dollar. On a more positive note, domestic employment continued to grow--increasing over 2 million from the 1985 level--and inflation and interest rates declined. The unemployment rate dropped slightly to 7.0%. Of direct importance to the minerals-producing industries, expenditures for consumer and producer durables as well as investment in residential construction each grew faster than the overall economy, but investment in nonresidential structures fell more than 10% from the 1985 level. The low level of real capital spending declined about 3% in 1986, following strong gains in 1985 and in 1984. Capital expenditures for new plant and equipment bottomed out in the second quarter and expanded in the third and fourth quarters. The decline in the first half of the year was partially attributed to uncertainty arising from the then-proposed tax bill. Also, there was little impetus for corporations to invest in new capacity since factories were operating at less than 80% of capacity during most of the year, in part because of continued competition from imports.

Jan 1, 1987

The U.S. economy continued to expand during 1986, although at the slowest pace since 1982 when the current business expansion began. Real gross national product (GNP) grew about 2.6%, slightly below 1985 growth but significantly below that of 1984. The continuing expansion was fueled pre-dominantly by rising personal consumption expenditures as well as growing Government purchases of goods and services. Fixed investment, although positive, failed to grow at the rate of the overall economy. The trade deficit, which increased more than 30% over the record 1985 level, continued to depress economic growth despite significant declines in the price of petroleum and the value of the dollar. On a more positive note, domestic employment continued to grow--increasing over 2 million from the 1985 level--and inflation and interest rates declined. The unemployment rate dropped slightly to 7.0%. Of direct importance to the mineral-producing industries, expenditures for consumer and producer durables as well as investment in residential construction each grew faster than the overall economy, but investment in nonresidential structures fell more than 10% from the 1985 level. The level of real capital spending declined about 3% in 1986, following strong gains in 1985 and in 1984. Capital expenditures for new plant and equipment bottomed out in the second quarter and expanded in the third and fourth quarters. The decline in the first half of the year was partially attributed to uncertainty arising from the then-proposed tax bill. Also, there was little impetus for corporations to invest in new capacity since factories were operating at less than 80% of capacity during most of the year, in part because of continued competition from imports.

Jan 1, 1986

By A Fitzpatrick

An integrated and dynamic water balance and water quality model for a proposed base metal mine has been created using the GoldSimÖ software package. The model was initially developed to address the water quality of the Tailings Management Facility pond and how pond discharge would potentially affect downstream water quality. However, due to changing project objectives, the modelling has expanded over time through an iterative process to encompass the full mine site and all receiving environment water quality. The model results have been used as part of the mine planning/design process but are also intended to be used for the environmental impact assessment and permitting process. The model has been built to allow for changes to input terms to be readily made as additional data becomes available. The model results serve as a prediction of water quality and the potential project impacts to the receiving environment water quality. This paper presents the process by which data from a wide variety of sources was used to construct the model and how the model results were used to guide development of a conceptual water management plan. In addition, the role of the model in establishing proposed discharge criteria will be highlighted. The model has been delivered as a fully functional application with the intent that the end-user can be other than the modeller, the development of this aspect shall also be covered. CITATION:Fitzpatrick, A, 2012. Predictive site-wide water quality modelling as a tool for mine planning and permitting - A recent Canadian example, in Proceedings Life-of-Mine 2012 , pp 71-78 (The Australasian Institute of Mining and Metallurgy: Melbourne).

Jul 10, 2012

By F. J. N. Bruwer, T. R. Stacey

The research described considers whether the variability in coal material strength, as derived through a series of uniaxial compressive strength (UCS) tests, could be used to indicate the variability in coal pillar strength. The aim is to be able to use a distribution of UCS tests as input into the coal pillar strength calculation. This will allow the pillar design to be expressed in terms of a probability of failure rather than as the commonly used safety factor. To achieve this, the bulk strength factor associated with commonly used pillar strength formulae was replaced with a distribution of UCS results divided by an adjustment factor. The factor was determined so as to ensure that the resulting bulk strength does not deviate from the statistically determined bulk strength published in the original formulae. This approach enabled pillar strength distributions to be obtained using industry-accepted strength formulae, subsequently allowing for a probability of failure to be calculated for a specific pillar design. Using a regional coal material strength curve as a baseline, coalfields in which the coal is stronger than the regional mean can be identified and the pillar designs optimized. This is based on the stronger coals achieving lower probabilities of failure at similar safety factors. The research has considered actual UCS data from multiple mines in the Mpumalanga coalfields of South Africa, and has proved that the variability in material strength between coalfields could allow for some optimization using the proposed approach. Based on the data used in the study, a 2.78% increase in extraction could be achieved. However, further research will be required to validate the results of the study in an underground environment.

May 3, 2023

By H. Simpson

Middelburg mine is an opencast export mine operating in the Eastern Transvaal Coalfields of the Republic of South Africa (Figure 1). The mine was commissioned in 1983, to its original design capacity of 4.25 million sales tonnes which was subsequently raised to 5.5 million sales tonnes in 1989. This increase in export tonnage was brought about through an increase in the mines export allocation and the revision of management strategies. The maintenance of this increased tonnage with minimal additional capital has been achieved through improved operational efficiencies and the upgrading of planning operations through the purchase of an Enhanced Mine Scheduling Software Package (EMSSP/MINESTAR). Simultaneously with this requirement were additional factors that contributed to the implementation of the EMSSP. The mine was owned by a joint venture of two partners and is operated by Middelburg Mine services (Rand Mines Pty. Ltd) (Figure 2). This structure of ownership required the evaluation of numerous planning options to be under-taken concurrently and hence the need for an efficient streamlined scheduling pack-age that could fulfill this requirement. Compounding the EMSSP needs further, were the problems experienced in the export market, these being the external pressure from current and new suppliers in the market place which obliged management to provide solutions to rapidly changing market conditions. These fluctuating market conditions have had a direct effect upon the mine&apos;s planning strategies because of the large percentage of the mine?s sales sold on the international spot market. The sanction factor for all South African exporters has had a negative effect upon price and this was compounded further through internal competition in 1987 via suppliers attempting to maintain the market position.

Jan 1, 1991