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By G. Sandoz, D. A. Meyn

The fracture surfaces of specimens of titanium a1loys which exhibited susceptibility to subcritical crack growth in a wide variety of environments, including aqueous solutions, alcohols, hydrocarbon gases, carbm tetrachloride, and dry air, were examined. The dominant and characteristic fracture mode was cleavage. which became mixed with an increasing porportion of ductile fracture modes (dimples. and so forth) as the applied KI was increased. The cleavage plane in all the alloys was oriented at 15 deg from (0001) of the a phase. No indices could be assigned to the 15 deg cleavage Plane because of uncertainty as to which zone it belonged to. SLOW crack growth and delayed fracture take place in certain titanium alloys at stress intensity (KI) lower than the critical values for complete fracture (KIc) found in short-time tests for plane strain fracture toughness.' Such crack growth may take place in very dry air on sustained loading if the alloy contains a sufficient amount of hydrogen and if the stress intensity factor is higher than a threshold value designated KIH.2 Similarly, a number of environments such as water, alcohols, and hydrocarbons may cause slow crack growth provided the stress intensity factor exceeds the threshold value KISCC. In this paper it is intended to examine some fracto-graphic and crystallographic details of this subcritical crack propagation process as related to alloy composition, environment, and stress intensity. It has been previously reported that subcritical crack growth takes place in titanium alloys in aqueous environments by cleavage.2"6 The observations to be described here indicate that in the alloys investigated the mode of subcritical crack propagation is more dependent on the stress intensity during cracking than on the environment. The fracture mode is illustrated with photographs of the fracture surfaces at a wide range of magnifications. The crystallographic character of the cleavage mode is also presented and the results discussed. EXPERIMENTAL APPROACH The studies were made on four titanium alloys— Ti-8A1-1Mo-lV, Ti-5A1-2.5Sn, Ti-7Al-1Mo-lV, and Ti-0.350. Of these four alloys the first two were mill annealed and fine grained. The latter two were noncommercial and coarse grained, having been prepared for other experimental purposes. The crystallographic studies required a large-grain size. Consequently the fine-grained alloys were an- nealed in a vacuum at 1950°F and furnace cooled to produce samples for these studies. This treatment did not alter the microstructure significantly except for grain size. All the aluminum-containing alloys showed primary a with a network of ß-phase particles within the a grains; Figs. 1 and 2 show this microstructure for the specific case of fine-grained Ti-8A1-1Mo-1V. The cracks in the figures will be discussed later. Fractographic studies were made on all the alloys, but the electron fractographs displayed herein are from the Ti-8A1-1Mo-1V alloy in the fine-grained condition. Emphasis was placed on this commercial alloy because considerable information was available on the stress-corrosion cracking characteristics in a variety of environments. Sufficient supplementary studies of the alloy after the anneal to produce large grains were conducted to establish that the stress-corrosion cracking susceptibility and fractographic characteristics were not changed significantly. Thus it could be inferred that the crystallographic data obtained exclusively on large-grained specimens would apply also to fine-grained specimens. Fracture surfaces for study were obtained using precracked cantilever beam specimens loaded as described by Brown. The precracked section of a specimen was immersed in the environment of interest. Values of initial stress intensity factor Kli were chosen slightly higher than KISCC in order to obtain a fairly extensive region of slow crack growth. Fracture surfaces representative of a variety of stress intensity levels could then be examined because in the cantilever beam test KI increases as the crack propagates.

Jan 1, 1970

By L. H. Schwartz, S. K. Lahiri, M. E. Fine, D. Chandra

WHILE the yield stress of solution treated Fe-Cu alloys increases rapidly with aging, a precipitate has only been directly observed in overaged samples.'-" This precipitate is essentially pure fcc copper, the E phase. Hornbogen2 suggested that the initial hardening is due to clustering of copper atoms in a bcc iron matrix. Hornbogen was not able to obtain direct confirmation for such bcc clusters by either X-ray diffraction or electron microscopy; however, he did observe analogous clustering in Fe-Au alloys7 where the scattering factor difference between iron and gold is more favorable than that for iron and copper. Fujii, Nemoto, Suto, and Monma4 recently observed diffuse electron diffraction scattering in a 2.5 pct Cu alloy aged at 500°C: for 30 hr. They suggested that part of the diffuse SCattering was due to a modulated structure of 26 to 42 wavelength parallel to the (110) matrix plane. This disappeared with further aging. However, 30 hr at 500°C is well beyond the peak strengthening in this alloy which occurred at less than I hr at 500oc.4 By transmission electron microscopy they could not distinguish between quenched samples and samples aged to maximum strength. The purpose of the present investigation was to try to obtain further confirming evidence for a precipitation stage preceding the formation of the fcc E phase. Change in Young's modulus was selected for study be- EXPERIMENTAL PROCEDURE The Fe-1.67 at. pct Cu alloy was supplied by the U.S. Steel Corporation in the form of a hot forged rectangular bar. The carbon and nitrogen contents of this alloy were 70 ppm and less than 10 ppm, respectively. The samples for Young's modulus measurements were made by cold swaging to & in. diam and cutting to 2 in. lengths. Tensile test samples 1 in. long were prepared by cold rolling to a thickness of 0.012 in. and then machining to gage dimensions of 0.5 by 0.012 in. For Mijssbauer studies, foils 1 mil thick were prepared by cold rolling. During fabrication the alloy was given intermediate anneals at 845OC followed by quenching. The specimens after preparation were solution treated at 1000°C (y field) or 845°C (a field) for 1 to 2 and 5 to 6 hr, respectively, in a dynamic vacuum of about 2 x lom6 torr. The tensile and modulus specimens were drop-quenched into iced brine through an aluminum foil which vacuum-sealed the bottom of the furnace tube. The MOssbauer foils were drop-quenched into helium gas cooled by passing through N2 (liq) in a copper coil. The grain diameter of the modulus and tensile specimens was between 0.025 and 0.035 mm. The samples were aged from 400" to 500°C in a dynamic vacuum of 2 x 10-6 torr and cooled by pulling into a water-cooled portion of the vacuum system. The dynamic Young's modulus was measured at room temperature using electrostatic excitation and detection of longitudinal vibrations.' Neglecting any length change during the aging, (this is expected to be comparatively small), the fractional change of Young's modulus (?Et)/Eav, is equal to twice the fractional change of the resonant frequency (?ft)/fav . ?ft =ft —fo where fo and ft are the resonant frequencies of a specimen before aging and after aging to time t. The measured f's, approximately 51,000 cps, were corrected to 25°C for any variation in the ambient temperature. The tensile tests reported were carried out at room temperature in an Instron with a 1000 lb full scale load cell. The flow stress reported herein corresponds to a strain of 0.002. The Massbauer spectra were taken in a constant acceleration model AM-1 NSEC spectrometer. The

Jan 1, 1970

By E. A. Hopkin

Laboratory tests conducted by the author. together with actual field experience in Canada. have indicated the magnitudc of some of the factors affecting ability of drilling mud to clean the hole. A correlation was observed between funnel viscosity and particle slip velocity. A relationship was also observed between the Bingham yield value of the mud and the particle .slip velocity. Laburatury tests indicated little change in slip velocity with funnel viscosities up to ahout 80 sec/qt or to yield values of aboul 15 lb/100 .sq ft. The lip velocity of the partic1e.s. studied vuried from 160 ft/min with water to neerly zero for high mud viscositie. lncreasing the mud density, creating laminar annular mud flow or rotating the drill pipe Mlay also improve the carrying capacity of mud. Field tests indicated that during fast upper-hole drilling, the rate of cutting removal must he sufficien to maintain the concentration of cuttings in the annulus at less than 5 percent by volume to prevernt balled-up drill collars and .stuck drill pipe. INTRODUCTION In the Mid-Continent area, hole cleaning during rotary drilling normally involves the use of low-density drilling muds with funnel viscosities varying from 26 to 80 sec/qt. However, frequently it is necessary to raise the mud viscosity in excess of 200 sec/qt to clean the hole. In addition, field experience has indicated that during very fast upper-hole drilling, annular mud velocities considerably in excess of the maximum slip velocity of the cuttings (normally considered to be 120 to 150 ft/min) are necessary to prevent balling the bit and drill collars. In view of these variations in mud properties and circulation rates necessary to clean the wellbore, a study of factors affecting hole cleaning and the relative magnitude of each was undertaken. The study was also prompted by the knowledge that a drilling mud with properties close to those of water would drill at rates several times that of a viscous mud. The carrying capacity of these near-water muds was not clearly understood, however, and many drillers were reluctant to use them. DISCUSSION EFFECT OF MUD PROPERTIES ON DRILLING RATE The effect of mud properties on drilling rate has been studied by a number of investigators. 1t is clear from their work that mud properties close to those of water will significantly improve drilling rate and bit life. If the water-type mud will also effectively clean the hole and cause no logging or formation impairment problems, the choice of an appropriate drilling fluid is clear—use water. Often, however, a water-type drilling fluid is not capable of effectively cleaning the hole. THEORY OF PARTICLE SETTLING The forces acting on a particle that is allowed to settle through a fluid quickly reach equilibrium, and the particle settles at some constant velocity referred to as the slip velocity. The slip velocity of any particle depends upon a number of factors such as density and viscosity of the fluid, the volume, specific gravity, shape and roughness of the particle, and upon the shape and area of the projected face of the particle at right angles to the direction of relative fluid solid movement. Brown' expressed this relationship for spheres in Newtonian fluids in the following equation. where v, = slip velocity D, — diameter of sphere p, = density of spheye p, = density of liquid /a = friction factor g = constant. The friction factor f, varies as the size, shape, surface roughness and slip velocity of the body, and as the density and viscosity of the fluid varies. WaddelV ound experimentally that where dynamic similarity exists. f, was a function of the particle Reynolds number (NRe) of the sphere (Fig. 1). Fig. 1 shows that f, for spheres varies from over 10 for low NKe numbers to 0.44 for N,,. over 600.

By KING HAMILTON GRAYSON

IRON and steel were the only basic industries in the Soviet Republic in 1928 that lagged behind the pre-war production on a comparative basis. This was due to the almost complete obliteration of all industrial units manufacturing iron and steel products during the Revolution, which resulted in the production of 1921 being only 4 per cent of that of 1913. By the end of 1927 the production of pig iron had reached a point equal to 70.5, per cent of 1913; steel was 84 per cent ; and rolled iron was 78 per cent. According to the. planned program of the Supreme Economic Council, the pre-war level of the iron and steel industry was to have been reached and exceeded during 1929-30. The production of the entire period from 1917 to 1928 in comparison with 1913, is shown in the following table

Jan 1, 1929

By E. P. Pfleider, Rolland L. Blake

IT is generally believed that the amount of diamond drilling will increase appreciably in the next decade, as the seaarch for minerals throughout the world becomes more difficult and intense. An attendant problem may be one of short diamond supply, resulting in higher bit and drilling cost. With this background, the U. S. Bureau of Mines' and the School of Mines at the University of Minnesota' have established comprehensive research programs in diamond drilling. One of the several aims is the design of a more efficient bit, which would lower diamond consumption and increase rate of advance, both essential in reducing drilling costs. The objective of the specific research problem" discussed in this paper was an investigation of the cutting action of the cliamonds set in a diamond drill bit, cutting action meaning the manner in which the diamonds cut or. loosen the minerals in the rocks being drilled. In the literature on cutting action such descriptive terms are used .as: grinding, wearing, cutting, breaking, shearing, scraping, melting, and chipping. These actions were seldom described or defined. Grodzinski describes the cutting action of a single diamond in the shaping of certain types of material as "breaking out chips of the material." Brittle mate-. rials break as small separate chips, and tough materials, because of heat generated, give a continuous chip. Deeby said about diamond drills: "When diamonds are forced into the formation and rotated, they either break the bond holding the rock particles together, or they cause conchoidal fracture of the rock itself. The former action occurs when drilling in sandstones, siltstones, shales, etc. and the latter action when drilling in chert, flint, or quartz." He said that diamonds cut on the "grinding principle" but he does not define or elaborate on this action. The cutting action of diamonds on glass was first investigated about 1816 by Dr. W. H. Wol-laston, an English physicist. The best glass-cutting diamonds have a natural or artificially rounded cutting edge. This edge first indents the glass and then slightly separates the particles, forming a shallow and nearly invisible fissure. Since none of the material is removed, this action is one of splitting rather than cutting. No other reports of research work on the cutting action of the diamond were found, and further work was considered justified and advisable. It is impractical, even if possible, to observe directly the cutting action of a diamond drill bit in rock; therefore it was necessary to devise an indirect method. It was believed that a study of the following three observations would lead to a better understanding of the cutting action: 1—the appearance of the minerals or rock surface in the bottom of the hole, 2—the size, shape, and other characteristics of the drill cuttings, and 3—the condition of the diamonds in the bit. The cutting action in a particular rock probably varies with bit pressure and speed. If the bit were slowly lifted off the rock, the effect of decreasing pressure might obliterate those bottom hole characteristics that are specific at the test pressure. Likewise, if the drill were stopped with the bit still in contact with the bottom of the hole, then decreasing speed effects would tend to obliterate the characteristics at the set test conditions. Therefore, in order to preserve those cutting effects impressed on the rock at test conditions, it seemed necessary to lift the bit off the bottom of the hole almost instantaneously once drilling conditions, i.e., revolutions per minute, pressure, and water flow became constant. In addition to observing the cuttings, the bit, and the bottom of hole, it seemed desirable to collect some quantitative data for purposes of correlation with the observations and for a record of bit performance. Consequently such data as revolutions per minute, force applied, and rate of advance of the bit were recorded. Six rock types, listed in Table I, were chosen for the tests. It was felt that these rocks had most of the variable characteristics of texture, bonding, and mineral hardness met in the common rocks generally being drilled. The sandstone was so poorly cemented as to be friable, even though most of the cement was silica. The limestone, though well cemented, was quite porous. Originally it was planned to conduct the tesk work with a full-scale drill unit, using EX bits, 7/8-in. core, 11/4-in. OD. The drill worked well, but was too cumbersome for rapid, accurate drilling of many short holes (1 ½-in.) in varied rock types. A new

Jan 1, 1954

By R. E. Schilson, F. H. Poettmann

The direct determination of the stabilized performance behavior of low capacity, slowly stabilizing gas wells is extremely time-consuming and wasteful of gas. From both field experience and theoretical considerations, a test procedure has been evolved by which the stabilized hack pressure behavior of such gas wells can he predicted without having to revert to long time flow tests. The method consists of using the isochrona1 test procedure to establish the slope of the back pressure curve, "n", and the short time variation of the performance coeficient, "C", with time. From this short time transient flow data and theoretical considerations, the value of C at large times can he established. By assuming the radius of drainage of a well to be half the distance between wells, one can calculate the stabilization time for various well spacing patterns. Once the stabilization time for a given .spacing has been determined, the value of C can be calculated and the stabilized back-pressure curve can he. establbrhed. The calculated perfortnance coefficient as a function of tinze was cotnpared to the experintentally measured values for a number of gas wel1.e. The deviation of the calc~*lated from the cxperimental res1tlt.s vary depending on the set of short time experimental points used to evabrate the parcrttzeters of the equation. The longer the tirne far the flow test data user1 in the calculations, the better was the agreement with the experimental results. The time necessary to obtain this data from well tests varies consitlerably, depending on the physical natrcre oj the rt3scrvoir under consideratiot~. INTRODUCTION For many years, the U. S. Bureau of Mines Monograph 7' has served as a guide for testing and evaluating the performance of gas wells by means of the back-pressure method. The back-pressure performance of a gas well is expressed by the following equation: Q~ CW-W.........(I) where the characteristics of the back-pressure equation are determined by C, the performance coefficient, and n, the exponent which corresponds to the slope of the straight line when Q and {P* - PN2) are plotted on logarithmic paper. Q is gas flow rate at standard conditions, and P, and P, are equalized and flowing bottom-hole pressures, respectively. Prior to the development of the back-pressure tcst, the "open flow" capacity method of testing a well was common. By this method, the wells were flowed wide open to the air and the flow rate measured. Such procedure was wasteful of gas and did not provide information on the deliverability of the gas to the pipe line. Monograph 7 Procedure The back-pressure method of testing wells was developed to overcome these shortcomings. Although much has been learned regarding the laws of the flow of gas through porous formations, the original development of the back-pressure relationship was based entirely on empirical methods. The back-pressure behavior provides the engineer with information essential in predicting the future development of a field. It permits him to calculate the deliverability of gas into a pipe line at predetermined line pressures, to design and analyze gas gathering lines, to determine the spacing and number of wells to he drilled during the development of a field to meet gas purchasers' requirements, and to solve many other technical and economic problems. As described in Monograph 7, the flow-after-flow method of back-pressure testing, when applied to fast stabilizing and usually high capacity wells, correctly characterized the behavior of the well. However, as the value of the gas at the wellhead increased, small capacity gas wells having slow rates of stabilization became economically operable. The flow-after-flow method of testing could not be used to describe the behavior of these slowly stabilizing wells. The procedure of Rawlins and Shellhardtl for establishing the back-pressure behavior of a gas well was based on the rcquirement that the data be obtained under stabilized flow conditions; that is, that C is constant and does not vary with time. C depends on the physical properties of the reservoir, the location, extent and geometry of the drainage radius, and the properties of the flowing fluid. In a highly permeable formation, only a very short period of time is required for the well to reach a stabilized condition, and, consequently, the requirements for the test procedure described in Monograph 7 are met. For a given well, n is also constant

By F. C. Bond

SIZE of grinding media is one of the principal factors affecting efficiency and capacity of tumbling-type grinding mills. It is best determined for any particular installation by lengthy plant tests with carefully kept records. However, a method of calculating the proper sizes, based on correct theoretical principles and tested by experience, can be very helpful, both for new installations and for guiding existing operations. As a general principle, the proper size of the make-up grinding balls added to an operating mill is the size that will just break the largest feed particles. If the balls are too large the number of breaking contacts will be reduced and grinding capacity will suffer. Moreover, the amount of extreme fines produced by each contact will be increased, and size distribution of the ground product may be adversely affected. If the balls added are too small, grinding efficiency is decreased by wasted contacts that are too weak to break the particles nipped; these largest particles are gradually worn down in the mill by the progressive breakage of corners and edges. Ball rationing is the regular addition of make-up balls of more than one size. The largest balls added are aimed at the largest and hardest particles. However, the contacts are governed entirely by chance, and the probability of inefficient contacts of large balls with small particles, and of small balls with large particles, is as great as the desired contact of large balls with large particles. Ball rationing should be considered an adjunct or secondary modification of the principle of selecting the make-up ball size to break the largest particle present. Empirical Equation In 19521,2 the author presented the following emerical equation for the make-up ball size: B - ball, rod, or pebble diameter in inches. F = size in microns 80 pct of new feed passes. Wi - work index at the feed size F. Cs - percentage of mill critical speed. S — specific gravity of material being ground. D == mill diameter in feet inside liners. K - 200 for balls, 300 for rods, 100 for silica pebbles. Eq. 1 was derived by selecting the factors that apparently should influence make-up ball size selection and by considering plant experience with each factor. Even though Eq. 1 is completely empirical, it has been generally successful in selecting the proper size of make-up balls for specific operations. But an equation based on theoretical considerations should be used with more confidence and have wider application. The theoretical influence of each of the governing factors listed under Eq. 1 was accordingly considered in detail, as described below, and a theoretical equation for make-up ball sizes was derived. Derivation of Theoretical Equation Ball Size and Feed Size: The basis of this analysis is that the largest ball in a mill should be just sufficient to break the largest feed particle into several pieces, excluding occasional pieces of tramp oversize. In this article the size F which 80 pct passes is considered the criterion of the effective maximum particle feed size. The smallest dimension of the largest particles present controls their breaking strength. This dimension is approximately equal to F. As a starting point for the analysis it is assumed that a 1-in. steel ball will effectively grind material with 80 pct passing 1 mm, or with F- 1000µ or about 16 mesh. The breaking force exerted by a ball varies with its weight, or as the cube of its diameter R. The force in pounds per square inch required to break a particle varies as its cross-sectional area, or as its diameter squared. It follows that when a 1-in. ball breaks a 1-mm particle, a 2-in. ball will break a 4-mm particle, and a 3-in. ball a 9-mm particle. This is in accordance with practical experience, as well as being theoretically correct. Confirmation of this reasoning is supplied by the Third Theory of Comminution," which states that the work necessary to break a particle of diameter F varies as F. Since work equals force times distance, and the distance of deformation before breaka4e varies as F it follolvs that the breaking force should vary as F½ These relationships are expressed in Table I, with a 1-in. ball representing one unit of force and breaking a 1-mm particle. This establishes theoretically the general rule used in Eq. 1 that the ball size should vary as the square root of the particle size to be broken. Ball Size and Work Index: The work input W required per ton" varies as the work index Wi, and the

Jan 1, 1959

By W. F. Kieschnick, Eugene Harrison, W. J. McGuire

W. J. McGuire, et al, are to be commended for their undertaking of a mathematical solution of a very difficult problem. Unfortunately, however, a mathematical approach requires the application of several assumptions. These assumptions appear to be unrealistic and lead to answers which do not describe what actually happens when hydraulically fracturing oil and gas wells. Considering laboratory confirmation of breakdown phenomena, the authors appear to have tested their theories only on cement specimens and on samples of Austin limestone, much too small to provide any fracture system. This work resulted in the formation of vertical fractures. If the authors had tried similar experiments on thick walled cylinders made from almost any sandstone cores, they would have found that, using crude oil as the breakdown fluid, horizontal fractures would almost always occur, and at pressures much lower than any calculated. They would also find that by confining the fluid to within the bore (using oil base mud for example) on similar samples, the pressures required to burst the cylinders would be considerably higher and most of the fractures would be vertical. This breakdown pressure behavior has been duplicated in wells in Texas, Oklahoma, Kansas and in Wyoming. Considering field data the phenomena of different breakdown pressures for different breakdown techniques can be further illustrated. Most production and service personnel will agree that a breakdown can be more easily obtained if injection into a formation can be established prior to the occurrence of the breakdown. This is true whether the formation being treated is completed as open hole or as a perforated interval. This is clearly illustrated by a Lakota well in Wyoming, completed open hole at a total depth of 7,358 ft. An attempt to vertically fracture this well failed when a bottom hole pressure of 10,326 psi was insufficient to break down the formation. A non-penetrating type fluid (oil base mud) was in the well at the time the breakdown was tried." The oil base mud was then cleaned out of the well and replaced by a 30" API gravity crude oil. With this oil in the hole the formation breakdown was easily accomplished at a bottom hole pressure of 3,607 psi. This large difference in fracture pressures would be impossible according to the theories presented by McGuire, et al. The authors have used as an example the breakdown pressures experienced when acidizing Permian Basin wells. During acid treatments of limestone and dolomite the "breakdown" (drop-off in pressure) seldom occurs until some injection of acid has been accomplished. In these cases the breakdown is most likely to result from the chemical reaction of acid and rock in already existing vugs and fractures rather than from making a new fracture by hydraulic pressure. If this is true, then results in the Permian basin should not be used to validate the authors' calculations. *** AUTHORS' REPLY to ROSCOE C. CLARK and HENRY F. COFFER The purpose of our laboratory experiments in which thick-walled rock cylinders were hydraulically fractured was to determine the validity of the "thick pipe" formula for brittle materials, and not to predict nor demonstrate directly the orientation of field fractures. Our conclusions concerning field results resulted from calculations involving the "thick pipe" relationship as well as considerations of overburden stresses, rock strengths, and the geometry and dimensions of the field system. Clark suggests that had the models been more porous or contained weak bedding planes, horizontal fracturing would have occurred. This is undoubtedly true provided external stresses similar to those in the earth's crust are nor imposed. However, if we were going to design experiments to represent directly the field case we would impose the proper stresses on the models. It is generally recognized that the vertical compressive stress in the earth's crust arising from the weight of the overburden is approximately 1 psi/ft of depth. Then, as an example, even though a horizontal bedding plane has zero strength, the formation cannot be separated to form a horizontal fracture unless the hydraulic pressure exceeds the stress due to overburden. And in those cases in which the stress resisting vertical fracturing is significantly less than that resisting horizontal fracturing, vertical fractures should result, notwithstanding horizontal plane weaknesses. We agree that breakdown pressure will be less if the fracturing fluid penetrates the formation. In Appendix HI of our paper it is shown that leak-off reduces the pressure necessary to initiate either a horizontal or vertical fracture. It would be difficult to attempt to

Jan 1, 1955

By Donald D. Carr, Lawrence F. Rooney

Perhaps no other mineral commodity in this volume has as many uses as limestone and dolomite. These carbonate rocks are the basic building blocks of the construction industry, the material from which aggregate, cement, lime, and building stone are made. Carbonate rocks, and their derived products, are used as fluxes, glass raw material, refractories, fillers, abrasives, soil conditioners, ingredients in a host of chemical processes, and much more. Although the use of limestone and dolomite goes back before recorded history, new uses and products are continually being found. Carbonate rocks form about 15% of the earth's sedimentary crust and are widely available for exploitation. Found extensively on all continents, they are quarried and mined from formations that range in age from Precambrian to Recent. Reserves of carbonate rock are large and will last indefinitely, although high-purity deposits may be absent or have limited availability in certain areas. Limestone and dolomite are so useful and so abundant that in 1981 about 644 million tons were produced in the United States. In fact, about 75% of all stone quarried or mined in the United States was carbonate rock. Sand and gravel was the only mineral commodity produced in greater quantity. In this chapter we have sought to provide the general reader an overview of carbonate rocks, their composition, distribution, production, and uses, with emphasis on developments since 1960. In doing so we borrowed extensively from our report on the applied geology of limestone and dolomite (Rooney and Carr, 1971). The literature on the limestone and dolomite resources of the United States and Canada is so vast and diverse that one can easily become mired down in print. We tried to list those references that are easily accessible and that provide the economic geologist with a starting point for more intensive study. The table listing selected references on limestone and dolomite resources in Canada and the United States reflects our attempt to be selective rather than exhaustive and to direct the reader to the best sources of detailed information. Geology Mineralogy The principal carbonate rocks used by industry are limestone and dolomite. Limestones are sedimentary rocks composed mostly of the mineral calcite (CaCO3) and dolomites are sedimentary rocks composed mostly of the mineral dolomite (CaCO3 • MgCO3). Aragonite (CaCO3), which has the same chemical composition as calcite but a different crystal structure, is economically important only in modern deposits, such as oyster shells and oolites. Aragonite is metastable and alters to calcite in time. Some other carbonate minerals, notably siderite (FeCO3), ankerite (Ca2MgFe (CO3)4), and magnesite (MgCO3), are commonly found associated with limestones and dolomites but generally in minor amounts. Because of their similar physical properties, the carbonate minerals are not easily distinguished one from the other. Specific gravity,

Jan 1, 1983

By M. J. Pool, P. J. Spencer, R. V. Miner

The partial molar heat of solution of rhodiunz in liquid lin and Rh-Sn alloys has been measured as a function of rhodium concentration at 700" , 725" , 750" , and 775°K. The values at infinite dilution are very exo thermic, ranging from -29,950 to -28,260 cal per g-atom Rh at 700" and 775"K, respectively. At the solubility limit the heat of solution becomes sharply nzore exothermic and has permitted the liquidus boundary to be accurately established at each temperature. From the measurements in the two-phase region of liquid Plus RhSn, the heat of formation of RhSn,has been determined. A large increase in the heat of formation between 725" and 750°K suggests the possibility of a phase transformation in this compound. V ERY few measurements have been made of the heats of solution of transition metals in liquid tin; yet those values that are available are of great interest because of their very large exothermic nature. For example, the heat of solution at infinite dilution of palladium in tin has been measured as about -26,000 cal per g-atom (625" to 800°~),' while the heat of solution at infinite dilution of platinum is even more exothermic at -28,591 cal per g-atom (914°K).' These values are of the same order of magnitude as the heat of formation of many intermetallic compounds and as such clearly indicate that very strong bonding must exist between the solute and tin atoms in the liquid solutions. Rhodium might be expected to display similar solution characteristics in tin in view of its close electronic resemblance to palladium and platinum. Since no thermodynamic data are available for Rh-Sn alloys, rhodium was chosen as an appropriate solute for further investigation of transition-metal systems with highly exothermic heats of solution at infinite dilution. EXPERIMENTAL The heat of solution of rhodium in dilute Rh-Sn liquid solutions was measured as a functlon of rhodium concentration at 700°, 725", 750°, and 775°K. The liquid-metal solution calorimeter used in this work has been fully described elsewhere. At the start of each series of measurements, the solvent bath consisted of about 0.5 g-atoms of 99.9+ pct pure tin. To this was made a number of small rhodium additions, about 0.0005 g-atoms each, and the heat of solution was measured for each. Rhodium of 99.9+ pct purity in powder form was used in order to promote rapid dissolution in the liquid tin. The powder specimens were contained in tin foil stated to be purified (lead-free). This foil was found to have no measurable heat of solution in pure tin so the effect of any impurities could be neglected. The energy equivalent of the calorimeter was experi- mentally determined by dropping tungsten and/or tin specimens at intervals during each series of rhodium drops. The tungsten and tin used were 99.95 and 99.999+ pct pure, respectively. The only heat effect due to a tungsten drop, or a tin drop at low rhodium concentrations in the bath, is the sensible heat of the specimen. The heat content of these materials is well-established and is presented in the compilation of Hultgren et al. 4 As the rhodium concentration in the tin bath increases a small heat effect arises from the reaction: Sn (liquid, T) — & (soln, T) where T is the temperature of the tin bath. Where necessary, this heat of solution of tin was determined by Gibbs-Duhem integration, and corrections were made to the heat effect of the tin calibration specimens and the tin foil containing the rhodium specimens. RESULTS AND DISCUSSION The calculated values of the heat of solution of rhodium in dilute Rh-Sn alloys as a function of composition are shown graphically in Fig. 1. The values have been corrected for the sensible heat of the specimens, initially at 273"K, and apply to the reaction:

Jan 1, 1969

By Fred C. Bond

MANY grindability tests have been made in the Allis-Chalmers laboratory since the tabulated results were last published.1 The lists of standard ball-mill and standard rod-mill grindability tests are brought up to date in this paper, and considerable new material is added. This includes a tabulation of open-circuit grinding tests, with impact and compression-crushing tests. The standard ball-mill grindability tests (Table I) were made in a standard laboratory ball mill, 12 by 12 in. inside, running at 70 rpm, with a charge of 285 iron balls ranging from 1 1/2 to 3/4 in. in diameter and weighing 20,125 grams. Seven hundred cubic centimeters of minus 6-mesh stage-crushed dry feed was used, and the circulating load was maintained constant at 250 pct by adjusting the number of mill revolutions for each grinding period. Tests were made at all mesh sizes from 28 to 200, and the grindability is the number of net grams of screen undersize produced per revolution. The table includes 479 tests. It has been calculated that the ball mill does 52.3 joules of useful work on the ore per revolution2 with 93 joules input to the mill. The standard rod-mill grindability tests (Table 2) were made in a laboratory tilting rod mill, 12 in. in diameter by 24 in. long inside, with a wave-type lining, running at 46 rpm. The grinding charge consisted of six 1 1/4-in. diameter and two 1 3/4-in. diameter steel rods, all 21 in. long and weighing 33,380 grams. The mill was rotated level for eight revolutions, then tilted up 5° for one revolution, down 5° for one revolution, and returned to level for eight revolutions. It was discharged by tilting downward at 45° for 30 revolutions. Feed consisted of 1250 cc of minus 1/2-in. stone and the circulating load was maintained constant at 100 pct by adjusting the number of mill revolutions for each grinding period. Tests were made at all mesh sizes from 3 to 65, and the rod-mill grindability is the number of net grams of screen undersize produced per revolution. The table includes 105 tests. The impact and compression-crushing tests (Table 3) are given in a table that is an extension of a previous publication3 in which the method of testing was described in detail. Two equal hammers weighing about 30 lb each were arranged to strike simultaneous blows on opposite sides of the test specimen. Broken pieces of rock passing a 3-in. square and retained on a 2-in. square were set up with the smallest dimension between the hammers, and this distance was measured. The height of the hammers was increased slightly before each blow until the specimen broke, and the foot-pounds per inch of thickness was calculated. Ten pieces

Jan 1, 1947

By R. A. Nadler and

Specimens of Ti-155A (Ti-5Al-1.3Fe-1.3Cr-1.2Mo), Ti-6Al-4V and Ti-4Al-3Mo-1V were hydrogenated, aged to high strength levels, and subjected to notched stress-rupture tests and tensile tests at two strain rates. Both Ti-6Al-4V and Ti-155A were relatively unaffected by hydrogen in amounts exceeding those found in commercially -produced material. Inconsistencies in the data and a relatively low initial level of ductility prevented an accurate determination of the hydrogen tolerance for aged Ti-4Al-3Mo-1V, but there was no evidence that this alloy was any more susceptible to hydrogen embrittlement than Ti-155A and Ti-6A1-4V. THE metallurgical literature has been considerably enriched in the past five years by the publication of the results of many investigations which were concerned with the effects of hydrogen on titanium and its alloys. Although it is not the authors' intention to provide a complete literature survey of this vast field, it is worthwhile to summarize briefly some of the important factors which influence the embrittling phenomenon in a - B alloys. Some of the influencing factors are those of strain rate, temperature, and composition. Hydrogen embrittlement in a - B titanium alloys is now believed to be a strain-induced phenomenon, and whether or not embrittlement occurs at high hydrogen levels depends largely upon whether sufficient time is available for hydride precipitation. This condition is fulfilled by tensile testing at very slow strain rates or by rupture testing smooth or, preferably, notched specimens over relatively long periods. Conveniently, maximum embrittlement occurs at or near room temperature. The presence B - stabilizing elements, in amounts sufficiently small to prevent an all -B or nearly all -P condition, promotes hydrogen embrittlement and this tendency is believed to be related to the amount of a - B interfacial area present. Aluminum, by increasing the solubility of hydrogen in A titanium, greatly reduces the danger of hydrogen embrittlement, while oxygen has been shown to have detrimental effects.' The preceding discussion is based on experimental work performed on annealed a - P alloys. This early work has shown that annealed Ti-6A1-4V and Ti-155A are relatively insensitive to the embrittling effects of hydrogen.' Although no work has been reported on annealed Ti-4A1-3Mo-lV, the alloy would be expected to exhibit somewhat the same insensitiv-ity as Ti-155A. In sharp contrast to the voluminous literature on annealed titanium alloys is the paucity of reported investigations on materials receiving commercial heat-treatment cycles. To date, the most notable contributions in this field have arisen from two separate investigations on Ti-3Mn-1Fe-1Cr-1V-lMo, the so-called 3Mn comple~. These investigations

Jan 1, 1961

By D. J. McAdam

In a series of papers the author and associates have discussed the influence of temperature on the tensile properties of metals.11-18 These papers present much information about the influence of temperature and the stress system on the conventional indices of mechanical properties, with special attention to the fracture stress. A recent study of the data, however, has revealed much additional information about the influence of temperature on the fundamental factors involved in the flow of metals. The present paper presents results of this study. Attention will be confined almost entirely to results derived from tension tests of unnotched cylindrical specimens at strain rates a little slower than those used in ordinary tension tests. According to a concept first presented by Ludwik and elaborated in recent papers by others,8,9,22,23 the mechanical state of a metal depends on the total plastic strain, but not on the temperature during straining, provided that the only structural changes are those essential to plastic deformation. In the summer of 1948, however, the author made the previously mentioned study of results of a general investigation by the author and associates and reached the conclusion that the mechanical state depends not only on the total strain, but also on the temperature during the straining. A number of diagrams were then prepared. These conclusions were presented without diagrams in a discussion last October of a paper by Dorn, Goldberg and Tietz.2 The metals used in the investigation on which this paper is based were Monel and oxygen-free copper. The Monel was supplied by the International Nickel Co. through the courtesy of Dr. W. A. Mudge. The copper was supplied by the Scomet Engineering Co. through the courtesy of Dr. Sidney Rolle. The data to be presented are based on results of tests at temperatures ranging between 165 and — 188°C. Description of the apparatus and methods of test are given in previous papers.1011'1"2 The present paper is the first part of the general discussion of the influence to temperature on the stress-strain-energy relationship for metals. The next paper will deal with metals that are subject to structural changes other than those induced solely by plastic deformation. Influence of Temperature and Plastic Strain on the Flow Stress of Monel and Copper For a study of the influence of temperature on the stress-strain relationship, flow-stress curves obtained with annealed metals at various temperatures will be compared with curves obtained with the same metals after cold drawing or cold rolling at room temperature. Diagrams thus obtained with Monel and copper are shown in Fig 1 to 8. Fig 1 to 7 show the variation of the flow stress with temperature and plastic strain; Fig 8 is a diagram of a different type, derived from Fig 4 to 7. In Fig 1 to 7 strain is expressed in terms of A0/A, in which A0, and A represent the initial and current areas of cross-section. Since values of Ao/A are represented on a logarithmic scale, abscissas are proportional to true strains; moreover, the true strains representing prior plastic deformation and those representing subsequent strain during a tension test are directly additive. Fig 1 shows flow-stress curves obtained with annealed Monel. Five of the curves are based on results of tension tests. Between yield and the maximum load, the flow was under longitudinal tensile stress; between the maximum load and fracture, the local contraction induced transverse radial tensile stress. The portions of curves designated F, therefore, represent flow with increasing radial stress ratio, the ratio of the transverse stress S3 to the longitudinal stress Si. Curve Fo is based on the ultimate stresses of specimens taken from bars that had been cold drawn various amounts.17 Since the tensile stress at the maximum load is unidirectional, curve Fo represents the course that a flow-stress curve would take if the stress during an entire tension test could be kept unidirectional. The flow-stress curve F obtained at room temperature (Fig 1) has been established accurately by numerous measurements of the diameter of the specimen during the extension from yield to fracture.17 At the time of the experiments, however, no apparatus was available for measuring the diameter during tension tests at low temperatures. Nevertheless, curves have been established to represent with sufficient accuracy the flow at low temperatures. Each flow-stress curve must be tangent to a curve U, which starts at a point representing the ultimate stress of annealed metal. Since the ultimate stress is based on the area of

Jan 1, 1950

By D. F. Gibbons

THIS investigation was undertaken in order to gain further information on the mechanism of plastic deformation of iron at temperatures between room temperature and 77.2°K, and also to contribute to our understanding of the low temperature brittleness phenomenon that occurs in both the tensile and impact tests. It was decided to study as pure an iron as possible and to use the tensile test in order to simplify the problem as much as possible. Two series of iron were used throughout the investigation: iron A, obtained from the National Research Corp. and supplied as gas-free, high purity iron; and iron B, obtained from Fast' and produced by melting iron under purified hydrogen. Typical analyses of samples of both irons, after fabrication, are given in Table I. The main difference in composition of these irons is in their nickel and oxygen contents. Nickel has a very large solid solubility in iron. However, oxygen has an extremely small solid solubility. An attempt was made to reduce the oxygen content still further by remelting under purified hydrogen. This was successful in reducing the oxygen content to -0.0006 pct. However, no results on tests of this iron are included, since difficulty was encountered in fabricating the ingot. Wire specimens 1 mm in diameter were used for the investigation. The wires were produced from the ingots, as supplied, by cold rolling and swaging, with intermediate vacuum anneals at 700°C. Care was taken to reduce preferred orientation in the wires to a minimum by keeping the percentage of reduction in area between anneals of 30 to 40 pct. X-ray transmission photographs showed no indication of preferred orientation. The final grain size of iron A was 20 to 24 grains per linear mm and of iron B, 16 to 18 grains per linear mm. All specimens were tested in the decarburized and denitrided condition unless otherwise stated. The procedure was to pass hydrogen, saturated with water vapor at 26oC, over the specimens at 730°C for 2 hr. The specimens were quenched in iced water and then annealed in a vacuum of - 5x10-6 mm Hg for one day at 560°C to remove hydrogen. Higher temperatures were used for the vacuum anneal but this temperature was found to be quite satisfactory in removing hydrogen. The specimens were slow-cooled after the vacuum anneal. Where it is stated that the specimens were carburized, the specimens were given an identical decarburizing treatment and then carburized by passing hydrogen, saturated with pure n-heptane vapor at 0°C, over the specimens at 730°C. The specimens were then homogenized at 730°C for 2 hr in dry hydrogen. After the carburizing treatment the specimens were given the same vacuum anneal as for the decarburized specimens. The carbon content was determined from the maximum value of the internal friction carbon peak.2 Apparatus Fig. 1 shows a diagram of the tensile apparatus, with half the cooling coil removed. The cooling coil enabled temperatures from 200" to 77.2oK to be obtained. The apparatus was designed to fit a standard Tinius Olsen Universal tensile machine. The apparatus consisted of the rigid plate A which carried two carefully machined hooks H. These hooks attached the plate firmly to the moving crosshead of the tensile machine. The lower grip B was rigidly suspended from the plate A by three stainless steel rods. The upper grip C was joined to the upper grip of the tensile machine by the stainless steel rod E. The lower part of the apparatus was surrounded by a Dewar vessel, which was made a sliding fit inside the brass collar D. The specimen was held in a clamp which fitted into the grip, forming a ball and socket joint (see inset a). This was to attain as nearly pure

Jan 1, 1954

By D. Cox, M. J. Duggin, L. Zwell

The carbides of approximate composition and Mn have been studied using X-ray diffraction techniques. Those carbides of the type (Fe,Aln)zC ave isostructural with cementite. The cell pararmeters a and c have minimum values at approximately 10 at. pd substitution of manganese for iron; no satisfactory explanation has yet been found for this phenomenon. The carbide fFeMn4)C has a monoclinic unit cell whose dimensions are close to those of ,11,15Cz A neu-troip-dij~ractiot~ study of (F'eAlrz4)C~ reveals that, like MnsCZ, it is isostructural with Pd5Bz. The iron and manganese atoms occupy the palladium atom sites, while the carbon atoms were found to have the same atomic coordinates as the hovon atoms. A neutrorr-diffraction study of indicates that the carbon-atom positions are very close to those occupied in (Fez.,ll/lr~,.3)C. In both carbides studied, tlre iron and manganese atomzs were found to be essentially randomly distributed, although, in the case of (Fe,.811fn1.2)C, it is possible that there may be a slight preference of manganese atoms for- the general (d) positions and a corresponding slight preference of iron atoms for the special (c) positions. It has been found that a complete range of solid solution exists between Fe3C and Mn3C at 1050°C,I although Mn3C becomes unstable when the temperature is reduced to 95O0C,' and can only be retained by rapid quenching. It is also known that a complete range of solid solution exists from Fe5Cz to M~SC~,~ although the stability range of carbides of the type (Fe,Mn)sCz as a function of the relative proportions of iron and manganese is not known. X-ray examinations of Oh-man's carbide3 and Spiegeleisenkristall,~ which have the approximate compositions (Fe3.67Mnl.33)C2 and (Fe3-,Mn,)C, where x lies between 0.4 and 1, respectively, have been made. The following carbides have also been studied: ] The lattice parameters determined during these investigations are listed in Table I. It is seen that carbides of the type (Fe,Mn)sCz have a monoclinic unit cell while carbides of the type (Fe,Mn)3C have an orthorhombic unit cell. It is evident that the variation of lattice parameters with manganese content is not linear for carbides of the type (Fe,Mn)3C. The coordinates of the atoms in (Fe2.7Mno.3)C have recently been determined by single-crystal analysis., The fractional atomic coordinates have been shown by Fasiska and jeffrey to be in good agreement withj those deduced from an earlier analysis of Fe3C by Lipson and etch.' However, it was impossible to determine whether iron and manganese atoms occupied ordered positions because of the small difference between the atomic scattering factors of iron and manganese. The atomic positions in Mn5Cz (Refs. 8 and 9) and Fe5C2 (Refs. 7 and 8) have been obtained only by comparisons made with the isostructural compounds P~SB~.' Since X-ray diffraction techniques were used in these investigations, accurate positioning of the carbon atoms, which have a low atomic scattering factor, was difficult. No attempt has been made to determine the atomic positions in the other carbides previously studied. It was felt that an investigation of the lattice parameters of a number of intermediate carbides of the types (Fe,Mn)sCZ and (Fe,Mn)& would be of interest. It seemed likely that a neutron-diffract ion study of such carbides would indicate whether ordering occurred between the iron and manganese atoms because of the large difference between the neutron-scattering cross sections of iron and manganese. It also seemed probable that such an investigation would provide a determination of the atomic coordinates of the carbon atoms. I) EXPERIMENTAL DETAILS Specimens, each weighing approximately 20 g, were carefully prepared according to the following proportions: The components were 500-mesh powders of 99.995 pct purity iron and spectroscopically pure carbon and a 200-mesh powder of 99.995 pct purity manganese. The component powders were intimately mixed by prolonged shaking, then each specimen was inserted into a spot-welded cylindrical container of tantalum foil, whose end was closed but not sealed. Each specimen in its envelope was then sintered at 1050° C for 24 hr in a thin-walled evacuated quartz capsule, such a time having been previously found sufficient for equilibrium to be attained.' Each specimen was then quenched in order to attempt to retain the high-temperature phase, as the literature indicates that transformations may occur on cooling. Debye-Scherrer X-ray photographs were taken of each specimen using a 114.6-mm-diam camera, Fig. 1, patterns 2 to 6. The exposure time was 6 hr using filtered iron radiation at a tube voltage of 40 kv and a tube current of 12 ma.

Jan 1, 1967

By P. R. Sperry

Aluminum alloy 3S was examined to determine the relationship between the applicable phase diagram and the microstructures produced under conditions tending toward non-equilibrium as well as equilibrium. The only phases present in the alloy were aluminum, (Mn,Fe)A16, aAI(Mn,Fe)Si, and silicon. It was found that the alloy behaved essentially as though it belonged to the Al-Mn-Si system. The effects of various solidification rates and of various homogenization treatments were studied. Particular attention was given to a nonuniform precipitate distribution which reflected the solute distribution in the cast aluminum solid solution. Some possible explanations for the solute distribution based on characteristics of the Al-Mn-Si or AI-Mn systems were discussed. ALUMINUM alloy 3s (aluminum plus 1.25 pct Mn) is one of the oldest and the most widely used of the wrought aluminum alloys. In spite of this fact, there are some aspects of the metallography of 3s which have not been revealed or which have not been brought to the attention of the metallurgist working with the alloy. The purpose of this investigation was to demonstrate in general how some of the fundamentals of equilibrium phase diagrams may be applied to the study of an alloy and to show specifically the origin of certain metal-lographic features of aluminum alloy 3s. It is hoped that this information can serve as a basis for solving specific problems relating microstructure to such things as mechanical properties, recrystallization, grain growth, and surface finish. Literature Survey Considering that alloy 3s contains only one alloying addition, manganese, it seems reasonable that it would have characteristics pertaining to the A1-Mn system. However, those characteristics might well be altered by the impurities in the commercial alloy, especially iron and silicon. Therefore, not only the simple binary but also more complex systems had to be reviewed. One characteristic of the A1-Mn system worthy of mention is its sluggishness, which leads to conflicting solubility data1,'2 and makes it relatively easy to supersaturate the solid solution and shift the eutectic composition during rapid solidification.3'4 Another characteristic is the strong influence which small amounts of impurities, especially iron, exert in lowering the solid solubility of manganese. Undissolved manganese is ordinarily present in an intermetallic phase, MnAl,, having an orthorhombic crystal structure. A metastable precipitating phase designated as "G" has been reported but small amounts of iron and silicon suppress this phase entirely." Raynor has established that MnA1, can dissolve a considerable amount of iron by a substitution of iron for manganese atoms up to 50 atomic pct.'" Consequently, in the Al-Mn-Fe liquidus diagram, the eutectic valley separating the primary aluminum and (Mn,Fe)Al, fields is fairly close to a constant Mn + Fe content." This is an important consideration in regulating the composition of 3s and also other aluminum alloys with high manganese additions."

Jan 1, 1956

By J. Nutting, D. H. Kirkwood

IT is clearly of great metallurgical interest to know the conditions of temperature and gas composition under which different products will form in sulphide roasting. A method of graphically presenting the relevant thermodynamic data will be demonstrated, such that these conditions are readily obtainable. The products of sulphide roasting may be oxide, sulphate, or oxysulphates, and the equilibria between these may be represented in the following way: 2 MSO4 = MOMSO4 + SO3 MO-MSO4 = 2 MO + SO3

Jan 1, 1961

By Mark R. Lamb

THE treatment of slime is of special interest to those engaged in cyaniding gold- and silver-ores. The usual practice is to make as small a percentage of slime as possible. In many instances the slime is given no treatment, but is impounded in dams in the hope that the future will develop some method of economically treating this product. The filter-press was the first step upward from ordinary decantation, but, on account of heavy labor-charge and high cost, its use has been limited to high-grade material. The slime at the Combination mill at Goldfield averages perhaps $20 per ton, and although the values are quickly dissolved the filter-press installation was not entirely satisfactory, resulting in the erection of the canvas-cell filter, developed by Chas. Butters and his staff. This filter is a great improvement over the filter-press, and the following description of it will be of interest to those engaged in treating slime produced in crushing ores, especially in view of the fact that, by the use of this filter, slime can be treated at a lesser cost and with a higher percentage of gold- or silver-extraction than in the ordinary treatment of sand; and, moreover, the initial outlay for an all-sliming plant is less than that for the ordinary sand- and slime-plant. The economy of construction of the Butters filter-plant is clearly shown in Figs. 1 and 2, which illustrate the 20-frame filter of 40 tons capacity, now being installed by the Nevada Goldfield Reduction Co. In this construction the slime-pump is so connected that it can pump to or from either tank or the filter by changing the valve-settings-a combination which is necessary in this particular installation, since sufficient fall is not available for filling and discharging the filter by gravity. Ordinarily, a slime-plant comprises: a filter-box with frames

Jan 1, 1907

By R. M. Foley, Hidesaburo Kurushima

About 25 pct of all Japanese pyrite comes from the Yanahara mine on Honchu Island. For the past 40 years lack of an economical recovery process forced the operator, Dowa Mining Co., to sell the pyrite for its contained sulfur only, leaving 7 million tons of high grade iron ore waste. Introduction of the Dorrco FluoSolids roasting system enabled the company to make use of its iron sulfides and also to find a market for the mine's considerable pyrrhotite production. The success of this project is of great importance to Japan's economy since it points the way toward material increase in her natural resources. Prior to use of the present process, no practical method was available for commercial exploitation of known large tonnages of pyrrhotite, while the copper content of the even larger deposits of pyrite made the roasted residue unacceptable for the iron and steel industry. In Europe, large amounts of copper-bearing pyrite cinder are given a chloridizing roast for separation of the copper and iron but this treatment is not economical in Japan where salt is imported from as far away as Egypt.

Jan 10, 1958

By S. G. Cupschalk, J. J. Wert, R. A. Buchanan

The uniaxial deformation of thermally ordered and disordered polycrystalline Cu3Pt was studied by means of the X-ray line - broadening analysis according to Warren and Averbach and the extension of this analysis to ordered fcc materials by Mikkola and Cohen. Because of the heat treatment history, extinction had a pronounced effect on the X-ray spectra of ordered and disordered C%Pt at small plastic strains. After an appropriate correction for extinction, the long-range order in thermally ordered ChPt was found to decrease at a slow constant rate with plastic strain. Furthermore, the antiphase domain probability increased at a constant rate to 17.5 pct strain. The effective particle size behavior indicated that the stacking fault energy is lower in ordered than in disordered Cu3Pt. Analysis of the stress-strain curves shouled that ordered Cuzt has a slightly lower yield Point but a much higher work-hardening rate than disordered Cu3Pt. THE presence of long-range order in a solid-solution alloy has a marked effect on its mechanical properties. While this effect has been known qualitatively for many years, it was not until recently that detailed investigations have been performed to determine the exact role long-range order plays in this strengthening mechanism. The development of an advanced, quantitative. X-ray diffraction analysis by Warren and Averbachl and the extension of this analysis to the L1, type super lattice by Mikkola and cohen2 have greatly accelerated research in this field. The research reported in this paper consisted of two primary phases. The first phase was to determine the effect of long-range order on the tensile properties of polycrystalline Cu3Pt. This objective was accomplished by comparing the stress-strain behavior of thermally ordered CusPt to that of thermally disordered CusPt. The second phase of the research was to determine the difference between the atomic arrangements in thermally ordered and thermally disordered Cu3Pt as a function of uniaxial deformation and thereby gain a deeper insight into the mechanism by which long-range order affects the tensile properties. This second objective was accomplished by applying the Warren-Averbach method of peak profile analysis to the X-ray diffraction patterns obtained from ordered and disordered Cu3Pt after given amounts of uniaxial deformation. EXPERIMENTAL PROCEDURE The Cu3Pt was prepared by vacuum melting and casting. After a homogenization anneal, the ingot was cold-rolled to sheet form. Two tensile specimens with gage sections of 2.50 by 0.500 by 0.115 in. were carefully machined from the sheet. The specimens were polished with a final step of 600-grit paper to insure smooth diffracting surfaces. Finally, one specimen was heat-treated to yield an average grain diameter of 0.016 mm and an initial degree of long-range order, S, of 0.825. The other specimen was water-quenched from above the critical temperature, 645"C, to yield an average grain diameter of 0.017 mm and zero long-range order. The heat treatment history of each specimen is shown in Table I. The tensile tests were performed utilizing a Research Incorporated Model 900.95 Materials Testing System. This unit employs a closed-loop feedback system which maintains a constant strain rate through an extensometer clipped to the gage section of the tensile specimen. A strain rate of 1.32 i0.02 x 10"4 sec-' was employed in testing both specimens. In the X-ray diffraction analysis, a General Electric XRD-5 diffractometer equipped with a pulse-height analyzer set for 90 pct efficiency was employed. The goniometer speed selected was 0.2 deg, 20, per min. Filtered Cu radiation was used for all peaks and all peaks were chart-recorded. Because of nonuni-form grain size. it was necessary to spin the specimens during X-ray analysis in order to obtain reproducible integrated intensities. The spinning rate was 2000 i100 rpm. The application of the Warren-Averbach method of peak broadening analysis to a diffraction pattern is very time consuming if done manually. In this research, the calculations involved were performed with the aid of a computer program by wagner.3 As reported by Wagner, the program is written in Fortran TV computer language. It was modified to Fortran I1 so as to be handled by the IBM 7072 computer at Van-derbilt University. In the X-ray diffraction analysis of uniaxially deformed Cu3Pt, the 100, 200. 400. 111, and 222 reflections were recorded from the initially ordered sample after 'plastic strains of 3.0, 6.0, 9.0, 12.0,

Jan 1, 1970