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By G. H. Lair, B. J. Langford, S. H. Rockwood

Volumetric reservoir analysis is dependent upon a firm relationship between porosity, connate water, and net pay. Capillary presswe data on core samples interrelate these three factors. It is shown that other reservoir problems may be resolved: (I) whether water production is indigenous or extraneous to the oil producing interval, (2) an oil-water contact is being approached, and (3) whether the oil-water contact is remotely situated. INTRODUCTION Volumetric analysis of carbonate reservoirs challenges the ingenuity of the reservoir engineer due to an absence of data on the minimum oil-bearing and oil-producing pores of a given reservoir rock. Many studies with this as the ultimate objective are available on the physical characteristics and fluid-flow behaviors of the carbonate reservoir. Some of these publications are: Hassler, Brunner and Deahl's1 study of the role of capillarity in oil production; Jan Law's' and A. C. Bulnes3 statistical treatment of core analysis data; Chalkley's method4,5 or estimating specific surface areas and porosity; Walter Rose's9 report on porosity, reserves, and recovery; Archie's78 tex-tural classification of carbonate rocks; Stewart and Spurlock's9 large core analysis; Purcell'sl0 mercury injection method in studying capillary pressure phenomena; Stewart, Craig, and Morse's" model multiple-phase flow test in investigating the relative permeability effect; and Fatt's12 illumi-nating network model of porous me-dia. The present study shows that the static reservoir parameters of porosity, connate water, and net productive thickness may be interrelated by utilizing a statistical arrangement of mercury capillary pressure data in addition to routine core analysis. This report summarizes the techniques that are followed to obtain these relations which are applicable to carbonate and sandstone reservoirs that do not contain clays. Capillary pressure statistical studies may be extended to include clay-bearing rocks by utilizing water capillary pressure data. POROSITY It is necessary that the laboratory determine the porosity value that includes the pores that are filled with hydrocarbons and water in the reservoir. This is the reservoir's total porosity. When true total porosity values are not available, correlations between porosity and capillary pressure characteristics lose their significance since they cannot be related through oil-base mud core data to the porosity and connate water values found in reservoirs. Methods used to obtain reliable porosity data will not be discussed; however, care should be exercised in selecting samples for measurement. In homogeneously porous formations, small core plugs sampled frequently will furnish adequate information. Where heterogeneity, particularly vugginess, is present, large-core porosity data may be required; in such reservoirs a suite of capillary pressure samples may be selected from small plugs. CAPILLARY PRESSURE Capillary pressure data provide a method for relating two carbonate reservoir parameters, porosity and connate water saturation, which in turn permit delineation of productive rock types necessary for selecting the third parameter, productive thickness. Specific applications of capillary pressure data to petroleum reservoirs were suggested by Archie8 and Pur-cell."' Their work demonstrated that in preferentially water-wet rock, porc

Jan 1, 1958

By W. T. Lankford, Low J. R., Gensamer M.

The problem of sheet metal formability is one which has received a vast amount of attention during recent years. In spite of the great amount of study and experimental work which has been devoted to this problem, many of its fundamental aspects are still not well understood. No generally satisfactory test has been developed which is capable of providing a reasonably adequate evaluation of sheet metal formability, particularly in borderline cases. Anomalies persistently arise in which a lot of supposedly satisfactory material exhibits excessive breakage when formed. This lack of suitable technological tests greatly hinders the design of new parts and the substitution of new materials in parts of existing design, frequently necessitating considerable costly trial and error experimentation on an actual production basis. When, at the beginning of World War 11, it became apparent that mass production methods of forming would bc required in the aircraft industry if the large numbers of military aircraft needed were to be supplied, renewed attention was focussed on the problem of sheet metal formability. It was widely recognized that the formability requirements placed on sheet metals had been steadily increasing in severity through the years, especially as greater and greater emphasis was placed on streamlined designs with the resultant freer application of doubly curved surfaces. The recent advent of several new high strength aluminum alloys of more limited formability than some of the older alloys also served to emphasize the need for new approaches to the problem. It seemed evident that progress in the technology of sheet metal forming would depend to a great extent on the development of more reliable criteria for the evaluation of formability. Although the problem is admittedly very complex, the most logical approach appeared to be through studies of the fundamental material properties of sheet metals, especially under complex stress conditions simulating those prevailing in actual forming operations. Most previous investigations have failed to make full use of existing knowledge of fundamental material properties, and in many instances have attempted correlations with properties which are now recognized as being ambiguous in nature. This investigation was undertaken, therefore, with the aim of developing methods for the study of the flow and fracture of sheet meta1s under combined stresses, and of making use of such methods as might be developed in

Jan 1, 1947

By Low J. R., Gensamer M., W. T. Lankford

The problem of sheet metal formability is one which has received a vast amount of attention during recent years. In spite of the great amount of study and experimental work which has been devoted to this problem, many of its fundamental aspects are still not well understood. No generally satisfactory test has been developed which is capable of providing a reasonably adequate evaluation of sheet metal formability, particularly in borderline cases. Anomalies persistently arise in which a lot of supposedly satisfactory material exhibits excessive breakage when formed. This lack of suitable technological tests greatly hinders the design of new parts and the substitution of new materials in parts of existing design, frequently necessitating considerable costly trial and error experimentation on an actual production basis. When, at the beginning of World War 11, it became apparent that mass production methods of forming would bc required in the aircraft industry if the large numbers of military aircraft needed were to be supplied, renewed attention was focussed on the problem of sheet metal formability. It was widely recognized that the formability requirements placed on sheet metals had been steadily increasing in severity through the years, especially as greater and greater emphasis was placed on streamlined designs with the resultant freer application of doubly curved surfaces. The recent advent of several new high strength aluminum alloys of more limited formability than some of the older alloys also served to emphasize the need for new approaches to the problem. It seemed evident that progress in the technology of sheet metal forming would depend to a great extent on the development of more reliable criteria for the evaluation of formability. Although the problem is admittedly very complex, the most logical approach appeared to be through studies of the fundamental material properties of sheet metals, especially under complex stress conditions simulating those prevailing in actual forming operations. Most previous investigations have failed to make full use of existing knowledge of fundamental material properties, and in many instances have attempted correlations with properties which are now recognized as being ambiguous in nature. This investigation was undertaken, therefore, with the aim of developing methods for the study of the flow and fracture of sheet meta1s under combined stresses, and of making use of such methods as might be developed in

Jan 1, 1947

By F. E. Harris

IT has been said that carbon is "ubiquitous" with reference to iron alloys. Certainly at temperatures where carbon and iron form the solid solution, austenite, it may be readily added to, or removed from, steel surfaces. Differences in concentration potential are the sole requirements for carbon flow. To measure the rate of carbon flow, a proportionality factor, D, is arbitrarily expressed in units of (length)2 X (time)-1. If the term used for carbon concentration is weight pct C, as is customary, then a factor must be found, changing this term to one of weight per unit volume when the diffusion coefficient, D, is used in flow formulae. The determination of D values for carbon is complicated since D varies with the concentration itself. The Matano method has often been employed for systems involving a variant D. This method expreses an evaluation of unsteady flow whereby .0 is given a relationship with the numerical value of C-X slopes. The author contends that this solution is incorrect, except when D is invariant. The primary purpose of this paper is to support this contention. Here diffusion is controlled under two conditions of flow. The first part deals with the steady state; the second part considers unsteady flow under definite restrictions. Beginning with the fundamental concept, the development is continued until equations for C-X curves are obtained. The solution is then subjected to experimental data. It is realized that a great amount of reliable data is required for a rigorous proof. No such research is claimed. Instead an attempt is made to set forth the physical significance of the diffusion mechanism, and particularly to express the factors which influence the flow of carbon in austenite. Of equal significance is the implication that gaseous media may be employed with reasonable accuracy in making these determinations. THE STEADY STATE DIFFUSION OF CARBON IN AUSTENITE The fundamental experiment defining diffusion in solid metals may be readily performed for carbon in gamma iron.1 For this experiment a thin plate of medium carbon steel is chosen. The surfaces of this plate may be taken as two parallel planes which may be considered infinite for points well in the center of the planes. These two planes are held at different concentrations of carbon at a constant temperature. This temperature, 1700°F for example, is one where the single phase, carbon in solution with gamma iron, exists. When these conditions are held for a sufficient time the concentration of carbon of the different points of the solid settles down towards its steady state value. At points well removed from the ends, the concentration will remain the same along planes parallel to the surfaces of the plate.

Jan 1, 1947

By R. D. Pehke, J. F. Elliott

The solubility of nitrogen in liquid pure iron has been measured as a function of pressure and temperature. Sieverts' Law is obeyed at all pressures up to 1 atm and the temperature coefficient of solubility is 8 x 106 %/°C. The solubility of nitrogen in liquid iron alloys has been studied by the Sieverts' method and by the sampling method. The solubility is decreased by Al, C, Co Cu, Ni, 0. Si atzd Sn, and it is increased by Cr, Cb, Mn, Mo, Ta, W, and V. Interaction coefficients eN, for the effect of these elements on nitrogen dissolved in liquid iron, are reported. 1 HE effect of nitrogen upon the properties of steel may be desirable or not, depending upon the composition, processing treatment, and the use of the product. To understand the behavior of nitrogen in the various steelmaking processes, it is necessary to have reliable fundamental data pertaining to the solubilities, rates of solution, and chemical reactions of nitrogen in steel, particularly in the liquid state. In an effort to extend our knowledge in these areas, an investigation of the thermodynamics and kinetics of the solubility of nitrogen in liquid iron alloys was undertaken. The thermodynamic aspects of the study are presented in this paper. The solubility of nitrogen in liquid pure iron has received the attention of several investigators whose results are summarized in Table I. Reasonable agreement among the more recent researches1-l8 is found for the solubility of approximately 0.044 wt pct at 1600°C and 1-atm pressure of N,, for the temperature coefficient of solubility being small but positive, and for adherence to Sieverts' law for pressures up to 1 atm. The solubility of nitrogen in a number of binary liquid iron alloys has been studied: Fe-A1,3 Fe-As,16 Fe-C,3, 71 18, 19, 20 Fe-Cr ,5,7, 10, 11, 15 21 Fe-Co,'4, '6 Fe-Cu,'6 Fe-Mn,10, 11, 18 Fe-Mo,13, 16 Fe-Ni,10, 11, 13, 15, 16 Fe-O,16, 17, l9 Fe-p,7 Fe-S,l6 Fe-Sb,16 Fe-Se,lg Fe-Si,4) 97 '77 187 19>2° Fe-Sn,16 and Fe-v.51 '3 A few ternary systems have been studied, most of which were reviewed by ~angenberg" who applied the methods suggested by WagnerPz3 Chipman,24 Darken and G~rry, and Morris and Buehl26 for predicting solubilities in multicomponent alloys. Langenberg22 and Humbert and Elliott15 compared the solubilities computed from simple binary alloy systems with experimental measurements of nitrogen solubilities in liquid iron containing two or more alloying elements. The interaction coefficients for nitrogen in

Jan 1, 1961

By R. E. Hobart

THE development of the Hauto power plant and the claims made by various engineers that electricity was more economical than steam for power purposes in and about the mines; led the Lehigh Coal and Navigation Co., in 1911, to conduct a test to ascertain the consumption of power used by a large steam hoisting engine. The question being one on which no reliable information could be found, and the opinion of various engineers differing to such an extent, it was decided that a test under actual operating conditions was necessary. This was arranged for at one of the collieries. The engine selected was a 30 by 60-in. (76.2 by 152.4-cm.) piston-valve engine of modern type, and comparatively new. The boiler plant from which the hoist received its power was about 600 ft. (182 m.) from the engine, the latter being fed by a 10-in. (25.4-cm.) steam line insulated with magnesia pipe covering. Two batteries of the boiler plant, aggregating 1200 hp., were cut off and fed into a separate steam line leading direct to the hoist engine. Steam auxiliaries, consisting of feed-water pump and blowers, were fed by the boilers in test, and their consumption charged against the hoist. A barrel weighing device was installed, as it was felt that this would be the most accurate way of determining the consumption of water. The fuel, No. 3 buckwheat, was carefully weighed, and every precaution was taken to make the test accurate in every particular. The test was run for 1 week, or a total of 168 hr. A record was kept of the number of trips hoisted or lowered, and continuous indicator cards were taken. One particular set of cards was taken with the hoist operating balanced, and also with the hoist operating with no counterweight other than the empty cage, the coal in the car being weighed. The speed of the hoist was taken by means, of a graphic recording instrument which registered the number of revolutions every 5 sec. From this record, speed-time curves were plotted.

Jan 2, 1918

By R. P. Alger, J. S. Wahl, C. W. Johnstone, J. Tittman

To overcome the problems of mudcake and hole irregularities, the new compensated formation density logging device employs two detectors spaced at different intervals from the source. The detector at the shorter spacing is particularly sensitive to the density of material immediately adjacent to the face of the pad. The contribution of this material, which includes mudcake and minor wall irregularities, affects the response of each detector to a different degree. The signals from both detectors are combined to give automatically a density correction which is added to the uncompensated density information from the detector at the longer spacing. Both the compensated density measurement and the amount of compensation are then recorded on the log. With the unwanted borehole effects removed, the measurement is recorded directly in terms of bulk density on a linear scale. INTRODUCTION Formation density is a very useful and diagnostic parameter for formation evaluation. When the matrix lithology is known, porosity is accurately and readily computed from density data.'-" hen the density measurements are used with other porosity-responsive measurements, such as Sonic and Neutron logs, both lithology and porosity may be defined with good accuracy."'" In the past, however, application of density logs was complicated by the need for corrections to obtain true formation densities from the log values. Mudcakes on permeable formations and roughness of the borehole walls have been particularly troublesome since, in each case, the device was prevented from direct contact with the formation. ~ttempts have been made to correct for such conditions by making a simultaneous caliper log, estimating pad standoff, and assuming knowledge of the composition of the intervening materials. While such corrections are frequently accurate, they are not always reliable because of uncertainties in the pad standoff and material composition. A further complication was the necessity for manual conversion of log counting rates to density. The newly introduced compensated formation density logging device (FDC) automatically corrects for mudcake and minor wall irregularities, and provides a log that is scaled directly in bulk density. The purpose of this paper is to discuss briefly the equipment, to explain the theoretical considerations involved in the compensation, and to present field examples of this new log. COMPENSATED FORMATION DENSITY EQUIPMENT The FDC tool incorporates two major advances over previous density logging devices: a second detector, spaced closer to the gamma ray source to provide increased sensitivity to material close to the pad; and the automatic derivation of a corrected formation density from the two counting rates. The arrangement of the source and detectors is shown in Fig. 1. The longer spacing detector is at the same spacing from the source and gives the same response as the single detector in the previous uncompensated device. By appropriate combination of the outputs of the two detectors, the FDC tool is able to compensate in large measure for effects of mudcake and hole rugosity. This is achieved, as will be described later, without specific knowledge of either the thickness or the nature of material between the pad and formation. To obtain corrected formation density, the two count rates are automatically processed by an analog computer in the surface panel. More specifically, the computer derives a value of uncompensated formation density from the long-spacing counting rate, computes a density correction p using both the long and the short spacing count-

Jan 1, 1965

By L. D. Wooddy, W D. Moore

Water drives may be natural or artificial. For the case of a natural water-drive reservoir, the volume of water influx corresponding to reservoir pressuve-production history may be calculated by means of the material bal-ance-unsteady state equation. For the artificial water drive or pattern injection water flood, the material balance equation is sufficient to determine the reservoir pressure-production, injection history. However, for both cases to reliably predict reservoir performance, it is necessary to calculate the saturation distribution of the reservoir fluids at different stages of depletion. This paper consolidates the various equations used to determine the performance of water-drive reservoirs into one approach that is applicable to any reservoir that has a natural water drive or is to be produced by water injection. The numerical procedure has been simplified considerably, and the solution is adaptable to either hand calculations or digital computer operation. INTRODUCTION A water-drive reservoir is one in which the predominant source of energy for production of oil results from the encroachment of water into the oil zone. The water drive may be either natural or artificial. Water invasion takes place by expansion of water in the aquifer as a result of the pressure decline transmitted from the oil reservoir. Because water is practically incompressible, a natural water drive requires an extensive aquifer containing a volume of water many times the volume of oil in the reservoir. Provided the permeability of the formation is reasonably high, a moderate decline of reservoir pressure usually causes adequate water influx. If the reservoir pressure decline becomes excessive because of low permeability or a high rate of production, part of the oil production may result from the liberation of dissolved gas which may cause excessive quantities of free gas to be produced along with the oil. In extreme cases, reservoir pressure may decline so rapidly as to cause the oil to be produced predominantly by dissolved gas drive, leading to abandonment of the reservoir before the advancing water has had a chance to flush the entire reservoir. In reservoirs that are not connected to extensive aquifers, an artificial water drive may be created by injecting water into or around the lower edges of a reservoir. Sometimes water is injected within the reservoir through wells interspersed among the producing wells. When injection is carried out in such fashion, the reservoir is in effect subdivided into a number of separate producing segments, in each of which the displacement of oil is simultaneously taking place, and the shapes of which are fixed by the well patterns employed. Because of the natural tendency of oil and water to segregate according to their relative densities, water will tend to migrate downdip regardless of where injected. Accordingly, the injected fluids can be retained within the intended reservoir segments only when conditions are unfavorable to their migration under the force of gravity. Pattern injections are usually confined, therefore, to reservoirs having low structural relief, or low permeability, or when production and injection rates are such that fluid movements in the reservoir are dominated by the pressure gradients rather than by the force of gravity. In the case of natural water-drive reservoirs and artificial water-drive reservoirs in which the water is in-

Jan 1, 1958

By Pedro Bolsaitis, Paul M. Sullivan

The activities of zinc in liquid Zn-A1 alloys in the temperature range of 1000' to 1160°K were determined by means of the isopiestic technique. The results are in reasonable agreement with the previous values determined by the radioactive tracer dew point method. Slightly negative excess entropies of mixing were found. The measured activities suggest that the a function for these alloys is not a constant, as has frequently been assumed, but is better represented by two straight line segments of different slopes. THE thermodynamic data of Al-Z alloys have been the subject of investigation by many authors. However, in the recent compilation of such data by Hult-gren et al.,5 the conclusion is reached that none of the previous measurements are reliable or consistent with each other. Since then two new sets of measurements have been reported,2, 8 both laying claim to good accuracies, yet showing very significant disagreement in the measured activity coefficients of zinc in Zn-A1 alloys and in the magnitude and sign of the excess entropies of mixing. Information on the thermodynamics of the system is of particular interest for interpreting results of the Bureau of Mines research on refining of zinc by distillation of die cast scrap. The present study was undertaken as an independent check on the recent values obtained from dew point measurements 8 and electromotive force measurements.' The isopiestic method, first introduced by Herasy-menko,' was chosen because it is one of the few thermodynamic techniques which, in principle at least, is carried out under perfect equilibrium conditions. The disadvantage of the method is that data are obtained at various random compositions and temperatures, which makes the interpretation somewhat cumbersome. The random errors arising from temperature gradients and fluctuations appear to be comparable to those of other techniques, as evidenced by the magnitude of the estimated uncertainties in the calculated integral thermodynamic quantities. An advantage of the method is that each run provides a relatively large amount of data; thus, it is less time consuming than those methods where samples of each composition must be measured individually. EXPERIMENTAL PROCEDURE The activities of zinc in liquid Zn-A1 alloys were measured by the isopiestic method which uses the principle of the tendency toward equilibrium of a series of samples at various temperatures under an imposed vapor pressure of the more volatile component (i.e., zinc, in this case). The basic apparatus is shown in Fig. 1. The small graphite crucibles that contained the metal samples were 15 mm ID and 13 mm deep with approximately 3 mm wall thickness. Each crucible was notched at the top and bottom to allow easy passage of vapor from one to another. At the beginning of an experiment each crucible was weighed on an analytical balance. Zinc and aluminum of spectrographic purity (99.999 pct) were then weighed into the crucibles. The amounts of each metal added were such that initial compositions varied above and below equilibrium values. One crucible contained only pure zinc. After weighing, the crucibles were stacked with the pure zinc on the bottom (corresponding to the lowest temperature in the furnace) and richest aluminum alloys at the top. The entire stack of crucibles was then placed in the silica glass container "E" which was then sealed under vacuum. Sixteen to twenty calibrated Chromel-Alumel thermocouples were attached by wire along the length of the silica tube, one or more adjacent to each sample.

Jan 1, 1970

By M. T. Simnad, C. E. Birchenall, M. H. Davis

The transport mechanisms previously proposed in wustite and hematite have been confirmed by movement of inert markers during the growth of these phases. By similar measurements the mechanism for magnetite has been determined. The rates of growth of multilayer scales on iron, magnetite and hematite on wustite, and hematite on magnetite have been measured. All rates are found to be diffusion controlled under the conditions of the experiments. MANY studies have been made of the rate of oxidation of iron in both the low temperature, film-forming region and the high temperature scaling range. The interpretation of the results has been greatly complicated by the formation of multi-layered scales containing the various phases of iron oxides. In most cases no attempt has been made to determine the composition of the scale and its dependence on time of oxidation in a quantitative manner. However, several recent studies'" have taken this variable into account. It is generally agreed that the growth of thick scales on iron follows a parabolic rate law for the increase of total thickness. In the thin film region agreement has been claimed with parabolic equations of various forms and with logarithmic equations. The discussion here will be concerned only with the formation of scales of considerable thickness, greater than several microns, although some of the conclusions may be applicable to thin films. A brief survey will be given of the solid phases of the iron-oxygen system, the stability of wiistite in bulk and in thin films, the relative densities of iron in the metal and oxides, and the kinetics of forma- tion of oxide scales on iron. New experimental results will be introduced at several points during the discussion, and an attempt will be made to formulate a mechanistic picture consistent with these observations. The experimental techniques are described in an appendix. Iron-Oxygen System Fig. la shows the iron-oxygen phase diagram. Wüstite ("FeO"): The lowest of the oxides has a NaCl type cubic lattice formed by close-packing of the relatively large oxide ions with the smaller iron ions arranged in the interstices. The ionic radius for 0' is about 1.3 A, while Fe++ has a radius of about 0.8A and Fe"' about 0.7A. Its range of existence on the phase diagram is very wide for an ionic material of this type and does not seem to include the stoi-chiometric composition FeO. These deviations appear to be well established as vacancies on iron lattice sites by the comparison of lattice parameters and densities for a series of compositions by Jette and Foote.4 The limits of the field are disputed, but the work which seems most reliable" indicates vacancies in from 5 pct to more than 12 pct of the iron sites. In order for the crystal to be electrically neutral, there must be electron holes which correspond to the formation of trivalent iron ions equal to twice the number of vacancies. Magnetite ("Fe³O4"): Considerable disagreement also exists in the literature about the stable range of the Fe³O4 field. Some diagrams show the Fe³O4 composition to be lower in oxygen than the formula indicates, but Schenck6 gives a diagram from the

Jan 1, 1952

By H. D. Attra

The effects of phase behavior on reservoir economics is a function of both fluid composition and flow properties of the reservoir rock. In some operations such as gas injection into volatile crude-oil resertoirs, the stock-tank recovery derived from vaporization of the reservoir oil by the displacing gas can approach the recovery derived from the displacement mechanism. In other cases of gas-cycling reservoirs, the dry injected gas can effectively vaporize a retrograde liquid phase, allowing a substantial reduction in cycling pressure with no appreciable loss in recoverable liquids. Because of these effects, it is essential that phase behavior be considered in developing the economics of any secondary recovery program in which gas is the injected fluid. The inclusion of phase behavior in conventional calculation procedures has been retarded because of the complexity of the calculation and the large volume of laboratory data required. However, by applying high-speed computers and new laboratory techniques, the phase-behavior displacement calculation has been developed to economically include this basic fundamental of gas displacement. The calculation is performed by reducing the reservoir to a series of one-dimensional segments. Phase changes, saturation distributions and changing fluid properties are evaluated for each segment as a function of time. The relative volumes of equilibrium liquid and gas displaced from each segment are determined by a trial-and-error procedure to meet the requirements of phase equilibrium and two-phase flow. The produced fluids are flashed through conditions of surface separation approximating those in the field to obtain actual values for stock-tank liquid recovery and producing gas-oil ratios. Based on the results of successful field studies, the technique appears to have application to reservoirs producing under a nonmiscible gas-displacement mechanism and should substantially increase the reliability of this type of analysis when phase behavior is an important factor. Predicting reservoir performance under gas-injection operations involves three basic principles of reservoir engineering. First, a material balance exists between the displacing gas and the displaced reservoir fluid. Second, the oil and gas flow in relationship with the two-phase flow property of the reservoir rock; and third, phase equilibrium exists between the displacing gas and the native reservoir fluid. To illustrate these three principles, consider the one-dimensional reservoir shown in Fig. 1. The reservoir is divided into a series of segments representing a one-dimensional cross section between an injection area and a producing area. So represents the oil saturation, Cw represents the connate-water saturation and Sg represents the displacing gas saturation. The gas is moving into Segment 1 and the reservoir oil is being produced out of Segment N. Consider the case when gas has just moved into Segment 1 and the displacing gas-oil contact is at position a-a The volume of oil pushed out of Segment 1 is a function only of mobility and the two-phase flow properties of the reservoir rock, that is, Kg/Ko and µo/µg. Behind the front if the injected gas is foreign to the reservoir oil (a nonequilibrium gas), there will either be vaporization of the oil into the gas or condensation of the gas into the oil. In either case the net result is a mass transfer and a re-distribution of composition and liquid and gas saturation. If the displacement is at constant pressure and is continued, the gas moves into new segments, proceeding succesively into Segment 2, 3, etc. The total volume of oil produced from the reservoir

By T. S. Lundy, J. Askill

J. Askill and T. S. Lundy (Oak Ridge National Laboratory)—Moore has proposed a simple correlation between the activation energy for solute diffusion in fee metals and the atomic volume of the diffusing species using both tracer and chemical-diffusion data. The first point we wish to make is that it is wrong to attempt such correlations without separating the two types of data. Activation energies obtained using tracer techniques cannot, in general, be directly compared with those obtained from chemical-diffusion experiments. A knowledge of the variations of chemical potentials with composition is necessary and chemical-diffusion data must be corrected for such variations, even for dilute alloys, before a comparison is possible. Secondly, for nickel, iron, aluminum, and lead, Moore neglected a number of recently published works, e.g., Al26 and Mn54 in aluminum,46 cr51 in nickel,47 and Fe55 in iron.48 These and others should have been used. If one considers only the tracer diffusion data, there is left no simple correlation of the type proposed by Moore. For example, in the case of solute diffusion in aluminum the only reliable tracer data listed by Moore are for zinc and aluminum. The nickel, cobalt, and iron data of Agarwala49 are representative of near-surface diffusion-not bulk diffusion. In conclusion we suggest that for tracer diffusion data in these fee metals there is no justification for proposing any simple correlation of activation energy with atomic volume using the limited data that are presently available. R. H. Moore (author's reply)—The writer has proposed a simple correlation which is useful in estimating the internal consistency of diffusion data. The accuracy of the correlation is, of course, a function of the number of data available. Lundy and Askill take issue with the writer's decision to include data for the systems: nickel, iron, aluminum, and lead, taking the view that the data are insufficient to establish the correlation. The writer agrees that the data are insufficient and has cau- tioned the reader that the correlation is tentative in these instances. Inclusion of these systems is justified by the interesting trends the correlation reveals. For example, the distinction between interstitial and substitutional diffusion may be clearly established, Figs. 4 and 6, and the behavior of transition elements and nontransition elements is indicated to parallel the behavior illustrated by the silver and copper systems. Lundy and Askill have singled out the aluminum system for special criticism. The writer chose to discuss the data for diffusion of iron, nickel, and cobalt in aluminum in the terms used by the original authors9 and has included the comments of Balluffi.l0 The writer did not intend the reader to conclude that data for these elements should correlate with data for the nontransition elements as Lundy and Askill have evidently assumed. Iron, nickel, and cobalt form intermetallic compounds with aluminum, so the activity coefficients of these elements may be expected to deviate markedly from unity. This may account for the low frequency factors observed by Hirano.9 The nontransition elements, on the other hand, yield relatively simple phase diagrams with no evidence for intermetallic compound formation as in the case of iron, nickel, and cobalt. In dilute solution, Henry's law should be closely obeyed, the activity coefficients should not deviate markedly from unity, and the effect on the activation energy should not be large. The accuracy of the correlation applied to these systems will improve as the data is extended. It does not seem unreasonable to conclude that its extension should be encouraged.

Jan 1, 1965

By Klaus M. Zeitler

INTRODUCTION One cannot discuss the subject of "Raising Capital in the 1990s" without having a crystal ball and a reliable fortune teller. This discusses "Raising Capital for the 1990sW, which is what Metal1 Mining Corporation did in 1987. We raised (3166 million in an IPO (initial public offering) that was originally planned for only some (375 million. Why did we raise this money? How did we do it, and why was this IPO so successful? The answer to the first question requires a description of the promoter and principal shareholder of Metal1 Mining Corporation, Metallgesellschaft AG. This West Germany company was founded over 100 years ago in Frankfurt as a trading house for non-ferrous metals. It has developed into one of the world's leading firms in the non-ferrous metals industry. The company's worldwide activities in all areas of that business include mining, smelting and refining, the trading of ore, base and precious metals, the production of semi-fabricated and finished metal products, and the provision of engineering, construction and transportation services. Metallgesellschaft AG has also diversified into many other areas, the most significant of which are manufacturing, the production of chemicals, the development and application of technical and chemical processes, the planning and construction of industrial plants, and activities associated with finance, insurance, shipping and freight forwarding, and the provision of barter and countertrade services. The company maintains a network of more than 100 offices around the world; it employs some 25,000 people in its worldwide activities. [Selected financial data, as of September 30, 1987: (MetaUgesellschafl AG) Working Capital DM 8%.6 million CDNS 672.8 million Long Term Debt DM 568.0 million CDNS 426.2 million Sales DM 14.9 billion CDNS 11.2 billion Net Income DM 87.7 million CDN$ 65.8 million Market Capitalization at April 27, 1987: (Metallgesellschafi AG) Outstanding Shares ~,6oo,o0o DM 1,875 million CDNS 1,407 million] For the fiscal year ended September 30, 1987, Metallgesellschaft AG had, on a consolidated basis, net sales of DM13.3 billion, net income of DM100 million and total assets of about DM7.0 billion. Metallgesellschaft AG established the American Metal Company, a predecessor of AMAX Inc. one of the largest ventures undertaken by Metallgesellschaft AG then was the acquisition, through the American Metal Company, of a majority share in the Mexican Penoles mine, a large lead and silver deposit. METALLGESELLSCHAFT AG Figure 1 depicts an overview of Metallgesellschaft AG before and at the time the initial public offering of Metall Mining Corporation's shares took place. It shows that mining was far from the centre of Metallgesellschaft's activities. Therefore, it could not expect to receive an extraordinary amount of board attention. Thus, mining investments not only had to compare well with other mining opportunities, but were in fierce competition with projects from other divisions. In addition, mining was sub-divided into two separate departments -- domestic and international. These two departments had different management styles and diverging interests which resulted in a minimum co-operation between them. The domestic mining department was primarily interested in the survival of the local mines. These mines were characterized by rising operating costs and declining ore grades. It took a special commitment and continued ingenuity to develop and adjust to local conditions, and modern mining and processing technology, in order to remain competitive, or at least survive within the international mining industry. Local management

Jan 1, 1990

By Michael Tenenbaum, C. C. Brown

Numerous investigatiolls have been carried out to determine the total oxygen present in the basic open-hearth bath and the results of these studies have clearly defined the factors controlling the bath oxidation during the period prior to deoxidation. On the other hand, very little is known concerning any changes in total oxygen content that may occur during deoxidation, tapping and teeming. The need for such additional data is continually becoming more apparent, since there is no strict technical basis for selecting any particular deoxidation practice. Many steels are made for the same purposes and to the same chemical analyses by widely differing practices. Apparently, each practice is satisfactory. However, it would be desirable to establish more definitely the effects and relative merits of the individual practices. Such information could be used to give a definite foundation for evaluating various types of deoxidation. In order to approach this problem, a study was made of the changes effected in the total oxygen content of several grades of steel by representative open-hearth deoxidation practices. To carry out such studies, it was necessary that a reliable testing procedure for determining the oxygen content of the steel bath and the ingot mold be established. Before entering on any extensive testing program, consideration was given to the available methods of obtaining steel samples and analyzing for oxygen. Testing Procedures It has been shown5-9 that in solidified steel oxygen may ekist in combination with iron, manganese, silicon, aluminum, titanium, chromium and carbon, depending upon the amount of these elements in the sample. In molten steel, however, there has been no direct distinction of the various oxides. Accordingly, the total active oxygen in the molten steel bath is designated by the symbol 0. During deoxidation, various oxides may be present either in solution in the molten steel or as solid or liquid inclusions. Any sample taken from the molten steel for oxygen analysis will include both the dissolved and the mechanically included oxygen. Similarly, samples of the solid steel product will also include the two forms of Oxygen' Methods of Samplirzg the Molten Steel Bath In order to obtain a true oxygen analysis from liquid steel, this element must be completely retained in the solidified sample. Several methods have been proposed whereby this stabilization can be effected. As early as 1915, Kichlinelo proposed a method for testing molten steel, in which the oxygen combined with some added aluminum. The concentration of the resultant aluminum oxide was determined by chemical analysis. This method was

Jan 1, 1945

By J. B. Newkirk, A. H. Geisler

The alloy Cu4Pd has a disordered face-centered-cubic structure when quenched from temperatures between 478ºC and the melting point (about 1100°C). Below 478ºC an ordered phase is stable. The results of a Debye-Scherrer X-ray analysis indicate that the ordered phase has a tetragonal unit cell described by the space group C24h — P42/mt with 2 Cu in 2a, 2 Cu in 2f, 4 Cu in 4j (x = 0.2, y = 0.6), 4Pd in 4j (x = 0.4, y = 0.2), and 8 Cu in 8k (x = 0.1, y = 0.3). The orientation relationship between the face-centered-cubic phase and the ordered tetragonal phase is given by: [100],,. // [130]al,. COO1Ia.d.//COO1I,,.. • The behavior of Cu,Pd is typical of ordering alloys except that the transformation is very sluggish. The increase in hardness and the microstructural and X-ray diffraction effects are interpreted in terms of coherency strains caused by the ordering. AN anomalous construction in the Cu-Pd phase diagram (Fig. 1) was reported in 1939 and has been allowed to stand without further published attention since that time. The odd figuration about the composition 10 to 27 atomic pct Pd is derived mostly from the work of Jones and Sykes.1 Evidently several features of this binary system require further study if the constitutional forms are to be well understood. The present paper includes a study of one of these features, that is, the crystal structure of a single ordered alloy containing nominally 20 atomic pct Pd. This choice of composition was suggested by the work of Harker and associates who determined the structure of Ni4Mo2 and Ni4W.3 The nature of the ordering process in Cu4Pd was studied also by observing the hardness, microstructure, and Debye-Scherrer patterns of specimens which had been aged at various temperatures after quenching from an initial disordering treatment. Experimental Methods A 20 gram ingot of Cu4Pd was made by melting spectrographically standardized copper from Johnson, Matthey, and Co., and commercially pure (99.5 + ) palladium in an argon-filled quartz tube. Chemical analysis showed that the ingot contained 80.0 atomic pct Cu. The ingot was rolled about 60 pct to a strip 0.060 in. thick and was homogenized for 16 hr at 950°C in low pressure argon. Rods cut from the rolled strip were worked into wire 0.015 in. in diameter, and specimens for hardness and microscopic examination were cut from the remaining strip. All specimens, with the exception of some of the wire, were given an initial disordering treatment by heating for 16 hr at 950°C, followed by water quenching. A 10 cm length of as-drawn wire was water quenched after being held in a temperature-gradient furnace4 for 89 days. Room-temperature Debye-Scherrer photograms were then made at points along the wire to determine the temperature below which the ordered phase was stable. Although the accuracy of temperature determination in the gradient was only about ±10 °C, the temperature gradient was sufficiently gradual that the sensitivity was much better and locations which had differed by as little as 1°C could be distinguished. An analysis of the crystal structure of the well ordered alloy was made by X-ray diffraction using a specimen cut from this wire. The change of Debye-Scherrer pattern as ordering progressed was studied by using isothermally aged samples of initially disordered wires. The wires were sealed under low-pressure argon in small quartz tubes for heat treatment. After the aging treatment, the tubes were quenched in water and photograms were made at room temperature in a 10 cm diam camera using filtered Cu kX. (A = 1.540511) Hardness was measured on a Vickers hardness tester using a 10 kg load and 2/3 in. objective lens. Reported values are the average of at least three impressions made on flat specimens 0.060 in. thick. After the hardness of a heat-treated sample had been measured, it was resealed in low-pressure argon and returned to the furnace for continued aging at the same temperature. In this way, two samples served for all aging times at each temperature. Hardness specimens which had been aged 500 hr or more were used for metallographic examination after the final aging treatment. A dilute potassium-dichromate etching solution was used.

Jan 1, 1955

By Zay Jeffries

IT has been known for a number of years that certain aluminum alloys could be hardened by quenching from a temperature of about 500° C. Immediately after quenching the total increase in hardness is not marked, but the hardness increases gradually so that after about 4 days, it has reached a maximum. The actual increase in hardness probably continues for weeks or even months, but in order to detect the differences after 4 days, it is necessary to make tests at rather long intervals. So far as is known to the authors, Alfred Wilm was the first to discover these properties in aluminum alloys.1 The aluminum-copper alloys are amenable to this treatment and the addition of magnesium makes the heat-treatment effect more pronounced. Until recently heat treatment of aluminum alloys has been confined to the worked alloys in the form of sheet or extruded products. The alloy that has been most used in this connection is known commercially as duralumin and has about the following composition: copper not exceeding, 6 per cent.; magnesium not exceeding, 2 per cent.; manganese not exceeding, 1 per cent.; commercial aluminum,2 balance. A typically good alloy in this range consists of: copper, 4 per cent.; magnesium, 0.5 per cent.; manganese, 1 per cent.; commercial aluminum, balance. This alloy can be rolled into sheets or extruded; and in this condition it can be heated in the ordinary furnace atmosphere to a temperature near 500° C. and quenched in water. After about 4 days it will have the following physical properties: tensile strength 55,000 to 60,000 lb. per sq. in. (3866 to 4218 kg. per sq. cm.); elongation in 2 in., 15 to 20 per cent. No harmful effects of the furnace gases or even the air, if the heating is affected in an electric furnace in an air atmosphere,

Jan 9, 1919

By W. A. Gibson, Zay Jeffries

It has been known for a number of years that certain aluminum alloys could be hardened by quenching from a temperature of about 500' C. Immediately after quenching the total increase in hardness is not marked, but the hardness increases gradually so that after about 4 days, it has reached a maximum. The actual increase in hardness probably continues for weeks or even months, but in order to detect the differences after 4 days, it is necessary to make tests at rather long intervals. So far as is known to the authors, Alfred Wilm was the first to discover these properties in aluminum alloys.' The aluminum-copper alloys are amenable to this treatment and the addition of magnesium makes the heat-treatment effect more pronounced. Until recently heat treatment of aluminum alloys has been confined to the worked alloys in the form of sheet or extruded products. The alloy that has been most used in this connection is known commercially as duralumin and has about the following composition: copper not exceeding, 6 per cent.; magnesium not exceeding, 2 per cent.; manganese not exceeding, 1 per cent.; commercial aluminum,2 balance. A typically good alloy in this range consists of: copper, 4 per cent.; magnesium, 0.5 per cent.; manganese, 1 per cent.; commercial aluminum, balance. This alloy can be rolled into sheets or extruded; and in this condition it can be heated in the ordinary furnace atmosphere to a temperature near 500° C. and quenched in water. After about 4 days it will have the following physical properties: tensile strength 55,000 to 60,000 1b. per sq. in. (3866 to 4218 kg. per sq. cm.); elongation in 2 in., 15 to 20 per cent. No harmful effects of the furnace gases or even the air, if the heating is effected in an electric furnace in an air atmospherc.

Jan 1, 1921

By Edward E. Hicke, Charles A. Stickels

The dihedral angle of liquid-lead inclusions in solid nickel has been measured as a function of temperature from 371 to 816 C at zero pressure. and as a function of pressure up to 50,000 psi at 317 and 593 C. Thermodynamic expressions are derived for the temperature and pressure coefficients of interfacial tensions in this system. A method is presented for obtaining relative tempmature and pressure coefficients from dihedral-angle data, and the resulting values are used in conjunction with the thermodynamic relations to calculate excess properties of the solid-liquid interface. Defining the position of the solid-liquid interface by the condition that the excess concentration of lead is zero, the excess concentration of nickel and the excess entropy at the solid-liquid interface are: where y(593) is the interfacial tension of the nickel grain boundary saturated with lead at 593C and zero pressure. These values are considered reliable to within a factor of two. While similar calculations are not possible for the nickel grain boundary, it is shown that the expwimental results imply that the excess concentration of nickel at the boundary and the excess entropy of the boundary are negative. THE temperature and composition dependencies of interfacial tensions in metal systems have been reported by several authors in the last 15 years.'-' No measurements of the effect of pressure on inter- facial tensions in metals have been reported, although measurements made on nonmetallic gas/liquid and liquid/liquid systems suggest that a measurable effect should exist.6-l2 The purpose of these experiments was to determine the influence of temperature and hydrostatic pressure on dihedral angles in the Pb-Ni system. The equilibrium dihedral angle is expressed as a ratio of interfacial tensions by2: where 8 is the dihedral angle, yGB is the Ni-Ni grain boundary tension, and SL is the solid Ni-liquid Pb interfacial tension. Three sets of ex- periments were run to determine: a) the effect of temperature from 371" to 816°C on the dihedral angle at zero pressure; b) the effect of pressure up to 50,000 psi on the dihedral angle at 371° and 593°C. THERMODYNAMICS The properties of interfaces of most direct concern are the interfacial tension, excess concentrations of various components at the interface and the excess entropy of the interface, S°. These quantities are defined following Gibbs. Consider an interface separating phases a and 0. Construct two surfaces C, and C such that all properties of phase a remain uniform up to C, and all properties of phase 0 remain uniform up to C. The region between C, and Cp is the interfacial region. Let a "dividing surface", D, be constructed in the interfacial region such that it passes through all points "similarly situated with respect to the condition of the adjacent matter."l3 Surface D divides the interfacial region into a portion a' adjacent to phase adjacent to phase . Let a closed curve, C, enclose an area A in D. Erect normal lines to D along C; these lines and the surfaces C,, Cp, and D then enclose volumes V and V' in the two halves of the portion of the

Jan 1, 1964

By Charles C. Russell, Glenn C. South

A primary factor in the selection of coals for making coke at high temperatures is the amount of pressure the coal will exert upon the oven walls when carbonized in modem by-product ovens.l-3 This factor has become increasingly important in recent years because of demands for increased throughput and higher coke yields. As a result of the increasing necessity for assessing coals in regard to their carbonization pressures, more and more research work is being directed toward developing labora-tory-scale tests for predicting carbonization pressures and for showing how blends of coking coals may be modified to eliminate pressures liable to damage coke-oven walls. The specific reasons have never been defined as to why some blends of coking coals develop dangerous carbonization pressures and others do not, therefore specific procedures to eliminate dangerous pressures could not be prescribed. It is true that the major processes occurring during carbonization have been described by many writers,4 and there is a general agreement on the ideas of the changes that occur as coal is transformed into coke. However, there is little evidence relating the processes occurring during carbonization to the pressure developed and the available evidence is very incomplete and often contradictory. This deficiency arises in part from the fact that the fundamental causes of pressure development have not been defined, at least from experimental data, and in turn this is the result of the lack of satisfactory testing equipment for studying the problem. Operating experience has shown several ways of reducing or eliminating dangerous pressures,5 but probably none of these ways is applicable to all blends of coking coals, and such empirical methods often lead to production of inferior coke. They are trial-and-error methods, not designed to remedy specific causes of pressure development. The number of laboratory-scale testing devices designed to estimate the carbonization pressures that result when coals are carbonized in commercial ovens is probably well over a hundred.6 Few of these instruments have remained in use for any length of time, either because the predictions were not sufficiently accurate or, more often, because too many exceptions were found. Often testing devices are operated without any attempt at duplication of conditions in commercial ovens and erroneous and contradictory results are frequently obtained. Therefore, it has seemed desirable to study the properties of coal itself and to determine any relationships between the properties of the coal and the carbonization pressures developed. Relationship between Properties of Coal and Carbonization Pressure Attempts to determine relationships between various properties of coal and carbonization pressures, for the most parti have been indirect because no satisfactory and reliable laboratory testing device has

Jan 1, 1944

By F. E. Harris

It has been said that carbon is "ubiquitous" with reference to iron alloys. Certainly at temperatures where carbon and iron form the solid solution, austenite, it may be readily added to, or removed from, steel surfaces. Differences in concentration potential are the sole requirements for carbon flow. To measure the rate of carbon flow, a proportionality factor, D, is arbitrarily expressed in units of (length)2 X (time)-'. If the term used for carbon concentration is weight Pet C, as is customary, then a factor must be found, changing this term to One of weight Per unit volume when the diffusion coeficient, D, is used in flow formulae. The determination of D values for carbon is complicated since D varies with the concentration itself. The Matano method has often been employed for systems involving a variant D. This method expresses all evaluation of unsteady flow whereby D is given a relationship with the numerical value of C-x slopes. The author contends that this solution is incorrect, except when D is invariant. The primary purpose of this paper is to support this contention. Here diffusion is controlled under two conditions of flow. The first part deals with the steady state; the second part considers unsteady flow under definite restrictions. Beginning with the fundamental concept, the development is continued until equations for C-x curves are obtained. The solution is then subjected to experimental data. It is realized that a great amount of reliable data is required for a rigorous proof. KO such research is claimed. Instead an attempt is made to set forth the physical significance of the diffusion mechanism, and particularly to express the factors which influence the flow of carbon in austenite. Of equal significance is the implication that gaseous media may be employed with reasonable accuracy in making these determinations. The Steady State Diffusion of Carbon in Au~tenite The fundamental experiment defining diffusion in solid metals may be readily performed for carbon in gamma iron.' For this experiment a thin plate of medium carbon steel is chosen. The surfaces of this plate may be taken as two parallel planes which may be considered infinite for points well in the Center of the planes. These two planes are held at different concentrations of carbon at a constant temperature. This temperature, 1700°F for example, is one where the single phase: carbon in solution with gamma iron, exists. When these conditions are held for a sufficient time the concentration of carbon of the different points of the solid ment down towards its steady state value. At points well removed from the ends, the concentration will remain the Same along planes parallel to the surfaces of the plate.

Jan 1, 1948