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By J. R. Low, M. Gensamer

During the course of an investigation into the drawability of automobile-body sheet steel, it became apparent that certain advantages would be possessed by a deep-drawing steel with a very low yield strength, but almost normal tensile strength, provided a reduction in yield strength could be accomplished without sacrifice of ductility and fine grain size. It would also be desirable to have this steel free from strain-aging; that is, free from the tendency for the yield point to return after it had been eliminated by straining. It was known that, wet hydrogen treatment acted to eliminate aging, but as carried out by other investigators, at relatively high temperatures, this treatment resulted in grain coarsening. It has been established by some of the experiments reported in this paper that all of these objectives can be accomplished by treatment in wet hydrogen at a low temperature, in the neighborhood of 72o°C. The time of treatment is relatively short (about 3 hr, to completion with 20-gauge sheet) if the water content of the hydrogen atmosphere is made fairly high, say aimost 30 per cent by volume. At this temperature. the grain size remains fine in the time required. After this treatment sheet reduced the usual amount by cold-rolling has a grain size of about A.S.T.M. No. 7. The yield point is completely eliminated, and the yield strength is reduced to a very low value, in the neighborhood of 15,000 lb. per sq. in. The stress-strain curve departs gradually from the straight-line elastic part after the fashion of nonferrous deep-drawing alloys. This elimination of the yield point is accompanied by a verv low indentation hardness; for example, 32 Rockwell B or even lower. There is no aging as measured by increase in the yield strength upon aging at room temperature or for 3 hr. at 200°C., after stretching about 10 per cent in tension. Tensile strength is only slightly less than the box-annealed sheet, seldom dropping much below 40,-ooo lb. per sq. in. Elongation is practically unchanged, or increased slightly. The term yield point, as here used, refers to the phenomenon of heterogeneous deformation occurring at a substantiallv constant load, observed principally in annealed low-carbon steels. If such a steel is loaded in tension, the load increases steadily with elastic strain, suddenly falls, fluctuates slightly about some constant value for a time and then begins to rise again as deformation proceeds. The maximum initial load before the sudden drop is generally called the "upper yield point"; the value of the load where deformation proceeds at a

Jan 1, 1944

By J. R. Low, M. Gensamer

During the course of an investigation into the drawability of automobile-body sheet steel, it became apparent that certain advantages would be possessed by a deep-drawing steel with a very low yield strength, but almost normal tensile strength, provided a reduction in yield strength could be accomplished without sacrifice of ductility and fine grain size. It would also be desirable to have this steel free from strain-aging; that is, free from the tendency for the yield point to return after it had been eliminated by straining. It was known that, wet hydrogen treatment acted to eliminate aging, but as carried out by other investigators, at relatively high temperatures, this treatment resulted in grain coarsening. It has been established by some of the experiments reported in this paper that all of these objectives can be accomplished by treatment in wet hydrogen at a low temperature, in the neighborhood of 72o°C. The time of treatment is relatively short (about 3 hr, to completion with 20-gauge sheet) if the water content of the hydrogen atmosphere is made fairly high, say aimost 30 per cent by volume. At this temperature. the grain size remains fine in the time required. After this treatment sheet reduced the usual amount by cold-rolling has a grain size of about A.S.T.M. No. 7. The yield point is completely eliminated, and the yield strength is reduced to a very low value, in the neighborhood of 15,000 lb. per sq. in. The stress-strain curve departs gradually from the straight-line elastic part after the fashion of nonferrous deep-drawing alloys. This elimination of the yield point is accompanied by a verv low indentation hardness; for example, 32 Rockwell B or even lower. There is no aging as measured by increase in the yield strength upon aging at room temperature or for 3 hr. at 200°C., after stretching about 10 per cent in tension. Tensile strength is only slightly less than the box-annealed sheet, seldom dropping much below 40,-ooo lb. per sq. in. Elongation is practically unchanged, or increased slightly. The term yield point, as here used, refers to the phenomenon of heterogeneous deformation occurring at a substantiallv constant load, observed principally in annealed low-carbon steels. If such a steel is loaded in tension, the load increases steadily with elastic strain, suddenly falls, fluctuates slightly about some constant value for a time and then begins to rise again as deformation proceeds. The maximum initial load before the sudden drop is generally called the "upper yield point"; the value of the load where deformation proceeds at a

Jan 1, 1944

By G. Nabi, W-K. Lu

Cylindrical specimens of natural dense hematite were reduced to magnetite at atmospheric pressure in H2-H2O mixtures of known composition over the temperature range 1084° to 1284°K. The rate of reduction was measured by the rate of movement of the interface between hematite and magnetite. The diffusion of gases through the gaseous boundary layer, the magnetite layer, and the interfacial chemical reaction were all considered in the interpretation of experimental data. The mass transfer coefficient through the boundary layer was calculated using accepted correlations. Values of the chemical reaction rate constant and the diffusivity of hydrogen in the magnetite phase were determined. THE present investigation is concerned with the reduction kinetics of natural hematite to magnetite by H2-H2O mixtures in the temperature range 1084" to 1284°K at atmospheric pressure. This reaction is the first step in the series of topochemical reactions in the process of reducing hematite to iron. Kinetic information of the simple steps such as hematite-magnetite transformation is necessary in order to have a better understanding of the complex processes of hematite reduction in iron-making. It also has direct industrial significance because magnetic roasting is one of the most important methods in benefication of lean ore.&apos; Although many technical papers have been published on the process of magnetic roasting and iron oxide reduction, very little information is available in the literature concerning the fundamental nature of hematite reduction to magnetite by reducing gases. Hansen et al.2 reduced the dense synthetic pellets of high-purity oxide in CO-CO2 mixtures and determined the reaction rate by weight-loss method. They were able to interpret most of their results by applying the interfacial area control theory developed by Mckewan.3 In contrast, Wilhelm and St. Pierre,4 who studied reduction of hematite to magnetite in H2-H2O mixtures by weight-loss method, stressed that the resistance of the porous magnetite layer to the diffusion of gases cannot be neglected in consideration of the overall reaction rate. In the present study the contributions of interfacial chemical reaction, diffusion of gases through the magnetite phase, and the gaseous boundary layer to the overall reaction rate will be considered. APPARATUS AND PROCEDURE Hematite Specimens Preparation. Natural hematite ore from Vermillon range of Northern Minnesota was selected for the present investigation because of its high purity and thermal stability. Chemical analysis of five samples gave the following average values: 67.52 pct total iron (96.62 pct Fe2O3, 0.28 pct FeO, 0.03 pct metallic iron), 2.53 pct SiO2, <0.07 pct MgO, 0.03 pct CaO, 0.05 pct combined mixture, 0.07 pct loss on ignition, and 0.34 pct other. Cylindrical specimens of 0.93 cm in diam and 2.7 cm in length were drilled from slabs of ore with a water-cooled diamond core drill. These specimens were heated to 1000°C and furnace-cooled. Specimens with silica pockets developed large cracks. The uncracked specimens were heated a second time, and their surfaces were carefully examined with a microscope. Those with hairline cracks or surface inhomoaenitv-- were rejected. Preparation of H2-H2O Mixtures. H2-H2O mixtures were prepared by the combustion of H2-O2, mixtures in a pyrex glass chamber in the presence of a catalyst. Alumina pellets coated with palladium, supplied by Englehard Industries, were used as the catalyst. Purified grades of hydrogen and oxygen were used which were repurified by usual techniques. Hydrogen before entering the combustion chamber was passed through an activated alumina H2O absorption bulb, with copper turning at the top. The cover of this bulb was not made pressure-tight so that any pressure development in the hydrogen line would cause the cover to blow off and also the copper turnings would act as a flame arrester in the case of a flashback from the combustion flame. Oxygen flow rates were measured with a bubble flow meter after purification with 1 pct accuracy. Hydrogen flow rates were measured by "precision wet test meter" and the amount of unburnt hydrogen was accurately measured by a bubble flow meter, after condensing water vapor in the gaseous stream. The Pyrex glass bulb contained concentric Vycor glass tubes as shown in Fig. 1. Oxygen was prevented from diffusing into the hydrogen line by threading platinum wire through pores at the combustion end of gas inlet tube. The glass bulb was heated with a Kanthal heating wire pasted in asbestos paper. The surface temperature of the bulb was measured with a thermocouple and adjusted to remain at approximately 350°C. The gaseous reaction chamber also served as a preheater for gases to avoid thermal segregation. The following sequence of operation was adopted. 1) Nitrogen was passed through the outer concentric tube to purge the catalyst bulb of oxygen. 2) Hydrogen was introduced through the inner tube until a steady flow was obtained. 3) Oxygen was then introduced into the nitrogen stream passing through the outer tube. 4) When combustion had commenced and a flame was visible over the platinum wire, the N2 was turned off.

Jan 1, 1969

By G. Derge, Ling Yang, G. M. Pound

The specific conductance and its temperature dependence were measured over the entire composition range of the molten Cu2S-CuCI system. At a typical temperature of 1200°C, 10 rnol pet of the ionically conducting CuCl reduced the specific conductance from about 77 ohm-lcm-l for pure Cu2S to about 32 ohm -1cm -1, and 50 mol pet CuCl reduced the conductance to that for pure CuCI—about 5 ohm 1cm1. The nature of electrical conduction in molten Cu2S, FeS, CuCI, and mixtures was studied by measuring the current efficiency of electrolysis at about 1100°C. The Cu2S, FeS, and mattes were found to conduct exclusively by electrons, but addition of 1 5 wt pet CUS to Cu2S produces a small amount of electrolysis. Addition of CuCl to Cu2S suppresses electronic conduction, and ionic conduction reaches almost 100 pet at a CuCl concentration of about 50 mol pet. These facts are interpreted in terms of electron energy level diagrams by analogy to the situation in solids. RESULTS of electrical conductivity studies on molten Cu-FeS mattes as a function of composition and temperature have been reported.&apos; The specific conductances ranged from about 100 ohm-&apos; cm-&apos; for pure Cu2S to 1500 ohm-&apos; cm-1 for pure FeS. This is in sharp contrast with the low specific conductance of molten ionic salts for which the transfer of electricity is by migration of ions in the field. In general, these ionically conducting molten salts, such as NaC1, KC1, CuC1, etc., have a specific conductance of the order of magnitude of 5 ohm-&apos; cm-&apos;. It was concluded on the basis of this evidence that molten FeS and Cu,S exhibit electronic conduction. Pure molten FeS has a small negative temperature coefficient of specific conductance, resembling metallic conduction, while pure molten Cu2S has a small positive temperature coefficient, resembling semi-conduction. The molten Cu2S-FeS mattes follow a roughly additive rule of mixtures, both with respect to specific conductance and temperature coefficient. Savelsberg2 has studied the electrolysis of molten Cu2S and Cu2S + FeS. He concluded that while molten Cu2S is an electronic conductor, there is some ionic conduction in molten Cu2S + FeS3 owing to the formation of the molecular compound 2Cu2S.FeS and its dissociation into Cu1 and FeS2-1 ions. The present work does not verify his results. Chipman, Inouye, and Tomlinson" have studied the specific conductance of molten FeO and report a high specific conductance, about 200 ohm-1 cm-1 of the same order of magnitude as that found for molten mattes, and a positive temperature coefficient. They interpret these results in terms of p-type semiconduction in the ionic liquid by analogy to the situation in solid FeO.1 imnad and Derne&apos; detected appreciable ionization in molten FeO by means of electrolytic cell efficiency measurements. In order to verify the conclusion that electrical conduction in molten Cu2S and mattes is electronic, and to gain further insight into the structure of molten sulfides, the following investigations were carried out in the present work: 1) The specific conductance, s of the molten system Cu2S-CuC1 was measured as a function of temperature over the entire composition range. As discussed later, molten CuCl is an ionic substance. It was thought that if molten Cu2S were simply ionic in nature, addition of small amounts of CuCl might not have a catastrophic effect in lowering the high conductance of the Cu2S. On the other hand, if much electronic conduction occurs, addition of the ionic CuCl should have a large effect in destroying the electronic conduction. 2) The electrolytic cell efficiency of the following molten systems was measured at about 1100°C in specially designed cells: Cu3; Cu2S + FeS, 50:50 by wt; FeS; Cu2S + CuS, 15 wt pet; Cu2S + CuC1, 5.9 to 46.4 mol pet; and CuC1. This gives a direct measure of the fraction of current carried by ions in these melts. Further, the cell efficiency, extrapolated to zero ionic current, is given by cell efficiency = (s leasile + s elexstronic). [1] s lucile for molten CulS would be expected to be no greater than that for molten CuC1, whose s lonle is about 5 ohm-&apos; cm-1, as will be seen in the following. u,.,,.,.,.......for molten Cu,S is of the order of 100 ohm-&apos; cm-&apos;.&apos; Thus, a large increase in cell efficiency from 0 to values of 10 to 100 pet upon addition of CuCl to Cu2S would indicate destruction of the electronic conductance. Conductance Measurements Experimental Procedure—The apparatus and experimental method were the same as those described in detail in connection with the study of electrical conduction in molten Cu,S-FeS mattes.&apos; A four terminal conductivity cell and an ac poten-

Jan 1, 1957

By W. M. Gray

The determination of the state of stress existing at a point in any solid body requires the determination of the six components of the stress tensor relative to a convenient set of axes. One general method of doing this in rock is to make strain or deformation measurements associated with local relief of stress in a number of boreholes penetrating a limited volume in the vicinity of the point at which the stress is to be determined. The stress is assumed to be constant throughout this volume. Leeman1 has described a number of the instruments that have been developed for making deformation or strain measurements in boreholes. Panek2 has described the application of statistical methods to the determination of the average components of ground stress from sets of measurements made in boreholes by means of deformation meters of the type developed by the U.S. Bureau of Mines. In this paper attention is given initially to the strain cell developed for use on the flattened ends of boreholes by the Council for Scientific and Industrial Research (CSIR) , South Africa. The equations required for evaluating the average components of ground stress from measurements made by means of the CSIR strain cell are derived. These equations are similar in form to those applying to measurements made by the borehole deformation meter developed by the U.S. Bureau of Mines. The results of an investigation of the properties of the equations are described. These results are important in the practical use of both the instruments mentioned above and any instrument based on similar principles.

Jan 1, 1968

By S. F. Emmons

The first paper which appears in the published Transactions of our Institute is that read by our respected Secretary at its first meeting in Wilkes-Barre in May, 1871. It is entitled " The Geographical Distribution of Mining Districts in the United States," and presents a brief but masterly review of what was known of the distribution of our deposits of useful minerals, particularly the metals, not only from a geographical but from a geological stand-point. At the request of Dr. Raymond I agreed, somewhat hastily, perhaps, to write for this occasion a brief sketch of the geological distribution of the deposits of the useful metals in this country, in the light of the increased knowledge of the present day. In the time given no personal investigation was possible, and as it was therefore out of the question to attempt to make anything that could be considered an original contribution to the history of our ore-deposits, I have been obliged to limit myself to an examination of such pub-

Jan 1, 1894

By Fritz Wald, Martin Weinstein, Herbert E. Bates

ONE of the methods for preparing a fiber-reinforced composite is that of directional solidification of a eutectic melt. Its major drawback lies in the fact that the matrix in this case is a pure constituent (or a terminal solid solution of fixed concentration) and the ratio of fiber to matrix volume is fixed. Only very few, if any, pure metal matrices possess technically useful properties, and the necessity for alloying is the inevitable consequence of this fact. Nevertheless, structures similar to binary eutectic ones can be obtained in higher-order systems by pseudoeutectic reactions of a type where crystallites and a solid-solution matrix are simultaneously precipitated from a liquid.1"3 Trials to solidify such alloys directionally—the matrix consisting of a solid solution of chromium, nickel, and silicon and the reinforcing phase being Cr3C2— will be reported elsewhere.4 This note describes constitutional studies carried out primarily to determine -the position of the pseudoeutectic and further to insure that no undesirable reactions between the reinforcing phase and components of the matrix occur in the solid state at elevated temperatures. In the investigation it was specifically desired to reinforce a matrix which is as close as possible to the 80/20 Nichrome composition with Cr3C2 to insure good oxidation resistance and improve the strength at high temperatures. No konstitutional information on the quaternary C-Cr-Ni-Si svstem in the range of interest was found in the literature. Nevertheless, the two most important ternaries c-cr-Ni5 and have been completely investigated. On C-Cr-Si some investigations have been conducted,7"9 and C-Ni-Si was studied with a few alloys in the course of the present investination 9 Eight alloys were prepared to investigate the changes in the constitution of the ternary C-Cr3Ce-yNi eutectic5 with silicon contents up to 10 wt pct. Five alloys were made with varying chromium and carbon contents around compositions containing -60 to 70 wt pct Ni and 5 to 6 wt pct Si. The latter alloys were designed to locate the y(Ni-Cr-Si-C)-Cr3C2 pseudoeutectic reaction and to determine the maximum chromium content in solid solution compatible with a two-phase equilibrium between y and Cr3C2 only. The alloys were prepared from electrolyte nickel, semiconductor-grade silicon, high-purity chromium (99.999 pct pellets), and semiconductor-grade graphite. They were melted inductively in high-purity alumina crucibles under flowing argon. Twenty-gram charges were prepared and investigated by optical microscopy (including hot-stage investigations), X-ray powder methods, and electronmicroprobe analysis. The alloys were studied in the as-cast condition and also after a homogenization annealing treatment of 48 days at 900°C in evacuated quartz ampoules. The structures of all alloys based on the ternary eutectic plus silicon in the as-cast state consisted generally of an inhomogeneous solid solution and at least two other phases. Annealing of the alloys showed that the solid solution was homogenized but that an incomplete solid-state reaction had taken place in all alloys with silicon contents over 5 wt pct. This reaction is clearly visible in Fig. 1 around the large Cr3C2 crystals (gray areas). Below this concentration, the alloys consisted of a homogeneous matrix and precipitates of Cr3C2 plus graphite, Fig. 2. The Cr3C2 was entirely unreacted here. From these observations and subsequent X-ray and microprobe investigations, it was deduced that the

Jan 1, 1969

By K. P. Gupta

In ternary systems containing silicon and transition elements, silicon was found to stabilize the a phase.1 At 1000°C a broad a-phase region was found in the Cr-Mn-Si system,1 which extends very close to the Mn-Si binary line (within about 3 at. pct Cr). Later investigation showed that, if chromium was replaced by small amounts (3 to 5 at. pct) of v,1 Ti, Nb, Ta, Mo, or w,1 the a phase or the R phase was again stabilized. Whether stabilization of complex crystal structures is possible by small additions of iron, cobalt, or nickel to Mn-Si is not known. Investigation of the Co-Mn-Si system4 shows the presence of several ternary phases, but at cobalt concentrations higher than 10 at. pct. In the present work the ternary systems Fe-Mn-Si and Co-Mn-Si were investigated at low iron and cobalt concentrations (< 7 at. pct Fe or Co) in search of phases of complex crystal structure. Experimental Procedure. The alloys were pre-pared by induction melting Co (99.8 pct), Si (99.9 pct), electrolytic Fe (99.88 pct), and Mn (99.9 pct) in recrystallized alumina crucibles under an inert atmosphere of argon gas. The alloy specimens were wrapped in molybdenum foil, sealed in argon-filled (at 1/2 atm) silica capsules and annealed for 72 hr at 1000" 1°C, and water-quenched. The use of argon-filled capsules was found to reduce the losses during annealing and the molybdenum foil prevented reaction of the manganese with silica. The total weight loss due to annealing and melting was found to be less than 1 pct for each alloy. The alloys were not analyzed chemically. A part of each annealed specimen was polished and examined metal-lographically using etching solutions, a) 0.1 to 0.5 pct oxalic acid solution in water, or b) 0.5 pct HNOs in methyl alcohol. Even though the etching solutions were quite dilute, the specimens were heavily etched and stained if the etching time exceeded more than 5 sec. Etching times found suitable were usually 1 to 3 sec. For X-ray diffraction work, a part of each annealed specimen was powdered by crushing after the exposed surface was ground off. X-ray diffraction patterns were taken with powder specimens using a 20-cm-diam asymmetrical focusing camera and Fe K radiation. These powder patterns were found quite satisfactory for the purpose of phase identification. However, X-ray diffraction patterns of single-phase alloy powder specimens were also taken using a Siemens diffractometer and filtered Cr-Ka radiation for greater resolution of the diffraction lines. Results and Discussion. In the Fe-Mn-Si system a new ternary phase was found at 1000°C between the p manganese and the MnsSi phase regions. The approximate phase relationships in this region of the 1000°C isothermal section are shown in Fig. 1. The new ternary phase, designated as the K phase, was found to occupy a narrow region about 1 at. pct in width and extending from Fee.oMn74.5Silg.5 to Fe.Mn.Silg.. At the high-iron content region, the K phase was found in equilibrium with the p manganese and the MnsSi phases but, at the low-iron content end of the K-phase region, the phases in equilibrium with the K phase were the p manganese phase, the MnsSi phase, and another new phase, designated in Fig. 1 as the v phase. The v

Jan 1, 1964

By Frank Woolley, John F. Elliott

The high-temperatzrre solulion calorimeter has been modified and an extension dynamic analysis of i/s transient behavior has provided an improved basis for interpretation of the experimental data. The partial molar heats of mixing, Him, of aluminum, copper, and silicon in the corresponding binary liquid-iron alloys haz,e been measured directly at 1600°C. The resulls in kcal per g-alom are: THE high-temperature1,2 solution calorimeter was planned and designed to obtain directly the partial molar heat of solution of solute elements in a variety of liquid metals that might serve as solvents. Results obtained in the temperature range of 1000° to 1200°C were described in the two earlier publications. However, a long-term goal in the program has been to obtain measurements at 1600°C of the heats of solution of elements in liquid iron. Hence, subsequent to the work last reported,2 and based on the experience of the first two studies, the calorimeter system has been modified and operated at 1600°C to obtain measurements on the Fe-Si, Fe-A1, and Fe-Cu systems. This paper reports the results of the work. THE MODIFIED SYSTEM It was clear from the earlier results that the most precise measurements can be obtained with the high-temperature solution calorimeter when the solution time is less than approximately 30 sec. This can be achieved only when the bath temperature, Tb, is above the melting (liquidus) temperature of the solute material being dropped. A second point of importance is that the refractories used earlier liberated considerable water so that, although the oxygen potential of the atmosphere in the system was satisfactory for work with nickel and copper, it was too high for work with iron, silicon, and so forth. Third, for work at 1550" to 1650°C, an improved temperature-sensing unit is needed. The modified calorimeter in its latest form is shown schematically in Fig. 1. The refractory brick and the Kanthal-Super heating elements which formed the furnace for heating the solute addition have been replaced by a smaller resistance-heated tube furnace which is insulated with bubble-grain alumina. A more elaborate gas purification system was installed to maintain a satisfactorily low oxygen potential in the atmosphere of the calorinieter chamber and of the solute addition furnace. As was the case in the earlier work, at the beginning of a run the calorimeter chamber is evacuated and then filled with purified argon so that a total pressure slightly in excess of 1 atm is maintained during the run. When either of the furnaces within the enclosure attains a red heat, 6 pct H2 is added to the atmosphere. Throughout the run, gas from the calorimeter is circulated through a de-oxidation system at the rate of approximately 5 standard liters per min; the changeover time for the gas within the calorimeter enclosure is approximately 1 hr. The essential stages of the deoxidation system are: a copper furnace at 600°C, a drying tube of anhydrite, and a titanium furnace at 800°C. The oxygen potential of the gas stream in the purification system can be monitored at several points by a stabilized zirconia oxygen cell. This cell is invaluable in detecting the development of leaks in the container vessel and the gas circulating system. The installation of the solute-addition furnace allowed either of two types of temperature-sensing units to be mounted on a water-cooled probe. The first was a small optical device using a silicon diode that was

Jan 1, 1968

By R. E. Reed-Hill, W. H. Hartt

The second order, {1011}-{1012} twin bands observed in critically deformed magnesium are often accompanied by an unusual form of kinking. These kinks, which lie adjacent to and run more or less parallel to the twin band habit, allow plastic deformation that occurs inside the twins to be accommodated into the matrix. The bend planes of these kinks cannot be explained using slip dislocations with Burgers vectors lying in the basal plane. A mechanism involving a single set of {1012)(1011) dislocations is proposed to account for the kink band habit, in good agreement with the experimental observations. It has been shown that second order, {101l)-{1012} twinning is an important phenomenon for magnesium deformed in the ambient temperature range. A detailed study of these twins has recently led to a better understanding of their irrational habit plane1 and their relation to fracture in magnesium.&apos; This paper is concerned with a third aspect of these twins; this being the kink bands, originally referred to as shear accommodation kinks,3&apos;4 that are often induced in the adjacent matrix. A twin band and its shear accommodation kink usually appear as a single macroscopic deformation, as seen in Fig. 1. The complex nature of these deformed regions is easily revealed, however, with the aid of a suitable microscope. It has been proposed374 that the primary function of these kinks is to accommodate the nonhomogeneous deformation that normally occurs in the second order twins. The twins tend to form as small discrete lamellae lined up approximately one in front of the other.&apos;&apos;&apos; Very large plastic deformations varying in magnitude from one twin to the next, may occur inside these lamellae. The kinks serve both to pass this strain into the matrix and to accommodate strain variations between twins. Deformation within the kinks, subsequent to their formation, has also been observed;3,4 and it is thought that this contributes to the strain accommodation process. The bend planes of the kinks have a habit near that of the twin band itself, which lies about 56 deg from the matrix basal plane. This orientation does not bisect the angle formed between the basal plane to either side of the bend plane and therefore differs markedly from the kink structure usually observed in hcp metals.&apos;-&apos; The present paper reports the results of a study of shear accommodation kinks and proposes a possible solution to the problem of the orientation of the bend plane. EXPERIMENTAL PROCEDURE Rectangular cross-section tensile specimens with axes parallel to (1010) and faces to (1210) and (0001) were cut and acid machined from a high-purity magnesium single crystal. After straining on an Instron machine the specimens were examined by both optical and replica electron microscopy. The basic procedures have been previously reported."&apos; EXPERIMENTAL RESULTS AND DISCUSSION Structure of Shear Accommodation Kinks. A typical, large (101l)-{1012) twin band with its associated shear accommodation kink is shown running diagonally from upper left to lower right in Fig. 1. In this optical micrograph the twin band, consisting of a row of small, highly deformed twin lamellae, is visible as a narrow dark band along the left hand edge of the structure in question. The kink lies to the right of the twins, is approximately three times as wide, and is lighter in contrast. The horizontal slip lines visible on the general specimen surface correspond to matrix basal slip and are parallel to the tensile axis. Note that on the right-hand side of the band these slip lines can be followed from the matrix into the kink. At the kink band boundary the slip lines suffer a clockwise rotation of about 25 deg. Rotations of this magnitude are not uncommon. When a surface such as that in Fig. 1 is given an etch1 making it sensitive to polarized light, it has been observed that the slip traces in the kink are parallel to the basal plane trace in this region. This is in excellent agreement with the assumption that these regions are indeed kinks. Fig. 2 is an electron micrograph of this same twin band and kink. Note that a row of small supplimation pits has formed

Jan 1, 1969

By D. R. Wieland, A. N. Barron, A. R. Hendrickson

A definite relationship has been found between the reactivity of flowing hydrochloric acid and its shear rate in a carbonate fracture. Both flow velocity and fracture width affect the acid reaction rate. Laboratory studies were conducted on acid reactivity at different flow velocities through horizontal-linear fractures, using 15 per cent hydrochloric acid at 80&apos; F and approximately 1,100 psi. Fracture width varied from 0.02 to 0.20 in. These data provide a new basis from which the spending time and penetration of the acid can be estimated. Equations were derived expressing the relationship between injection rate, fracture width, acid concentration, time and fracture height, for linear and radial fracture systems. Because the penetration of the acid before spending is closely related to the extent of productivity increase resulting from an acidizing treatment, these data provide a valuable insight into some of the controlling factors that must be taken into consideration during treatment preplanning. INTRODUCTION Acidizing of carbonate reservoirs to improve production characteristics has been widely practiced since 1932. Originally, it was assumed that the acid uniformly penetrated natural formation pores and flow channels, enlarging them and thereby increasing their flow capacity. Little consideration was given to the reaction rate of the acid, or how far it would penetrate away from the wellbore into the formation, before spending. It has been shown recently&apos; that, unless fractures are present in the rock, very little penetration is attained before spending, and the benefits of the acidizing treatment are largely confined to the immediate vicinity of the well-bore. Therefore, acid treatments may be classified into two categories: (1) matrix acidizing, in which the acid flows through multiple: formation pores; and (2) fracture acidizing, in which the bulk of the acid travels through fractures in the rock, whether natural or induced. When acidizing treatments are conducted at pressures of sufficient magnitude to open and extend such fractures, it is often desirable to inject a propping agent to hold the fracture open after the treating pressure has been released, thus providing a highly conductive flow channel through the rock. In some cases where acid attack produces surface irregularities, the resulting flow passages can be sufficient to provide high conductivity without use of a propping agent. During injection, the acid dissolves the carbonate rock with which it comes in contact until it is spent. Deeper penetration of the spent acid into the formation produces no appreciable benefit (if no propping agent is used) because the unetched fracture faces will rejoin when treating pressures are released and the fracture will "heal", with negligible resultant conductivity. The spending time of the acid thus becomes important in determining how far from the wellbore the improved-conductivity zone extends. The spending time of acid after injection into carbonate rock depends on the rate at which the acid reacts with the rock. This in turn is controlled by a number of factors, as previously reported. 2-4 These include temperature, pressure, acid concentration, rock composition, injection rate and the area-volume relationship between the acid solution volume and the surface area of the formation flow channels through which it penetrates. Thus, in matrix acidizing, where an extremely large area is exposed to the acid, spending is rapid. In contrast, spending time is prolonged in open fractures, where the area-volume ratio is much lower. Thus, greatest penetration of the acid before spending will be achieved in acid-fracturing treatments where fractures are held open by hydraulic pressure. Perkins and Kern5 have presented equations for fracture-width determinations which, for hard formations such as limestone or dolomite, indicate that high injection rates and viscous fluids are required to create wide fractures. Increasing the injection rate will, in itself, produce faster reaction rates. Because this type of reaction can be considered to be first-order diffusion-controlled, &apos;-&apos; the rate of movement of fluid past the rock surface affects the thickness of the diffusion layer, in turn affecting the reaction rate. Staudt, et al,10 report the results of tests in which a rotating marble cylinder was immersed in acid, and the weight loss at different speeds determined. Such test conditions, however, are not truly indicative of the situation in a formation during an acid-fracturing treatment. This paper reports reaction-rate data obtained under

By I. S. Servi, N. F. Graves

THE etch-pit technique has long been used to reveal low-angle boundaries and, in general, the distribution of dislocations in high-purity metals. Often this technique is amenable to quantitative computations,&apos; but generally is affected by some degree of uncertainty because etch pits can be nucleated not only by dislocations, but also by inclusions and surface imperfections. In the case of high-purity zinc, Gilman&apos; reported that etch pits were produced only after adding 0.1 atomic per cent of cadmium to the crystals and aging at a proper temperature. In the course of a study of low-angle boundary behavior in zinc crystals, etch pits were produced in 99.999 pct zinc without resorting to intentional contamination, by etching in a 50:50 HNO3H2O solution. The technique consists of etching until the heat of reaction brings the solution to a boil, and in displacing the etchant rapidly with a stream of cold water, thus avoiding staining. The undesirable features of this technique are the heating effect and the removal of a considerable amount of material. A high polish is obtained if staining is controlled properly; therefore, the etch pits appear very clearly and with excellent contrast when observed under dark field illumination. This is especially helpful for cursory examinations at low magnifications. Several observations were made on zinc crystals of square cross section, having one face parallel to the basal plane and one edge parallel to a slip direction, Fig. l(a). When such a crystal was bent about 2" over a glass die and annealed at 390°C, a sharp boundary was formed only in the inner portion of the crystal. Near the two prismatic faces, perpendicular to the basal plane, the boundary was always branched into many lower-angle boundaries, Fig. 1(b). Attempts to move the main boundary by applying a shear stress to the specimen invariably resulted in curving the boundary, the branched ends remaining pinned in their original location, Fig. 1 (c). The prismatic face, when etched as indicated in the foregoing, revealed the existence of numerous (sometimes up to 30) fine boundaries, nearly parallel to each other. Pre-existing lineage boundaries, generally parallel to the direction of growth, were barriers to the formation of polygonization boundaries and interrupted the continuity of the substructure, Fig. 2. Twin

Jan 1, 1959

By Charles T. Holland

IN connection with the problems of bumps in coal mines, much has been written concerning the manner in which roof action and methods of mining enter [ ] into the pressure effects observed but little has been written upon the relationship between the physical properties of coal and associated rocks and the cause of bumps. The question of the strength of bituminous coal, of its elasticity, and the energy it can absorb and store has been neglected to a considerable extent in connection with this problem. STRENGTH OF COAL It is generally agreed that one of the necessary conditions for bumps is a strong [ ] coal.1 Therefore, since the strength of coal is an important factor, a study of this property should be made. Several investigators have conducted laboratory experiments upon the strength of coal, using small cubical or rectangular test pieces. An attempt will not be made to go into all of these investigations, but they are briefly summarized in Table I. Experiments of the U. S. Bureau of Mines8 and of Lawall and Holland (pp. 27 and 34 of ref. 9), however, indicate that regardless of whether the coal is a soft and weak semibituminous (Pocahontas No. 3) or a hard and strong bituminous (Stockton) coal, the maximum ultimate unit stress developed in laboratory tests is higher in small cubes than in larger cubes of the same coal.

Jan 1, 1942

By E. Amott

A test is described in which the wellubility of porous rock is measured as a function of the displacement properties of the rock-water-oil system. Four displacemet operations are carried out: (I) sponlaneous displaceti?ent of water by oil, (2) forced displacement of water by oil oil in the same system using a centrifuging procetllrre, (3) spontaneous displacement of oil by water. and (4) forced displacment of oil by water. Ratios of the spontaneous displacement volumes to the total displucenlent volumes are used as wettability indicates. Cores having clean mineral surfaces (strongly preferentially water-wet) show displacement-by-waler ratios approaching 1.00 and displacement-by-oil ratios of zero. Cores which ([re. strongly preferentinlly oil-wet give the reverse resu1ts. Neutral wellability cores show zero values for both ratios Fresh cores from different oil reservoirs have shown wettabiltties in tlris te.st covering rrlti~ost 111e conlplrtt, range of thr: te.st. Notvever, nlo.s/ of tlle fresh California cores tested were slightly prcfere111icrlly wclter-\vet. The chrrnge.~ in coro u (&apos;liabilities, as indicated hy this te.st, r~.sril/ing from various CO~P hanrlling procedures tt,ere oh.served. In sonie ca.sc,s /Ire ~,cttahilitio.c. of fresh cores were changell by drxi:~g or 11y e.x/rclct ing with iolcreiii~ or. dioxunc~; in o/h~r cases they were 1101 changed. Co~ltrrc/ of cort,.s ~.ith filtrc~t~c. from water-base rlrilling rilrrrls crlrc.sed littlc change in we/ /ahility ivhile contnct with filtrates frorii oil-hus~ ri1rlcl.s tlecrrascrl the prefcrerlcc, of the, cores for )I.NI Usitig thi,s test to ri.crl~lute n~r~ttubili~y, N .vt~ldy was iilarle of /lie correlmtio~i of wettability with wa/erfloocl nil recovery for orttcrop Ohio sand.stone and for Al~ln-tlunl. Resul/.v indicate thml no single correlatioti between these factors applies to different porous rock syste~n. It is thought that diflerences in pore gen~netry resrrlt in diflrrerrce.~ in this correlurio~z. INTRODUCTTON Most investigators who have reported on the wettahility of porous rock have described such rock as prcferentially water-wet or preferentially oil-wet. In some cases a third classification, neutral wettability. has been used. The efficiency of water floods in each of these wettability groups has been described in numerous publications. Several methods for characterizing porous rock wet tability more precisely have been reported,&apos; " but it appears that because of one weakness or another. none of these has been generally accepted. Early in our studies in this field, it was found that the displacement efficiency of oil by water in a particular porous rock having a strong preference for water was quite different from that in a similar rock having only a moderate preference for water. Thus, there appeared to be a need for a practical, reasonably precise wet tability test. one which could classify porous rocks into 10 to 20 different groups rather than the two or three broad groups listed above. The test developed to meet this need is described in this paper. Also, changes in wettability, as indicated hy this test, resulting from various core handling procedures are discussed. Finally, data showing the corrclation of wettability with waterflood oil recovery for two different types of cores are presented and discussed. Some confusion has resulted from the failure of certain writers to define clearly some of the wettability terms they have used. Accordingly, the following commcnts concerning definitions are offered. The wc t ta-hility of a solid surface is the relative preference of that surface to be covered by one of the fluids under consideration. It is felt that this is the generally accepted definition. The fluids being considered must bc specified (or understood) before the term wettability has any significance. In the work reported here these fluids are water (3 per cent brine) and oil (kerosene). The term preferential wettability is sometimes used, but we think that the word preferential is redundant here and should not be used. Tn line with the definitions of Jennings&apos;, a preferentially oil-wet solid surface is regarded as a surface which will show an oil advancing contact angle less than 90" (measured through the oil) in the water-oil-solid system. Oil will spontaneously displace water, if both are at the same pressure, from such a surface. A preferentially water-wet surface is analogous. This is consistent with the wettability definition above. As Jennings has said, frequently the term oil-wet is used to mean the same thing as preferentially oil-wet. However, oil-wet also has been used occasionally referring to an oil-covered surface when the availability of water was limited. To avoid confusion from this source, we do not use the terms oil-wet and water-wet. DESCRIPTION OF WETTABTLITY TEST The following points were considered desirable in a wettability test for our purpose. 1. The test should be a displacement test resembling

By John C. Bilello

The brittle fracture behavior of cold-worked sintered tungsten was studied over the temperature range 4.2° to 298°K using a high-sensitivity strain measuring system and electronfractography. Similar observations were made on a swaged electron beam zone-refined monocrystal. In sintered tungsten irreversible plastic deformation was observed during cyclic load-unload tests at stress levels well below the fracture stress for all temperatures, but general microyielding could be detected only down to 202°K. For the zone-refined samples macroyielding occurred at all test temperatures with evidence for twinning below -202°K. The fracture stress of the sintered tmgsten was virtually independent of temperature, while the zone-refined crystal showed a 2.3 times increase over the same temperature range. Electronfractography confirmed the presence of numerous rod-shaped and spherical submicroscopic voids which ranged in diameter from 1400 to 4300A in the sintered tungsten; no voids could be found in the zone-refined tungsten. Contrast effects observed on the replicas in the vicinity of certain voids indicated that plastic deformation could be induced by the local stress concentration. It has been suggested that the presence of these voids may be responsible for the low-temperature brittle failure of sintered tungsten. Based m this suggestim und on the evidence obtained here, a dislocatim model is presented to account for the brittle behavior of sintered tungsten. In this model slip, which is induced by the local high stress concentration in the region at the edge of a favorably oriented void, could cause the void to grow to a microcrack of critical size. STUDIES of brittle fracture in bcc metals have led to the well-known experimental relationships between grain size, yield stress, fracture stress, and temperature which have formed the basis for the various dislocation pile-up1-3 or interaction4&apos;= models for slip-induced microcrack nucleation. While microcracks can be nucleated by deformation twins,6,7 there has been no direct evidence furnished by transmission electron microscopy to support conclusively either the Zener pile-up or Cottrell dislocation reaction models for producing micro-cracks in all "brittle" materials. In addition to the "inverse" grain size relationship for yield and fracture stresses the cottrel14 theory predicts that the fracture stress below the transition temperature should behave in a fashion similar to that of the yield stress above this temperature. Such behavior has been verified for several bcc metals.8-10 With reference to both grain size effects and the tem- perature dependence of the fracture stress below the transition temperature, the behavior of sintered tungsten appears anomalous. Early work by Bechtold and Shewrnon 11 showed no apparent temperature dependence of the fracture stress below the ductile-brittle transition temperature (DBTT). They attributed this result to the intergranular nature of the fractures observed. More recent work by Wronski and Four-deux12&apos;13 on considerably purer material did not show any systematic relationship between the fracture stress and temperature below DBTT. The dependence of flow and fracture stresses on grain size is also not clearly established for sintered tungsten. Koo, for example, has shown that the DBTT for sintered tungsten depended chiefly on the annealing temperature and was relatively insensitive to the actual grain size achieved. Using electrofractography and transmission electron microscopy, Wronski and Fourdeuxl3 showed that numerous spherical and rod-shaped submicroscopic voids could be found in sintered tungsten but not in melted tungsten of nominally the same purity. They suggested that these voids could be responsible for the temperature insensitivity of the fracture stress below the DBTT. In the present work the temperature dependence of the fracture stress for high-purity commercially sintered tungsten has been determined. The presence of submicroscopic voids in sintered materials was confirmed, and these were studied in detail to examine the role they could play in nucleating brittle fracture. A dislocation model is suggested which could cause an inherent spherical void to lengthen into a Griffith crack of critical size. EXPERIMENTAL PROCEDURE Commercially sintered tungsten rod was obtained in the as-swaged condition from Sylvania. A zone-refined crystal was obtained from the same source. This crystal was grown by giving three zone passes (at 25.4 cm per hr) to a sintered rod of high-purity tungsten. The rod axis prior to cold working was -15 deg from the [110] direction. Originally the zone-refined rod was -6 mm in diam; it was reduced to -3 mm by eight swaging passes, at high temperatures, with each step having about the same reduction of area. The final swaging step gave a 7.5 pct reduction of area at 1050°C. All swaging operations were performed in a hydrogen atmosphere. For the sintered rod a similar working schedule was employed. Metal-lographic examination of the sintered material revealed that the cold-worked structure had an apparent grain diameter of -25 u transverse to the swaging direction (obtained by the intercept method). In the longitudinal direction cold-worked grains were approximately 1.5 to 2 times their diameter. No distinct fiber structure could be observed optically for the zone-refined rod. The cold-worked structure in the

Jan 1, 1969

By Yuri D. Tretyakov, Robert A. Rapp

The stoichiometry ranges ofNiOl+y and LiFe,O,-d were established by high-temperatwe electrochemical meas7rements in a stabilized-zirconia electrolyte cell. The results were consistent with doubly ionized cation vacancies in NiO,+y and interstitial lithium or iron ions in . The defect structure of the ternary ferrite was derived from the consideration of equilibration with respect to oxygen between the solid and the gas phase. The absolute magnitudes of defect concentrations were calculated. Pavtial molar enthalpies of oxygen in the compounds were calculated and interpreted in terms of the enthalpy of defect formation in these crystals. NICKEL oxide (NiO,,?) is a metal-deficient, p-type, extrinsic semiconductor whose properties are consistent with a structural model based on the presence of cation vacancies as the predominant ionic defect at sufficiently high temperatures and oxygen activities. A survey of previously reported conductivity studies and the presentation of some more recent conductivity measurements will be given in a later paper.&apos; The absolute magnitude of the equilibrium vacancy concentration in NiO has been reported from combined conductivity and thermogravimetric data of &apos; However, disagreement exists concerning the state of ionization of the nickel vacancies in NiO. Some authors3-&apos; have proposed that the predominant defects in NiO are singly ionized nickel vacancies and positive holes (h&apos;) formed by the reaction where, accorhng to the notation of Kroger and Vink,6&apos;7 Oq represents an oxygen ion on its normal lattice site. Other authors"27E have proposed that doubly ionized nickel vacancies (V{i ) and positive holes are predominant and are formed by the reaction One purpose of the present investigation was to establish the nature of the predominant defect in NiO,+? as well as its equilibrium concentration and thermo-dynamic properties at elevated temperatures and known oxygen activities. To accomplish this purpose, the coulometric titration of oxygen into and out of NiO was accomplished using a galvanic cell involving the calcia-stabilized zirconia electrolyte. The ternary oxide, lithium ferrite (LiFe50,-6) is ferromagnetic and has the inverse spinel structure LiFe,O,-d Thus, in the ideal stoichiometric XY204 lattice, Fe3 ions occupy one-eighth of the tetrahedral (A) sites, and i&apos; and Fe3&apos; ions share at random one-half of the octahedral (B) sites within the fcc sublattice of oxygen ions.g The structure and thermodynamics of the spinel structures have been comprehensively decribed.&apos;-&apos; A recent compilation of literature for lithium ferrite is also available. 15 Oxygen-excess LiFe50sis not expected to exist because both i&apos; and 17e3& ions are in their highest normal valency states (positive hole formation is not favorable). As will be discussed, equilibration of the ferrite crystal with oxygen of a surrounding gaseous phase will result in the introduction of equilibrium concentrations of ionic and electronic defects. In the Results and Discussion section a defect model for LiFe50,-d is proposed. This defect model is tested by high-temperature coulometric titration experiments. EXPERIMENT The oxygen activities in nonstoichiometric NiOl+, (and also LiFe,O,-d) were measured by means of coulometric titration with the high-temperature galvanic cell The critical characteristic of a suitable experimental cell is the complete isolation of the phase to be investigated in a minimum sized chamber which is free from extraneous sources and sinks for oxygen (leakage). Then oxygen is only admitted to the chamber or removed from it in known amounts by coulometric titration, which involves the passage of oxygen ions through the solid electrolyte with electrochemical oxidation and reduction reactions at the platinum contacts to the electrolyte. The experimental cell is shown in Fig. 1. The cap of the cell was the tip from a closed-end alumina tube, which was found to be leak-free from a helium leak-detector test. This alumina cap was about 1.2 cm OD and about 1.5 cm high. A Zircoa calcia-stabilized zirconia tablet (crucible lid) of 1.5 cm diam and 0.3 cm thickness served as the solid electrolyte. A Pyrex ring of about 0.1 cm thickness was placed between the electrolyte and the cap. The electrolyte tablet was painted with platinum paste on the entire outer face and on that part of the inner face which would be within the enclosure; these electrodes were further prepared by heating in air at 1000°C for 2 hr. Nickel oxide powder, listed as 99.999 pct pure, was purchased from Leico Industries, Inc. Pills of the NiO were cold-pressed and sintered at 1050°C for 3 hr in a platinum crucible. This poorly sintered NiO was crushed, and chunks were wrapped in 52-mesh Pt gauze (to catalyze the solid-gas exchange) for place-

Jan 1, 1970

By A. M. Tweedy, Felix A. Vogel

Since the publication of Prof. J. W. Richards&apos;s paper on The Schumacher Briquetting Process,&apos; this process has been in operation on a practical scale in two plants in the United States, and a few notes on the results obtained may be of interest. An experimental plant, consisting of a toggle press and a hydraulic press, was installed at Johnstown, Pa., together with the necessary bins and other equipment, and an attempt was made to briquette a mixture of Mahoning fines and blast-furnace flue dust, in a ratio of 7 parts of ore to 3 parts of dust. The results were unsatisfactory, in that the briquettes made were not hard enough to stand the severe handling to which they were subjected in the various transferences necessary between the press and the furnace. The trouble was due to several causes: First, the mechanical arrangement of the plant, which did not lend itself to a proper hardening and loading of the briquettes into the cars. Second, the irrrgular composition of the flue dust, as to both its lime and its fine coke contents. Third, the variability of the moisture content of the ore, and its viscosity when wet. Fourth, the ratio of fine ore to flue dust; that is, it was attempted to bind in too large a quantity of fine ore for the small amount of flue dust used. The first trouble could be overcome, and has been improved upon in the plant as it stands to-day. The second trouble was important, because of the small amount of dust used as a binder, and might have been easily removed by mixing the flue dust of the different furnaces, which, however, was never done. The third difficulty was more serious, but could have been partly solved, had it been possible to meet this along with the fourth difficulty, by decreasing the percentage of ore in the briquettes. It was proposed to do this, but it seems that the problem at the works in question is, primarily, to prepare Mahoning fines for reduction in the blast furnace, and only secondarily to use up the flue dust, and a solution along this line did not seem to be interesting.

Jan 1, 1914

By L. A. Rapoport

This paper is intended as an aid in the perfornzance and interpretation of experimental studies of multi-phase flow in porous tnedia. The mathenmatical formulation of incompressible, two-phase flow phenomena in a three-dimensional porous medium is presented. The resulting equations which account for the effects of gravitational and capillary forces are used to derive general scaling laws permitting accurate representation of water-oil displacement processes by means of experimental flow models. The significance of the scaling laws is examined in the light of the physical concepts of micro-behavior of fluids in porous media. Formulations of scaling con(litions obtained by application of dimensional analysis are reviewed. The principles of application and the limitations of flow model studies are discussed. INTRODUCTION Most of the problems facing a reservoir engineer involve three-dimensional, two-phase flow systems. influenced by the effects of gravitational and capillary forces. Present computational procedures generally do not permit the evaluation of such systems without introducing greatly simplifying assumptions. The actual reservoir configuration is usually replaced by one or a combination of several "equivalent" linear systems, and the effects of capillarity and gravity are considered to be negligible. While these assumptions may in many instances be justified, there are a number of important practical problems for which the results obtained on the basis of such simplifications remain open to question. An alternate approach to the problems that are not amenable to analytical treatment consists of seeking their solution by experimental laboratory investigations. The main difficulty of such investigations lies in the fact that multiphase flow phenomena subjected to gravity and capillary forces proceed differently at different rates and in systems of different dimensions. Therefore, unless great discrimination is exercised in selecting the conditions and interpreting the results of laboratory tests, such tests may be entirely misleading in regard to field applications. The purpose of this paper is to present a theoretical background for the performance of laboratory tests that are more nearly representative of field behavior. It will be shown that simulation of field behavior can be achieved by conducting laboratory tests under conditions selected in accordance with definite scaling laws. The formulation of conditions required for the scaling of two-phase flow experiments in porous media has been presented in the literature on several occasions.12,3,4 One of these formulations advanced by Leverett1, cannot be considered applicable to most practical cases inasmuch as it does not cover generally the regime of laminar flow. An adequate treatment, based on the consideration of dimensional analysis, was presented by Engelberts and Klinkenberg.2 Their treatment, however, is concerned primarily with the scaling of linear flow systems, and hence somewhat limited in scope. Accordingly, the formulation of more general scaling laws, applicable to three-dimensional systems, appeared to be warranted. It appeared, furthermore. desirable to establish this formulation on the basis of a mathematical treatment, rather than by application of general principles of dimensional analysis, to permit more explicit

Jan 1, 1956

By P. S. Kotval, H. Hahell

PRECIPITATION of M23C6 carbides at grain boundaries in austenitic steels and nickel base superalloys has received considerable attention. Hatwell and Berghezan,1 in a study using carbon extraction replicas, observed that dendritic growth of the M23C6 carbide phase occurred at grain boundaries in austenitic steels. Merrick and Nicholson2 have observed that in thin foils of "Nimonic 75" alloy the precipitation of M23C6 carbide was discontinuous at grain boundaries. Lewis and Hattersley3 have suggested that the type of M23C6 precipitation observed in Ni-Cr austenites represented the early stages of discontinuous precipitation. In a more recent study, Singhal and Martin4 have concluded that grain boundary M23C6 precipitation involves the migration of an austenite grain boundary by a mechanism very similar to that proposed by Tu and Turnbull.5 Tu and Turnbull&apos;s suggestion essentially involves grain boundary migration under the driving force arising from the difference between the surface free energies of the opposite sides of the plate. Thus, Singhal and Martin&apos;s observation4 implies that the first stage of discontinuous precipitation has taken place and that the grain boundary has provided a diffusion short circuit for the sub-stitutional solute atoms (principally chromium) which assist the preferential growth of the M23C6 plate in contact with the grain boundary. However, no evidence of the characteristic cellular morphology typical of discontinuous precipitation was found in Singhal and Martin&apos;s studies. Presumably, this is due to the low volume fraction (i.e., low supersaturation of solute) of the precipitating carbide reported by Singhal and Martin.4 Recently, Raymond6 has studied precipitation of the M23C6 carbide in Incoloy alloy 825 and shown that the growth of this phase is perpendicular to grain boundaries at 1400°F and in the plane of the grain boundary at 1700°F. The former aging condition resulting in discontinuous or "cellular" precipitation. Raymond has suggested an explanation for this behavior based on the relative rates of chromium diffusion by volume and grain boundary diffusion at the two different aging temperatures. In the course of a recent investigation on precipitation behavior in several &apos; strengthened superalloys, precipitation of M23C6 carbide at grain boundaries was studied. Under conditions of a high degree of supersaturation, the M23C6 carbide precipitates in the discontinuous or cellular morphology in Waspalloy (Ni-18.51 Cr-2.0 Fe-3.53 Mo-1.16 Al-2.78 Ti-0.09 C). This note presents electron microscopy observations of this mode of precipitation and its effect on mechanical properties. Specimens of Waspalloy after "solution" heat treatment at 1950°F show essentially single phase matrix with a few residual particles of "primary" MC carbide. In normal commercial heat treatments, precipitation of Ni3(Al,Ti) ordered fcc r&apos; phase occurs throughout the matrix and M23C6 precipitates as globular particles at the grain boundaries. In these commercial heat treatments the degree of supersaturation during precipitation is low. However, if specimens of Waspalloy are quenched from the solution heat treatment temperature of 1950°F and then aged at an aging temperature where the degree of supersaturation is high, discontinuous precipitation of M23C6 occurs. Figs. 1 and 2 show thin foil micrographs of Waspalloy, solution heat treated at 1950°F, water quenched, and aged at 1292° and 1598°F, respectively, for 500 Fig. 1—Thin foil of Waspalloy, solution heat treated at 1950°F for 1 hr, water quenched, and aged for 500 hr at 1292°F. Magnification 15,500 times. Fig. 2—Thin foil of Waspalloy, solution heat treated at 1950°F for 1 hr, water quenched, and aged for 114 hr at 1598° F. Magnification 13,000 times.

Jan 1, 1970

By J. Chipman, J. C. Fulton

Reduction of Si from slag to carbon-saturated iron is a very slow reaction. The rate is nearly independent of stirring but is accelerated markedly by increased temperature. In a slag containing 45 pct SiO2, 38 pct CaO, 17 pct Al2O3, the energy of activation is approximately 130 kcal. AT temperatures of 1450" to 1600°C occurring in the hearth of the iron blast furnace, chemical reactions are normally expected to proceed with great rapidity. Many of the reactions of iron making are indeed quite rapid and attain a condition corresponding roughly to a state of equilibrium before metal and slag are tapped from the furnace. The one conspicuous exception is the reaction by which carbon in the coke or metal reacts with silica in the slag to deliver silicon to the metal. From the practical viewpoint the slowness of this reaction is indeed fortunate, for if it were fast enough to reach equilibrium during the time of contact, it would be difficult to produce pig iron of the desired low silicon content. The one other reaction which under certain circumstances can also be very slow is the transfer of sulfur from blast-furnace metal to slag. In their laboratory exeriments on this reaction, Hatch and Chipman found that at least 5 hr were required for the attainment of equilibrium in a small crucible and that the reaction was not accelerated by increasing the rate of stirring. Subsequently it was shown2 that the presence of a reducible oxide MnO greatly inhibited sulfur removal, and finally that SiO2 itself3 under some conditions has sufficient oxidizing Dower to interfere with sulfur transfer. The kinetics of sulfur removal are now rather well understood,4 and it has been shown3,5 that when the silicon content of the metal is high enough to prevent reduction of SiO,, the desulfurizing process is rapid. Thus the observations of slow removal of sulfur are associated with the coincident slow reduction of SiO2, a process which transfers both silicon and oxygen to the metal phase. It is, of course, the oxygen which interferes with sulfur removal. An explanation of the slowness of the silica reduction reaction has not been obtained. The sulfur transfer process has been shown by the authors5 to be at least partially transport-controlled in the absence of interference by silica. Under the experimental conditions. 99 pct of the total transfer at 1600" was accomplished within an hour. Under similar conditions the reduction of silica may not proceed half-way to equilibrium6 in 8 hr. It is scarcely imaginable that silica reduction could be transport-controlled when the slag contains 44 pct SiO2 and the metal about 1.4 pct C. The experiments reported here show definitely that it is controlled by a slow step or steps in the chemical reaction. One might expect to find such a slow step in the breaking of SiO bonds or in the slow removal of oxygen by carbon at the slag-metal or metal-crucible interface or by bubble formation. EXPERIMENTAL PROCEDURE The experimental procedure used in this investigation was similar to that used in the study of the kinetics of desulfurization.5 A sketch of the graphite crucible and stirrer used in the current work is shown in Fig. 1. The quantity of metal and slag used was adjusted so that the slag-metal interface was located between the top and bottom blades on the stirrer; the top of the slag layer was above the top blade. The stirrer shaft was hollow down to the metal level and served as thermocouple well or optical pyrometer target. Stirring speeds of 0 to 500 rpm were used. A carbon monoxide atmosphere was maintained in the furnace by flushing the gas over the slag surface, except in run No. 220 when it was bubbled

Jan 1, 1960