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By C. Zener

HE present paper has its origin in an attempt A by the author, extending over the last several years, to understand the influence of the magnetic properties of the constituent atoms upon the various properties of metals and alloys. During this study the author has been impressed by the important role which magnetism plays in the many facets of metallurgy. In writing this paper he has attempted to convey to his colleagues some concept as to the role magnetism will play in the future development of metallurgy. In some of the subjects herein discussed, such as the influence of magnetism upon phase boundaries and upon mechanical properties, only general thermodynamical properties associated with ferromagnetism are introduced. For these subjects, it is not anticipated that any serious differences of opinion will be encountered. A discussion of the other subjects requires the adoption of a quite specific model of the magnetic structure of the metal. This model is discussed under "Magnetic Shells and Their Interactions." The simple model adopted in this paper has been criticized' as being mathematically inconsistent. This criticism has been ably silenced by Dr. W. J. Carr' of this laboratory. The simple model adopted in this paper has been criticized as being inconsistent with observations on magnetic structure using neutron diffraction. This apparent discrepancy has been resolved by Bader3 who has shown that the finite velocity of neutrons prevents neutron beams from detecting the anti-ferromagnetic structure of the type predicted by the simple model adopted in this paper. Magnetic Shells and Their Interactions An electronic shell of an atom is specified by two symbols.' The first symbol is a number and denotes the total quantum number of the shell. The second symbol is a letter and denotes the orbital angular momentum of each electron within the shell. The letters s, p, d, f are used to denote an orbital angular momentum of 0, 1, 2, 3, respectively, in units of h/2a, where h is Planck's constant. As we pass from scandium to nickel in the first transition period, we eraduallv fill the 3d shell. This shell is responsible for ferromagnetism and hence may be called the magnetic shell. The component of the orbital angular momentum along some specified direction must be an integral multiple of h/2~. Thus, for a d electron this component can have any one of the values —2, —1, 0, 1, 2 times h/2a. In addition to its orbital angular momentum, each electron has an intrinsic angular momentum, hereafter called spin, due to a rotation about its own axis. This spin can have either of two components + Yz (h/2a) or — Yz (h/2~) along the specified direction. A fundamental principle of atomic structure, known as Pauli's exclusion principle, is that no two electrons in a given shell may have identical components of orbital angular momentum simultaneous with identical components of spin. The 3d shell thus may have at most ten electrons. Since an intrinsic angular momentum of a charged particle is necessarily associated with a magnetic moment, the spin of an electron may be thought of as referring to either this angular momentum or to the associated magnetic moment. The electron configuration of the ground state of the isolated atoms in the first transition row is reproduced in Table I. The superscript represents the number of electrons in each shell. For the purpose of this paper, the most important principle of atomic structure relates to the manner in which electrons are added to a shell as the atomic number is increased. This principle, commonly known as the principle of highest multiplicity or more simply as Hund's rule, states that within a given shell the maximum number of electrons have spins pointing in the same direction as is consistent with Pauli's exclusion principle. An application of this principle to the d shell is illustrated in Fig. 1, where the spin structure is given for shells containing from 1 to 10 electrons. The physical basis of Hund's rule is the exchange energy discussed first by Heisenberg in 1928. This is a coupling between electrons with parallel spins (spins pointing in the same direction), the coupling

Jan 1, 1956

By R. L. McCarron, G. R. Belton

The diffusivities of oxygen and sulfur in liquid iron have heen determined hy a capillary technique in which the surface concentrations of the solutes were established by means of appropriate H2/H2 and H2S/H, gas mixtures. Total diffusate and concentration profile results are shown to be in good accord, yielding for- 1560 and Supporting results at 1660°C are also presented. The conditions necessary to avoid gas transport control in this type of experiment are discussed. IN spite of their importance in understanding the kinetics and mechanisms of refining reactions, the dif-fusivities of oxygen and sulfur in liquid iron are not well established. Accordingly, as a first step in studies of rates of solute absorption from the gas phase into liquid iron, new measurements of these diffusivities have been made and are presented in this paper. The only published results for the diffusion of sulfur in pure liquid iron are those of Kawai.' He used a diffusion couple technique in which two cylindrical specimens, one containing sulfur and the other with negligible sulfur concentration, were joined together and held in a refractory capillary. After an experiment, the sample was quenched and the concentration distribution of the solute determined. Kawai recognized that significant changes in solute distribution occurred during melting and freezing and he attempted to correct the concentration profiles for these effects to give a sulfur diffusivity of 4.6 x 10-6 sq cm per sec at 1560°C. The method of correction, however, was not rigorous and the uncertainty in this result cannot be easily assessed. Koslov et a1.2 have reported the diffusivity of oxygen in iron as 7.8 x 10"3 sq cm per sec at 1660°C. This value appears to be unreasonably high but, unfortunately, details of their experiments are not available. Shurygin and Kryuk have used the rotating disc method for a study of oxygen diffusion in liquid iron. In their experiments a silica disc was rotated in liquid iron containing oxygen, and the rate of formation of liquid iron silicate was determined by measuring the decrease in weight of the disc. On the assumption that the rate of dissolution was controlled by the diffusion of oxygen in the iron, the diffusion coefficient was computed to be 5.2 x sq cm per sec at 1550°C. However, the Levich equation, which was used to interpret the rate data, was originally de- rived for the case of mass transfer between a solid disc and a single-phase liquid. The hydrodynamic and diffusion boundary layers in the iron stirred by a disc, via coupling of the silicate melt, may be appreciably different from those predicted by Levich's equations. Recently, Novokhatskiy and Ershov, using an identical experimental method to that of Shurygin and Kryuk, obtained a diffusivity for oxygen in liquid iron of 1.22 x 104 sq cm per sec at 1550°C: no reasons were offered for the disagreement. Schwerdtfeger5 has also recently studied the diffusivity of oxygen in liquid iron. He reacted shallow melts of liquid iron, 0.5 to 1.0 cm deep and contained in high-purity alumina crucibles, with appropriate H20-HZ-He mixtures. The total sample was analyzed, without sectioning, to obtain the average concentration of diffusate. A value at 1610°C of D = 12(3) x 10-5 sq cm per sec was obtained from the results of twenty experiments.= Oxygen profile measurements, which were carried out in three additional experiments using long capillaries and the semiinfinite boundary conditions, indicated a diffusivity about half that computed from the shallow bath experiments. Possible sources of error in Schwerdtfeger's study will be discussed later in this paper. EXPERIMENTAL TECHNIQUE The essential arrangement of the diffusion cell is shown in Fig. 1. The liquid iron was contained in an alumina capillary, 3 to 4 mm diam and 3 to 9 cm long, which was supported by a hollow alumina pedestal and this whole assembly was held within a movable alumina reaction tube. This tube, which was about 7 mm in bore

Jan 1, 1970

By Evans B. Mayo

This study examines the structural framework of the Southwest for evidence of the four principal trends of lineament tectonics. It attempts to classify their intersections and compares the positions of those trends that appear most favorable with the positions of the presently known mining districts. This is a controversial topic. The author, in presenting his analysis, is aware that the study of lineament tectonics and relation of ore districts to regional structure is complicated by insufficient data and, unavoidably, by personal bias. The development of lineament tectonics has been summarized by Umbgrove.' Early attempts to fit ore districts into the Cordilleran framework were made by Billingsley and Locke, who do not refer to lineament tectonics, although their appoach is similar; they observe that heat and fluids, including the ore-depositing fluids, are most likely to rise at or near intersections of major structures where the crust is fractured, or weakened, to great depth. As a result of studies distributed over the earth-including ocean basins as well as continents— some tectonists recognize four dominant structural trends: 1) northwest; 2) northeast; 3) nearly east-west, or equatorial; and 4) nearly north-south, or meridional. Baker' proposed a theory to account for these trends and Sonder" called their world-wide arrangement the regmatic shear pattern. Moody and Hill proposed a much more complicated shear network which, although fascinating and perhaps ultimately useful, will not be followed here. In a recent review of deformation within the Cordillera, Wisser mentioned the four fundamental directions. The fact that many geologists deny the existence of the regmatic shear pattern implies that the fundamental structures are far from obvious. It may mean, also, that some geologists are not accustomed to examine regional and world maps analytically. The maps require much study, and certain features should be isolated on overlays. Even so, with the present limited knowledge, uncertainties remain. The following analysis is a qualitative experiment, subject to change as information accumulates, and should be supplemented by the western sheet of the Tectonic Map of the United States.' Many have probably gained the impression that the Cordillera of the West is oriented northwest-southeast and from this, unless experience rules otherwise, it is natural to assume that the structure likewise trends northwest-southeast. To an important extent this is true, yet anyone tracing off the western sheet of the Tectonic Map all the recorded north west-southeast structures may be surprised to see what a small part of the entire area they occupy. In southwestern U. S. (Fig. 1) the northwest-southeast structures are mostly restricted to the eastern, southern, and western parts. The eastern margin of the Cordillera from northern Colorado to the International Boundary near El Paso, Tex., is obviously determined by some structure other than northwesterly ones. In eastern Nevada and far southward toward the Gulf of California many mountain ranges, valleys, and faults are meridional. In the Rocky Mts. of Colorado, and at many places in southern Arizona, the crystalline Pre-Cambrian is foliated northeast-southwest. The Uinta Mts. of Utah trend approximately east-west, in much the same way as a broad belt of transverse, west-northwest structures—the Texas lineament of Hill",' and Ransome10 in southern Arizona, southwestern New Mexico, and southern California. It seems, then that there are four regmatic shear directions in the Southwest, but at many places they are discontinuous, and their projections must be inferred. To clarify these trends the four sets have been isolated into two systems: 1) northwest-northeast and 2) east-west-north-south (Figs. 2 and 3). Northwest-Northeast System: A number of prominent northwest-trending zones of structure are easily recognized. They are designated by circled Roman numerals, the northeast-trending structures by circled capital letters. Perhaps no two geologists would agree completely on the positions of all these belts. Names given below are for convenience only, and may be discarded where other names have priority. (I) The Sierra Nevada-Lower California belt contains the Jurassic-Cretaceous granitic massifs of the Sierra Nevada and Lower California. These

Jan 1, 1959

By F. G. Miller

Earlier studies of the mechanism of flotation showed that coal flotation products are contaminated by composite coal and mineral particles. The mineral matter of these particles has the dual effect of increasing the specific gravity and decreasing the coal surface of the particle available for bubble attachment. The present study showed that these particles are more loosely held in the froth column than the high grade coal particles, especially in the coarser sizes. Consequently, it is possible in the laboratory to produce a secondary concentration of the product by sprinkling the froth product with water to wash these loosely held particles from the froth during formation and recovery. To improve coal flotation, Bethlehem Steel Corp.'s research department has been working on a number of related problems based on both physical and chemical factors. The ideal is of course full understanding of the interrelated effects of both sets of factors. However, detailed studies1 of each physical factor have produced immediately usable results in commercial coal flotation and are providing the basis for our current work on the effect of chemical factors. Since these detailed studies provided the basis for the present investigation of froth sprinkling, their key results are highlighted here. For example, for a highly floatable bituminous coal, particle floatability was shown to be directly related to the percentage of coal surface available for bubble attachment and inversely related to particle mass. The data also underscored that in a given size range the addition of mineral matter to a particle makes flotation more difficult both by adding to the mass of the particle and by decreasing the amount of coal surface available for bubble attachment. Size also affects particle floatability in that as the size of particles of a given specific gravity increases, mass increases and the ratio between coal surface area and mass of the particle decreases. Flotation tests showed that any recovery of the coarsest particles in a coal feed is accompanied by progressively higher recovery of finer and finer particles. Consequently, fine high-specific-gravity particles are recovered with the coarse good coal. As a result, when the entire 14 mesh x 0 size range is processed by froth flotation, the good coarse coal is contaminated with high-specific-gravity fines. It is the small quantities of these fines that must be rejected from the froth product if flotation efficiency is to be increased above its present level. These principles and findings, together with work by Professor V. I. Klassen,' led to the froth sprinkling investigation reported in this paper. Working with Russian coals, Klassen had shown that froth sprinkling washes loosely held refuse particles from the froth and compensates for the water lost in the upper froth layers and increases its stability. As a result, the coal particles are retained in the froth and the loss of coarse coal from the froth is prevented. Decreases in ash of 0.1% to 1.5% were reported by Klassen. Working with a 28 mesh x 0 American coal, the author of the present paper found that single-stage flotation coupled with froth sprinkling produced a product comparable to that obtained by the more expensive method of roughing and cleaning.3 The investigation reported in the present paper expanded study of the effect of froth sprinkling to include each of the standard size fractions of a bituminous coal. The main findings are as follows: (1) For those size fractions between 48 mesh and 150 mesh, froth sprinkling produces a marked reduction in both sulfur and ash. In the case of those sizes coarser than 48 mesh, while this reduction is not as significant, the additional benefit of shorter retention time is still a plus for froth sprinkling. Furthermore, when froth sprinkling is coupled with normal fast flotation, the product purity approaches that obtained by starvation reagent feeding, i.e., gradual feeding of flotation reagent to float the recoverable coal from the pulp slowly in several increments. The specific results for these two size groups are: (a) 14 x 48 Mesh. By using froth sprinkling with fast flotation, product sulfur can be reduced to a level

Jan 1, 1970

By Yoichi Ishida, Ching-Yao Cheng, John E. Dorn

Tile creep behavior of a iron was investigated over the range of temperatures from 375° to 1150°K. Apparent activation energies for creep, obtained by the effect of sudden changes in temperature on the creep rate, revealed the presence of- four distinguishable regions. The creep behavior of a iron in Regions II (480° to 774°K) and 111 (775° to 1045°K) was found to correlate well with a model where the creep rate is controlled by the nonconsevualive motion. of jogged screw dislocations. An anomalous increase in the apparent activation energy fur creep in Region III was found to be in harmony with the known decrease in the free activation energy for self-diffusion over the Curie transformation range. Furthermore, the creep rates in Regions II and III were found to increase with stress due not only to the effect of stress on the activation energy but also to an increase in density of mobile dislocations. The evidence suggests that pipe diffusion along moving dislocations could be a significant factor over the lower temperatures of Region II. ALTHOUGH the creep properties of iron and steel have been of interest to metallurgists for some time and an extensive literature is now available, most of the published information has not produced much toward understanding the fundamental dislocation processes controlling the creep behavior of these materials. Whereas the activation energy for high-temperature creep of metals usually agrees well with that for self-diffusion, Sherby, Orr, and Dorn1 deduced from the data of Tapsell -and c1enshaw2 that the activation energy for creep of Armco iron is about 78,000 cal per mole and more recently Lytton and sherby3 reviewed the data of Jenkins and Mellor4 and found an activation energy for creep of a iron of 80,000 cal per mole. Both values are substantially greater than the activation energy for self-diffusion in a iron. Because of the limited data suitable for purposes of identifying the controlling dislocation mechanisms of creep and the unusually high activation energies for creep quoted above, it was considered desirable to reinvestigate in detail the creep of a iron by more reliable techniques which would provide sufficient data for determination of the creep mechanisms. EXPERIMENTAL TECHNIQUE Creep specimens were prepared from 3/8 by 4 in. iron bar stock of the following composition by weight percent: 0.001 C, 0.0120 0, 0.001 N, 0.004 S, and 0.003 P. The as-received bars were cold-rolled to a thickness of 0.100 in., annealed under argon for 30 min at 1113°K, cold-rolled to a final thickness of 0.063 in., and machined into tensile specimens having gage section 0.250 in. wide and 1.70 in. long. Finally they were recrystallized under argon at 1113°K. Specimens to be crept above 1060°K were recrystallized at 1173°K. Both recrystallization treatments gave the same reproducible equiaxed ASTM No. 4 grain size. Creep testing was conducted in machines fitted with Andrade-Chalmers arms that maintained a constant stress to within ±0.2 pct of the reported values. Deformation over the specimen gage section was sensed with a linear differential transformer and recorded auto-graphically as a function of time. The strains deduced from these measurements were sensitive to ±5 x 10-5. During creep the specimens were contained in an argon-filled chamber which was immersed in a temperature-controlled molten tin bath. Creep temperatures, as measured by thermocouples attached to the specimen, were maintained to within ±10.25°K of the reported values. For activation-energy determinations rapid changes in temperature of about 12°K were obtained within 30 sec by direct resistance heating of the creep specimen, and maintained to ±1°K of the reported values. Observations of structural details of specimens before and after creep tests were made by electrolytic polishing and etching in acqueous ammonium persulfate. EXPERIMENTAL RESULTS A typical example of the determination of apparent activation energies, Q, by the effect of small changes in temperature is illustrated in Fig. 1. Q is defined by

Jan 1, 1967

By J. J. Pfeiffer, A. A. Oming, J. W. Myers

The method used to process washery refuse for use as a building material aggregate is described. Results of studies made in investigating this process are summarized. The Bureau of Mines, in cooperation with two commercial coal producers, studied the feasibility of using coal washery refuse in manufacturing lightweight aggregate for building materials. Lightweight aggregate is used in making masonry products of lower bulk density than can be produced from conventional aggregate such as sand, gravel, and crushed rock. Among the advantages of lightweight masonry are reduced deadweight loads which reduce materials cost, lower cost of construction labor, and improved thermal and sound insulation. A process for making lightweight concrete aggregate from coal-washery refuse appeared attractive because of the availability of raw material which contained sufficient fuel to carry out the processing. Vast quantities of such refuse have accumulated throughout the coal-producing areas of the U.S. Generally it is stored in unsightly piles, creating a problem of air pollution and using space that could be devoted to other purposes. Conversion of the refuse to lightweight aggregate could remedy these conditions and effectively utilize a material normally considered waste. The first part of the program was conducted in cooperation with the Truax-Traer Coal Co. to determine the feasibility of utilizing refuse from the company's coal preparation plant at Ceredo, W. Va., in manufacturing lightweight aggregate. The aggregate was to be used in making cement blocks comparable to those known in the building trades as cinder blocks. This project was completed in 1954. Based on the results of this work, the cooperator began commercial development and formed the Trulite Corp. for this purpose. A plant capable of handling 120 tpd of refuse was constructed for approximately $120,000. It began operating on June 1, 1955.' In 1959 a similar study was made, under a cooperative agreement with the Carbon Fuel Co. of Charleston, W. Va., to rrscertain the suitability of refuse from its plant for conversion to lightweight aggregate. Under this program various modifications of equipment were made, and the investigation of process variables begun in the previous project was extended. PROCESSING METHOD The formation of lightweight aggregate from coal-washery refuse js governed, in part, by the fusion characteristics of the mineral constituents of the refuse when heated above their softening temperatures. The heat required for this purpose is supplied by the combustihle matter in the refuse. Some of this combustible matter is coal with high ash content, but most of it is carbonaceous shale. Whereas the density of aggregate prepared from clay or shale depends largely on expansion, the density of aggregate made from refuse is reduced by burning out the combustible matter and, to a lesser extent, by expansion of the residue in the fuel bed. A chain-grate stoker with an updraft air system was chosen for the conversion process in preference to a downdraft sintering machine or a rotary kiln. In down-draft sintering, which operates on the overfeed burning principle, the volatile matter from the refuse would distill off without burning, resulting in inefficient combustion and interfering with proper operation of the equipment. The rotary kiln is designed primarily for processing material that lacks, or is low in, combustihle matter with the heat of reaction supplied from external sources. Washery refuse, which is relatively high in carbonaceous material, could burn at the wall of the kiln causing clinker formation with resultant decrease in throughput and increase in production cost. With an up-draft air stream on a chain or traveling-grate stoker, the volatile matter would be carried into the ignition and burning zones where the heat could be utilized more efficiently. The combustion procedure in this type of unit may be considered as consisting of four periods: 1) initial ignition of the top of the bed by furnace radiation, 2) ignition travel through the bed until the whole bed is ignited, 3) burnout of the combustible material with high air rates to fuse and expand the residue, 4) a cooling period. Investigation of the process required a study of the variables that affect both

Jan 1, 1962

By F. V. Lenel, G. S. Ansell, J. A. Dromsky

The growth of aluminum oxide particles in a nickel matrix was studied eve?. the temperature vange of 2140° to 2470°F. The instability of the dispersed alumina was shown to be independent of the crystal structure of the alumina. The activation energy for the growth of the dispersed alumina was found to be 84.7 1 2.0 kcal. The particle radius increased as a function of time. These results indicate that the growth is not diffusion controlled. It is believed that the rate controlling mechanism is the dissolution of the aluminum and oxygen atoms into the nickel lattice. THE strength properties of alloys which consist of a finely dispersed second phase in a metallic matrix depend upon the spacing between the second phase particles. It is therefore desirable to achieve very fine dispersions in these alloys. Furthermore, to retain the properties these very fine dispersions must be stable during fabrication and service of the alloys. The best known of the dispersion strengthened materials, the SAP type alloys, which consist of a dispersion of aluminurn oxide in aluminum, have exceptionally good stability up to the melting point of aluminum. There is evidence, however, that other dispersion strengthened alloys, even those consisting of refractory oxides in a metal matrix, may be less stable. This investigation is concerned with the stability of Ni-Al2O3 alloys in the temperature range in which these alloys are usually fabricated. The mechanical properties of Ni-Al2O3 alloys at elevated temperatures have been previously investigated by Crelnens and rant,' and Gregory and Goetzel.2 The behavior of these alloys in stress rupture tests appears to indicate that at temperatures below 1800°F they are highly stable. There is some doubt, however, as to their stability at the higher temperatures used during the conventional fabrication. Cremens and Grant, in preparing their test alloys, cousolidated, by powder metallurgical techniques, nickel powders as fine as 0.13 µ diam and alumina powders as fine as 0.018 µ. Metallo- graphic examination of the alloys following fabrication revealed that none had interparticle spacings of less than 2 µ. Considering the size of the original component powder particles, it is likely that the dispersions coarsened during fabrication. Gregory and Goetzel, in their studies of extruded alloys of 80 pct Ni—-20 pct Cr matrixes with nonmetallic dispersion, observed a definite coarsening of the alumina dispersions in alloys sintered at 2280°F as cantrasted to those sintered at 2000oF. Similar observations on the spheroidization and growth of thoria particles finely dispersed in a nickel matrix were made by D. K. Worn and S. F. Marton.3 As a result of such coarsening, much of the effort expended in the preparation of very fine powder mixtures would be lost. The mechanical properties of the alloys which had experienced coarsening would be expected to be poorer than if the original dispersions had been retained. EXPERIMENTAL PROCEDURE Ni-Al2O3 alloys were produced from powder prepared by a coprecipitation technique. Aluminum hydroxide and nickel hydroxide were coprecipitated from chloride solutions of the metals. The mixed hydroxides were calcined to form metal oxides and the nickel oxide in the mixture was selectively reduced to nickel by treating it in hydrogen. Specimens were compacted from the resultant powder which consisted of a fine, uniform mixture of aluminum oxide and nickel particles. The compacts were sintered by resistance hot pressing4, a densification technique which requires exposure times of the order of only a second or less at elevated temperatures. A conventional sintering process was not used, since the temperature required for densification would have to be in the region in which the stability of the dispersion was to be studied. A series of hot pressed specimens were treated in hydrogen at temperatures from 2140o to 2470°F (1171o to 1355oC, for times up to 120 hr. Changes in the microstructures were studied by electron microscopy using the two-stage preshadowed carbon replica method.5 In performing a lineal analysis on a series of micrographs from each specimen it was found more convenient to determine the mean free path between alumina particles rather than particle radii as the parameter of growth. Although these quantities are directly proportional for only spherical particles, the alumina particles in these alloys

Jan 1, 1962

By R. A. Wattenbarger, H. J. Ramey

A systematic study has been made of the application of the real gas pseudo-pressure m(p) to short-time gas well testing. The m(p) function can be used in real gas flow problems to account for the variation of viscosity and z-factor with pressure. A mathematical model was solved numerically to generate solutions of various real gas flow problems. The effects of turbulence, formation damage and wellbore storage were included in the model. The analysis of simulated well tests showed that the interpretation methods normally used for liquid flow are generally accurate when the m(p) is used. For practical rates without turbulence, the solutions are not rate-sensifive, and the flow capacity.. kh and skin effect can be determined accurately from either a buildup or a drawdown test. However, the kh calculated from a drawdown test can be significantly low when turbulence is present. A case was simulated in which this error wa.s 36 percent. Turbulence does not affect the determination of kh from buildup tests. The proper determination of kh from p and pa buildup and drawdown plots is developed by analyzing their relationship to the m(p) method. A simple equation is given that can be used for long-range gas well performance forecasting. This expression is compared with a method pre-sended by Russell et al." Introduction The formation flow capacity and wellbore damage condition can be determined for liquid-producing tests by means of buildup and drawdown tests. These tests make use of short-time pressure transient data rather than stabilized flow tests that are often used for gas well testing. Tracy' presented a method of gas well testing that was based on ideal gas equations and utilized p2 buildup plots. This method was shown to be good for low pressure wells. MatthewsZ suggested plotting p and using an average slope and average gas properties. This method was more successful on high pressure wells. Al-Hussainy, Ramey and Crawfordh howed that variation in gas properties could be simplified by using the real gas pseu do-pressure m(p). Some cases showed that the use of the m(p) function Original manuscript received in Society of Petroleum Engineers office Aug. 8. 1967. Revised manuscript received June 11, 1968. Paper (SPE 1835) was presented at SPE 42nd Annual Fall Meeting held in Houston Tex., Oct. 1-4. 1967, and at SPE 4th Annual Eastern Regional ~eetin$ held in Pittsburgh. Pa., Nov. 2-3. 1967. 0 Copyright 1968 American Institute of Mining, Metallurgical, and Petroleum Engineers, Inc. 'References given at end of paper. provided an accurate method of interpreting buildup and drawdown tests. Because the work of Al-Hussainy et al. included only a few actual cases, there was a need to explore the use of the m(p) method for a greater variety of flow conditions. The occurrence of turbulent flow around the wellbore often is an important factor in gas well testing. Swift and Kiel,' and Carter et a1.' treated turbulence in gas well testing, but did not consider the variation of viscosity and z-factor with pressure. This paper presents the results of an investigation of the application of the m(p) method to buildup and drawdown testing. Vhe investigation included the effects of turbulence, formation damage and wellbore storage. Gas Flow Equations To formulate the mathematical model, many of the assumptions usually used in well testing theory are applied. The system has radial geometry with a closed outer boundary and is composed of a horizontal porous formation that has uniform and isotropic rock properties and uniform thickness. Allowance is made, however, for a radial region of reduced permeability near the wellbore. This region represents formation damage. The geometry of the system is shown in Fig. 1. Darcy's law does not always apply to gas flow. A more general expression is needed for non-Darcy, or "turbulent" REGION OF DAMAGED PERMEABILITY SEALED BOUNDARIES ClRCULAR WELL BORE Fig. 1—Radial flow model. AUGUST, BT7

Jan 1, 1969

By Nicholas J. Grant, A. W. Mullendore

Measurements of grain boundary sliding were made on polycrystal and bicrystal tensile creep specimens of Al-2 pct Mg at 500oand 700oF. Grain and pain boundary orientation factors were studied with respect to their effect on the magnitude and direction of grain boundary sliding. It was concluded that a large part of the observed sliding may be caused by the operation of slip crossing the grain boundaries. Supporting evidence for this model from other investigations is also given. GRAIN boundary sliding in metals during elevated temperature deformation has frequently been thought of as an independent mechanism with its own characteristic stress, temperature, and structure dependence. Further, it is common to regard the grain boundary as the weak link during creep deformation since: 1) the strength of apolycrystallinemetal, as meas-sured by creep rate or rupture life, decreases at\a more rapid rate when sliding comes into action, and 2) grain boundary sliding, as revealed by "boundary darkening" and fold formation is often the first visible deformation. On the other hand, more quantitative observations, in both polycrystals and bicrystals, have shown boundary deformation to be closely related to the over-all deformation of the specimen. The curve of grain boundary contribution to elongation vs time takes a shape similar to that of the total creep curve,',' and the apparent activation energy for sliding appears in most cases to have the same value as that for total creep.3 In order to reconcile these observations, one is left with these possibilities: 1) The grain boundary and grain deformation rates are both controlled by the same rate limiting processes. 2) The rate of sliding is limited by accommodation deformation in the grains. 3) Sliding is the consequence of grain deformation. 4) Grain deformation is dependent on prior grain boundary deformation. Alternatives 2) and 3) would view grain boundary sliding as occurring in series with another deformation process which could be rate controlling. McLean prefers item 3) as the sliding basis,4 where subgrain rotation results from dislocation accumulation in the substructure at the grain boundary. Crussard and Friedel's5 treatment of grain boundary migration considers that dislocations are forced into the grain boundary and are dissociated into dislocations with Burgers' vectors nearly parallel to the grain boundary. The movement of these partial grain boundary dislocations then produces the sliding. The results of this investigation have indicated alternative 3) to be the most appropriate and it is this aspect of grain boundary sliding which will be analyzed. EXPERIMENTAL PROCEDURE Three types of tensile creep tests were employed in this study: (I) A constant load test of one very coarsegrained polycrystalline specimen at 500° F, 3600 psi, 1 pct per/hr strain rate to examine the grain and grain boundary orientation factors in grain boundary sliding. (II) Constant strain rate tests of polycrystalline specimens at 510o and 715o F, 2 pct per/hr strain rate to determine the grain boundary contribution to elongation at very low total strains. (m) Constant strain rate tests of bicrystals of various grain and grain boundary orientations at 500°F, 2 pct per/hr strain rate to obtain further information on the orientation effects in grain boundary sliding. The material used in the tests was a high-purity Al-1'.92 mg alloy, kindly furnished by Alcoa. Its composition is given in Table I. The alloy is a solid solution at the test temperatures. The coarse grained polycrystalline specimen for test (I) was prepared by recrystallizingthe machined specimen, straining it about 1pct intension at room temperature and annealing at 900°F for 8hr. The constant strain rate specimens (11), with an 178 in. sq. by 1/2 in. long gage section, were cut from 1/8 in. thick sheet of cold rolled and annealed (8 hr at 1000oF) material. The latter specimens were supported in a special jig during the milling operations to avoid any bending, and the gage section was subsequently heavily electropolished to remove the cold-worked surface layer. The bicrystals (ID) were machined in the same

Jan 1, 1963

By D. J. Rose, W. M. Armstrong, A. C. D. Chaklader

The wetiability of single crystals of MgO by specimens of vacuum-cast iron was studied using the sessile drop technique in vacuo at 1550ºC. Formation of FeO at the liquid-vapor interface caused the contact angle (6) to decrease from 117 to 65 deg during the first minute. After cooling, all specimens possessed a peculiar annular interfacial deposit of Fe,O,. Within the annulus the interface showed no sign of chemical attack. Chemical reaction occurred where the iron was alloyed with Ti, V, Cr, Nb(Cb), Ta, cmd Zr. Vnriation of 6 with alloy concentration was studied. Although vanadium md chromium improved the wettability of MgO by iron, the effect of Zr, Ti, Nb, and Ta was indeterminable because the 8 derived from sessile drop considerations was that of the metal against a restrictive peripheral volume of liquid oxide wetting the substrate. INTEREST in the nature of metal-ceramic interactions has been stimulated by progressive development of metal-ceramic combinations. One of the more valuable methods used for bond investigation is the sessile drop technique. Recent attempts to improve wettability through solute additions to the drop have revealed that the solute may a) react with the oxide creating new compounds at the solid-liquid interface, b) adsorb at the interface in a monolayer formation, or c) distribute throughout the drop uniformly causing no wettability variation. This work, part m in a series1,2 of investigations of interface reactions between metals and ceramics, embraces a study of the Fe-MgO system. MgO possesses a large negative free energy of formation (-173 kcal per g mole 0, at 1550ºC) compared to FeO at this temperature (-73 kcal per g mole 02). Hence the electronegativity difference between the drop and substrate allows selection from an extensive group of metal solutes with intermediate electronegativity differences that would, in the absence of chemical reactions, be expected to adsorb pre- ferentially at the solid-liquid interface. However, the high oxygen pressure of MgO in vacuo creates an oxidation environment which influences the solute behavior. Recent studies of the Fe-MgO systemJ74 have been restricted by chemical reactions that obscured observations. In view of the current interest in liquid-phase sintering of metal-ceramic combinations under partial oxidizing conditions,5 a more comprehensive study of the Fe-MgO system seemed beneficial. EXPERIMENTAL PROCEDURE 1) Specimen Preparation. Optical-grade MgO single crystals and Ferrovac vacuum-cast iron were used in all experiments. A spectrographic analysis of the iron indicated 0.008 C, 0.05 V, 0.005 Mo, 0.01 Ti, 0.002 S, 0.01 Cr, 0.001 Mg, 0.01 Si, 0.003 A1 (wt pct). The MgO crystals were cleaved along (100) planes into plates approximately 20 by 20 by 1 mm. Following annealing in vactto at 1100°C for 3 hr, surface irregularities were removed by a chemical etch in phosphoric acid at 100°C. The iron rod was machined into small cylinders approximately 0.250 in. in diam and length and these were carefully cleaned in organic and acidic solutions to remove surface impurities. All specimens were weighed to the nearest 0.1 mg. 2) Apparatus. The apparatus described in detail by previous authors1,2 was modified for this work. A Polaroid Land Camera was incorporated into the optical system so that drop measurements at ten-second intervals after melting were possible with ultra-high-speed self-develop ing film. 3) Procedure. The iron cylinders were placed upright on the MgO plates and positioned in the molybdenum susceptor. After furnace assembly the system was pumped down to a vacuum of 1 µ and flushed with H2 at 800°C. The system was again evacuated to 5 x 10-5 mm of mercury and the temperature raised to 1550°C within 2 min. The molten drops were photographed at appropriate time intervals. TWO min after melting the iron vapor pressure caused gas discharge, or corona, within the tube. The specimens were then slowly cooled to room temperature, sectioned and examined metallographi-cally. X-ray identification of interfacial reaction products was attempted by the powder diffraction technique. Alloying elements (Ti, V, Cr, Zr, Ta, Nb) were added to the iron in various concentrations up

Jan 1, 1963

By J. M. Tobin, L. H. Cadoff, L. M. Adelsberg

The reaction between liquid zirconium and graphite at temperatures above 2000 °C has been investigated. The reaction products were found to be carbon-saturated zirconium metal and ZrC which formed between the graphite and the metal. Parabolic growth behavior was observed for the ZrC Phase at all temperatures of this investigation. The parabolic growth constant at temperatures between 2000° and 2860°C was measured to be 1.83 exp 84,300/RT) sq cm per sec. The reaclion mechanism was proposed to be the rapid carbon saturation of the liquid metal and the formation of ZrC at the metal-carbide interface with diffusion of carbon through the ZrC, the "rate - determining step" of the reaction. The concentration-independent diffusion coefficient of carbon in ZrC, (DZrC), was expressed as 0.95 exp? 78,700/RT) sq cm per sec. This value mas calculated using the temperature-invariant ZrC phase fields proposed in the literature. The [ZrQiq)] — [ZYQiq) + ZrC] phase boundary over the temperature range 2000° to 2800°C was determined and the ZrC + C eutectic temperature was found to be 2890° ± 50°C. ThE Group IV and VB transition-metal refractory carbides are of interest because of their high melting points, high temperature strength properties, and relative inertness in certain corrosive environments. The present-day understanding of these materials, however, is limited by the general unavailability of accurate and reliable physical and chemical property data. This is due primarily to the difficulties associated with the preparation of suitable, high-density, high-purity carbide samples, and the achievement and control of uniform high-temperature environments. Accurate measurement of temperature is also an important factor limiting the reliability of reported data. The Zr-C reaction was selected for investigation because of the high melting point (>34000C) and favorable nuclear properties of the reaction product, ZrC. The direct reaction method afforded an opportunity to obtain kinetic data on the fully dense carbide. In this paper, layer-growth techniques were used to estimate the diffusivity of carbon in ZrC and to investigate the phase equilibria in the Zr-C system at temperatures above 2000°C. EXPERIMENTAL PROCEDURES AND DATA Crystal bar zirconium (99.9 pct Zr), purchased from the Nuclear Materials and Engineering Corp., and ZTA graphite (99.9+ pct) were used in this study. The analyses of the materials are listed in Table I. A schematic of the high-temperature carburization apparatus is shown in Fig. 1. The sample was a zirconium metal charge in a 1/2 -in.-ID graphite crucible which was capped with a tightly fitting graphite stern. The crucible and stem were designed to closely approximate black-body conditions. The graphite crucible was packed in lampblack (outgassed 1 hr at 2100°C) which provided insulation and thermal stability to the system. The inert atmosphere was maintained by a constant flow of high-purity argon or helium gas. The nozzle-diffuser section on the top flange was sufficient to prevent back-diffusion of air into the system. Chemical analysis of the carburized samples revealed oxygen and nitrogen concentrations of less than 44 and 20 ppm, respectively. The sample was heated by induction with a Westing-house 5 kw-450 kc power source. Temperature was measured with a Milletron two-color pyrometer which was sighted into the crucible by reflection from an overhead front surface mirror. Optical losses due to glass absorption and light reflection at air-glass interfaces were minimized by the employment of the dynamic-gas seal. It was found that no correction was required for the front surface reflector. The pyrometer was calibrated periodically against a U.S. Bureau of Standards tungsten ribbon secondary standard. Temperatures inside the crucible were also calibrated against the Zr-ZrC eutectic (1850°C)1 and the Nb-Nb2C eutectic (2335°C)2,3 temperatures and agreed to within ±10°C of these values. The temperature was controlled by manually adjusting the power input; variations of

Jan 1, 1967

J. Chipman (Massachusetts Institute of Technology, Cambridge, Mass.)—The fact that the experimental work has been applied to copper rather than iron and that the paper is presented to the Iron and Steel Division, I regard as rather significant. It shows the unity of metallurgy and the fallacy of trying to cut it up by metals. This result for the solubility of sulphur in molten copper correlates with Professor Schuhmann's finding that the published data on the other side of the copper-sulphur miscibility gap are also in error. I should like to ask the author to say a little bit more about the sulphur capacity of the slag. T. Rosenqvist (author's reply)-—I hope that Dr. Chipman will find the derivation of the expression for sulphur capacity more clearly explained in the printed version of the paper than in the oral discussion at the meeting. I feel that this quantity, which actually is the ratio of two activities, can be measured more easily than the individual activities. Even if the ratio CaO/CaS is chosen as the standard state, the expression can be used for any slag, even for slags completely free of lime, and it represents a way to put the desulphurizing power of all slag constituents into one bag. Some doubt has been expressed as to whether oxygen ions really exist in calcium oxide and in molten slags. From a thermodynamic view point that question is of minor importance. The term oxygen ion activity, or any activity for that matter, is defined rigorously by the equation: activity = exp u/RT, where u is the change in free energy connected with the transfer of one mol of ions from the standard state into the slag. Whatever happens to the ion in the slag is of no concern to the thermodynamicist. Regardless of whether the ion is "free" in the slag or not, or whether it is present in a very small amount, its activity can always be expressed, and for a thermo- dynamic calculation that is all we need. However, ionic activities will only be of some real value if they are simple functions of the slag composition, or can be measured easily. Concerning the real nature of the oxygen in the slag, my feeling is that the oxygen atom has a rather multiplex nature depending on how strongly it is tied by covalent forces or polarized by the other atoms or ions present. The oxides of iron, cobalt, and nickel differ from calcium oxide and blast furnace slags as to the amount of free electrons that can give rise to electronic conductivity. In slags we know that the conductivity is mostly ionic. The fact that reversible emf's can be obtained with oxygen electrodes in certain salt melts, indicates a significant amount of oxygen ions in these melts. But extended work, e.g. polaragraphic studies and measurements of transference number; are needed to obtain quantitative information about the real structure of the slags. D. E. Babcock (Republic Steel Corp., Youngstown, Ohio)—-With reference to the ion, it might be well to remember Dr. Moses Gomberg. All of his life he had no use for the ionization theory and he contributed greatly to the field of chemistry on the assumption there was no such thing as ions. I do not think we have to worry about whether the oxygen is ionic or not. I think one thing specifically should be brought to your attention and this I think is one of the important contributions of Dr. Rosenqvist. He pointed out what we know as oxygen potentials or what is described as oxygen potentials. I have used this concept for a long period of time and I want to state that if this concept is properly applied, it vitiates much of what we have in the literature, or makes our usual ideas regarding oxidation seem primitive. That one thing is more valuable than almost all the rest of the discussion as a fundamental basis on which to build a reasonable in-

Jan 1, 1952

By R. J. Burgess, A. R. Ramey, A. R. Adams

During interpretation of pressure buildup tests on gas wells in a tight dolomite gas reservoir, peculiar behavior was noticed. Two straight lines were apparent. Effective permeability to gas taken from either straight line was about the same, and the Miller-Dyes-Hutchinson dimensionless time check for the straight line was proper for both straight lines. Geological data indicated the likelihood of scattered trending fractures in the reservoirs. Since the first straight Iine yielded permeability values close to the geometric mean permeability from core analyses, it was postulated that the reservoir model was that of an acidized well completed in the tight dolomite, but that widely scattered hairline fractures caused the mean permeability of the reservoir distant from the well to be higher than the matrix permeability. Because all other studies of fractured reservoirs to the authors' knowledge assumed that the fracture matrix was dense enough to communicate directly with the well, no interpretative methods were available. The Hurst line-source solution for a radial change in permeability for interference between oil reservoirs was adapted to pressure buildup testing. The result indicated that the first straight line should yield the proper matrix permeability and wellbore skin effect. The second straight line may be extrapolated to obtain static pressure. The time of bend between the straight lines was used to estimate distance to a fracture. Application to field test data is shown. It is believed that the methods developed and the case history presented will add to present tools available for pressure buildup interpretation. Introduction Since the pioneer studies by Miller, Dyes, and Hutchin-son1 and Horner' in 1950 and 1951, well test analysis has become recognized as one of the most powerful tools available to both production and reservoir engineers. Well test analysis serves as a logical basis for well stimulation and completion analysis, and for long-term reservoir engineering. Since the early 19501s, much effort has been placed on the development of well-test analytical methods. Reservoir and well conditions of increasing complexity have been considered systematically to provide the analyst with a catalog of causes and effects. Matthews and Russella state that some 200 papers dealing with this subject have been published in the last 35 years. Developments in well test analysis appear to have originated in one of two ways. Either a physically realistic field condition was anticipated and analytical solutions for the condition achieved, or anomalous field test behavior was recognized and interpretative methods sought for the anomaly. In recent years, it has appeared that the latter has inspired an increasing number of studies. The analyst today finds an increasing number of known cause and effect studies available for well test analysis, the classic of which is that of finding the specific flow problem that generated the answer — the well behavior. Although it may be impossible to achieve this goal uniquely, the analyst often is able to select a useful interpretation that combines all known performance and geologic data — or to show that various logical alternatives would not significantly affect the interpretation. During a recent reservoir study, we observed gas well test behavior that did not appear to fit behavior described previously. Although it cannot be said that we have found a unique interpretation, we shall present in this paper the peculiar behavior observed, and describe the reservoir and interpretative methods developed. Reservoir Description The subject gas reservoir is a 9-mile-long, narrow dolomite reservoir lying within a limestone of Ordovician age. (See Fig. 1.) The dolomitized rock in the field consists of dark brown to buff, dense to coarsely crystalline, vugular dolomite containing numerous hairline fractures, many of which may have been closed in the reservoir and parted when cores were brought to the surface. Larger fractures are also apparent in core, but usually are filled and sealed with euhedral dolomite crystals. Portions of the north flank of the reservoir are known to be cut by a sealing fault downthrown to the north. Gas wells located near the fault have higher open flow potentials than those more distant from the fault. This is believed to be a result of higher permeability near the fault due to more extensive and open fractures. Detailed coring and core analysis have been performed on several of the wells in this reservoir. Fig. 2 presents permeability variation' plots for both horizontal and vertical

Jan 1, 1969

By D. J. Carney, E. C. Rudolphy

IT has been demonstrated that inclusion size, distribution, and composition affect the machin-ability of resulphurized steels. Merchant and Zlatinl concluded that large sulphide inclusions aided machining by forming a (lubricating) coating on the tool face. Boulger et al.² and Van Vlack³ noted that the size, distribution, and composition of the inclusions in the steel affected the machinability. Steel specimens containing large globular sulphide inclusions usually exhibited excellent cutting properties, while machinability was adversely affected by the presence of numbers of oxide-type inclusions. Consequently a thorough knowledge of all the factors which affect the inclusions in the final product is desirable. Since almost all the inclusions have their origin in liquid steel, it was necessary to begin a study of inclusions in free-machining steels by studying the inclusions and chemical segregation in the as-cast ingot. Very little information is available on the size, distribution, shape, and composition of inclusions in large, capped, free-machining steel ingots, particularly the B1113 grade. Gregory and Whiteley4 made a general study of the inclusions in a small, high sulphur, free-machining steel ingot. Also, numerous authors have described the solidification and segregation characteristics of the four basic types of steel ingots, namely, rimmed, capped,7 semikilled,7,8 and killed7,9,10 ingots. Most of these studies were made with plain carbon or low alloy, low sulphur steel. It was desirable to study not only the ingot but also the change in size, shape, and number of inclusions on rolling an ingot to a bloom and thence to a billet. This procedure was followed and it is hoped that this study may serve the dual purpose of adding to the general knowledge of ingot solidification as well as contributing to the knowledge of the size, shape, distribution, and composition of inclusions from the ingot to the billet in a high sulphur, free-machining steel. Procedure A 12.000-lb 23x35x75 in. slab ingot of the B1113 grade was cast, sectioned, and studied both macro-scopically and microscopically. An adjacent ingot from the same heat and of the same size was rolled to a 77/8x77/8 in. bloom and thence to a 21/2x21/2 in. billet. These various bloom and billet sections were also sectioned and studied macroscopically and microscopically. Sectioning the Ingot: The ingot herein described was obtained from the United States Steel Corp.'s South Works Bessemer Blow No. 0193, a B1113 mechanically capped heat. The 23x35 in. ingot (No. 2) was teemed according to normal procedures and after stripping and transportation to the rolling mill was not placed in the soaking pit but allowed to air cool in an upright position. When completely solidified, the ingot was cut into sections by means of a powder scarfing torch and further sectioned by saw cutting as indicated in Fig. 1. Cut No. 2 (1x10x12 in.) from sections A through H was cleaned thoroughly, macroetched in a solution of 50-50 water and hot muriatic acid and used to obtain a macrograph of a horizontal section from the surface to slightly beyond the center of the 23-in. ingot dimension. Cuts No. 5 and 3 (lx81/2xl0 in. each) from sections A through H were treated in a similar manner to obtain a macrograph of a horizontal section from the surface to slightly beyond the center of the 35-in. ingot dimension. The composite macrograph of these horizontal ingot sections, which shows a vertical section of the ingot from top to bottom, is shown in Fig. 2. It should be noted that sections No. 2 are normal to sections 5 and 3 in the composite. Drillings for chemical analyses were obtained from selected positions within the above-mentioned ingot sections as noted in Fig. 3. The oxygen content was determined by the vacuum-fusion method. Samples for microscopic examination were cut from

Jan 1, 1954

By Joseph L. Gillson

IT is a somewhat curious circumstance that the newest of the Institute's several awards should be conferred in the oldest of our several professional fields—for there is little question that geology and mining have antedated metallurgy and petroleum. Perhaps so did geophysics. After all, one geophysical instrument, the magnetic compass, has been successfully employed for many centuries. Young in years the Jackling Award may be, but in distinction—whether in terms of the title it bears or in terms of the first three medalists (Reno Sales, E. D. Gardner, and the late Father MaceIwane)—the Jackling Award is already one of the ranking awards of the Institute. The name of the 1957 medalist will bring added luster to an already illustrious roll. J. L. Gillson was born and brought up in Evan-ston, Ill., in an area characterized by singularly flat and uninspiring topography, and devoid of ore deposits and even of bedrock—a most unlikely environment from which to expect a future Jackling medalist to emerge! On graduating from high school he enrolled as a chemistry major at Northwestern University. At the end of his frkshman year his scholastic prowess had won him a prize of two chemistry books. But these went unread, for about this time he came under the spell of U. S. Grant, that distinguished and enthusiastic geologist who for many years was head of Northwestern's geology department. From thence forward Joe was a dedicated geologist. After a Navy interlude in World War I, he went to MIT for graduate work, substituted for the late Professor Palache for one year at Harvard in mineralogy, received his doctorate at MIT and stayed on as a meniber of the MIT faculty, where the geology department was headed by that great dean of all economic geologists, Waldemar Lindgren. In 1926 Joe spent a summer looking into barite deposits for the du Pont Co., and since 1928 he has been du Pont's chief geologist. Do men who lead such active lives in our exploration organizations generally have permission to publish their observations? Do they generally find time to publish, even if they have permission? Alas, no. But in Joe Gillson's case, yes. His bibliography runs to well over two score papers, including one of the very first papers in this country on the applications of petrography to problems of Portland cement. He has written on such topics as the origin of talc and the genesis of ilmenite, fluorspar, and alkaline rocks. Each year for a number of years his annual reviews of industrial minerals in the Mining Congress Journal have been eagerly awaited. - Do men who do so much to advance their science have anything left to give to the advancement of their profession? In Joe Gillson's case, yes. He has been a vice president of the Mineralogical Society of America and has been active in the councils of the Society of Economic Geology, particularly in their research programs. This year he is president of the American Geological Institute, currently perhaps the most challenging post to which a geologist may be elected. His contributions to AIME scarcely need documentation, so I shall mention only three—his chairmanship of the Committee on Democratization of the Institute, a committee whose work and recommendations brought a new and hopeful leaven to AIME at a time when this was sorely needed; his chairmanship, a few years ago, of the Industrial Minerals Division; and his current place on the Board of Directors as Vice President of the Institute. Such' a catalogue of contributions and achievement, even as condensed as I have had to make it, certainly merits the Jackling Award. Yet, implicit although not often expressed in the conferring of most awards is the hope that besides recognizing a distinguished career, it will provide stimulus for further effort. In the present instance, such an expectation is already well on the way to fulfillment. Only recently Joe Gillson has undertaken the editorship of a new and greatly revised edition of the Institute's Seeley Mudd volume, Industrial Minerals and Rocks. And so, Mr. Chairman, it is with pleasure and enthusiasm that I present Dr. J. L. Gillson for the 1957 Jackling Award.

Jan 1, 1958

By W. F. Domis, F. von Gemmingen, R. W. Whitmore, F. Garofalo

Constant-load creep-rupture tests at 1100°, 1300° and 1500°F were made on a Type-316, 18 Cr-8 Ni-ZMo, austenitic stainless steel to determine the relationship between ruptzire life and other aspects of creep behavior, The test program was designed on a statistical basis to check the variability of various creep properties and to determine emprically the dependence of minimutn creep rate and ruptzire life on initial stress, The dependence of the rupture life, tn, on the initial stress, a,, is found to be related to the stress dependence of the wrinimum creep rate, i,, and secondary creep strain, The instabilities or breaks found in the conventionual log - logt, plots can be traced to either a change in the linear dependence of log i, on log a, or to a change in secondary creep strain. In the alloy tested, rupture is found to be predowmanantly of the inter crystalline type, For this material, the secondary creep strain is -found to depend on the nature of the grain boundary precipitate, which affects grain boundary migration. It has been shown for a variety of metals and al tested at elevated temperatures under creep conditions that the empirical relationship exists between rupture life, tr, and minimum creep rate, i,. This relationship, in which C and a are in some cases independent of stress or temperature, shows in a general way that creep rupture depends on creep behavior prior to tertiary creep stage. To determine more fully the factors which control creep rupture requires, therefore, more knowledge of the relationship between rupture and creep behavior. To obtain such information, constant-load creep-rupture tests have been made on a Type-316, 18 Cr-8 Ni-2 Mo, austenitic stainless steel at 11000, 1300°, and 1500°F. These tests were designed on a statistical basis to determine the variability of the various creep properties measured and to determine empirically the stress dependence of minimum creep rate and rupture life. As a result of this work, various empirical relations have been established defining more closely the factor, C, in the relation between rupture life and minimum creep rate. This factor is found to be proportional to the amount of strain during secondary creep. The variation of secondary creep strain with rupture life and temperature is determined and its effect on rupture behavior discussed. The variability in minimum creep rate and rupture life is also dis- cussed and the dependence of these quantities on initial stress is interpreted. TEST MATERIAL AND TEST PROCEDURES The material tested is an austenitic stainless steel, AISI Type 316, of the following composition: C, 0.07 pct; Mn, 1.94 pct; P, 0.01 pct; S, 0.021 pct; Si, 0.38 pct; Cr, 18 pct; Ni, 11.4 pct; Mo, 2.15 pct; All 0.003 pct and N, 0.043 pct. The material was received as a hot-rolled 0.5-in. diam bar which was sectioned into 3.25 in. long blanks. Each blank was heat treated by holding 1/2 hr at 2000° F followed by water quenching. A 3/8-in. coupon was removed from each heat-treated blank and the microstructure was examined, the grain size determined, and the hardness measured. The microstructure was found to be uniform from blank to blank and the grain size was found to range from 4 to 6 ASTM numbers. Hardness, measured using a 20-kg load, varied between 127 and 148 Vpn. Tensile creep-rupture specimens having a 0.25 in. diam within a 1.5-in. reduced section were machined from the heat-treated blanks. Constant-load tests were made on these specimens at 1100°, 1300°, and 1500 F. Six specimens were tested at each stress level. Three creep-rupture machines were employed for tests at each temperature, two specimens being tested in each machine for each stress level. To minimize any effect due to local inhomogeneities along the length of the as-received bar the specimens were randomized before test. Each specimen was tested to rupture and in most cases an autographic extension-time curve was obtained. For very low stress levels at 1500°F no auto-

Jan 1, 1962

By F. A. Forward, V. N. Mackiw

The paper relates to the laboratory and pilot plant studies that have been carried out by Sherritt Gordon Mines Ltd., Metallurgical Research Div., in developing the ammonia pressure leach process for extracting copper, nickel, cobalt, and sulphur from high grade nickel concentrate produced from Lynn Lake ores, and describes in some detail the chemistry of the process. IT is well known'.' 2 that copper, nickel, cobalt, zinc, ferrous iron, and a number of other metals combine with ammonia in aqueous solution to form complex ions of the form [Me(NH,)x]"+. The stability and solubility of these ions depend on the concentration of the metal ion in solution, on the amount of NH, present, and on the amount and type of anions present, e.g., OH-, C0;-, NO,, C1-, SO;-. If ammonia is partially or completely removed from such solutions, for example by boiling, the soluble ammines tend to decompose and the metals precipitate as basic salts. These properties of metal ammines have found practical application in the commercial recovery of copper, nickel, and cobalt from a variety of ores. In an operation formerly conducted at Kennecott"" the mill tailing containing 0.80 pct Cu as copper carbonate was leached by percolation with an ammonia-ammonium carbonate solution to dissolve the copper as the ammine. At Calumet and Hecla0-" With a mill tailing containing about 0.4 pct Cu as metallic copper, it has been found necessary to aerate the ammonia-ammonium carbonate leach solutions between stages to oxidize the solubilized copper. At Nicarol"-" where nickel and cobalt occur as oxides and silicates which are not soluble in ammonia, the ore is heated to decompose silicates and to selectively reduce the nickel and cobalt to metal, leaving the iron as Fe,O,, and is then leached with ammonia-ammonium carbonate solution accompanied by aeration to dissolve nickel and cobalt as ammines. In each of these leaching operations, the anion present is CO, and the metals can be precipitated from the leach solution as oxide (copper) or hydroxide and carbonate (nickel and cobalt) by boil- ing off NH, and C02 both of which are recycled to treat a subsequent ore charge. The products-—oxide, hydroxide, or carbonate—can be treated by conventional smelting, calcining, or electrolytic methods to convert them to refined metals. Thus the procedures mentioned utilize the properties of the ammines to extract copper, nickel, and cobalt from nonsulphide ores and separate them from the leach solutions. These processes have the advantage that they can be carried out in closed vessels at atmospheric pressure, that the leaching solutions are specific for the desired metals, that the metals can be recovered as relatively pure compounds by the boiling operation, and that the NH, and CO, can be recycled. Also, as the leach solutions after boiling and filtration are substantially free of metals, NH,, and CO,, they can be discarded, thus facilitating the control of the water balance in the leaching circuit. When, as is the case with Sherritt Gordon concentrates, sulphides are treated with ammonia solution in the presence of oxygen, the leach solution contains SO,--, and an entirely different set of conditions is encountered. The ammine sulphates (of nickel, for example) are more soluble than the carbonates and can be only partially decomposed by boiling. The SO, and NH, in the solutions can not be recovered by boiling. Thus, despite the more favorable leaching conditions resulting from high solubility of the ammine sulphates, the recovery of metals, NH,, and SO, from the solutions must be effected by other means. Sherritt Gordon Nickel Concentrate Treatment The Ni-Cu-Co flotation concentrate produced at Lynn Lake'" contains 12 to 16 pct Ni, 1 to 2 pct Cu, 0.2 to 0.5 pct Co, 33 to 40 pct Fe, 28 to 34 pct S, 8 to 20 pct insoluble, and less than 0.02 oz per ton precious metals. The nickel is present chiefly as pentlandite, the copper as chalcopyrite, and the iron as pyrrhotite and pyrite. Most of the cobalt is thought to be present in pentlandite, although it is known that a small amount occurs as Co-Ni-pyrite.

Jan 1, 1956

By C. S. Smith

The solubility of tantalum at 1100°C is 0.025 pct in pure copper, 1.2 pct with 20 pct Ni, and 2.7 pct with -30 pct Ni. The solubility decreases with temperature, and the alloys are precipitation hardenable. A 79/20/1 Cu-Ni-Ta alloy reaches maximum hardness after aging at about 750°C and, if cold worked, does not recrystal-lize below that temperature. The alloys have good tensile properties at moderately elevated temperatures and, since they can be hot and cold worked nearly as easily as cupronickel, they are suggested for service at temperatures above the usual limits for copper alloys. MASSIVE tantalum dissolves extremely slowly in copper-rich alloys, and tantalum powder forms refractory surface layers which prevent its solution unless special precautions are taken. After many trials, an 80 pit recovery was consistently obtained by using a mixture of a finely divided tantalum powder with two or three times its volume of potassium tantalum fluoride (K2TaF1,-melting point about 750°C) poured in a slow stream directly on to the surface of the molten copper alloy, so that each particle would remain coated with flux and be immediately and individually wetted by the molten metal. A 50-50 nickel-tantalum alloy has a melting point of about 1400°C and small lumps of it slowly but satisfactorily dissolve in molten copper-nickel alloys as does commerical "ferro-tantalum." The high affinity of tantalum for carbon makes it necessary to avoid carbonaceous crucibles. COPPER-TANTALUM ALLOYS Castings with good shrinkage resulted when 0.1 pct Ta or more was added to copper melted under charcoal and cast in air. Copper in two-pound melts to which 0.1, 0.15, and 0.2 pct Ta had been added retained residual amounts, by analysis, of 0.025, 0.023 and 0.026 pct, respectively. The solubility of tantalum in molten copper at about 1200 °C is therefore about 0.025 pct. The electrical conductivity of these three samples was 100.48, 100.00 and 100.20 pct IACS in the annealed condition, and 98.24, 97.98 and 98.08 in the cold-drawn condition. Wires 0.080 in. in diam withstood from thirteen to nineteen reversed bends after annealing for 30 min in hydrogen at 850°C and the metal was therefore completely deoxidized. The annealing temperature corresponding to 50 pct loss of work hardness in 1 hr in a strip cold rolled 77 pct reduction was 175C, compared to 230°C for equivalent undeoxidized material. It is unusual for a decrease of recrystallization temperature to result from an alloying addition, and the tantalum probably combines with and removes some minor impurity in the original copper that restrained its recrystalliza- tion. No difficulty whatever was encountered in hot or cold rolling or drawing these ingots, and were tantalum not so expensive it would make an excellent deoxidizer for copper. COPPER-NIcKEL-TANTALUM ALLOYS Studies were made of the solubility of tantalum in a large number of copper-rich alloys, but of these significant solubilities were found only in the case of an iron alloy (which, however, separated into two immiscible liquid layers) and alloys of copper and nickel, in which tantalum is freely soluble and which were found to have interesting properties. The latter alloys are the subject of U. S. Patent No. 2,430,306 (1947). They are not at present available commercially. The Ternary Constitution Diagram—The copper-rich corner of the constitution diagram was constructed on the basis of microscopic examination of samples of various composition and treatments. The alloys were hot rolled to 0.25 in. from castings 0.625 in. thick, then annealed for 1 hr at 1000°C, quenched, cold rolled to 0.10 in. and finally annealed for various periods of time at temperatures between 800" and 1100°C in a nonoxidizing atmosphere and quenched. When the solid-solubility limit had been transgressed, particles of a compound believed to be Ni2Ta*were

Jan 1, 1960

By J. J. Gilman, T. A. Read

S URFACE effects in the cleavage of brittle crystals have been known for some oftime,1, 2 but our knowledge of surface effects in the plastic deformation of crystals is of relatively recent origin. In 1936, Roscoe9 eported his observations comparing oxidized and clean cadmium crystals. He found that cadmium crystals 1/2 mm in diam when covered by an oxide film less than 1200 atoms thick had critical shear stresses about twice those for clean crystals. This effect is anomalous in that it cannot be caused by the inherent shear strength of a normal oxide film. Either an interaction between the film and the crystal causes the large increase in strength, or the oxide film has an abnormal mechanical strength. Several authors since Roscoe have studied surface effects in crystal plasticity. A complete review of the literature will not be presented, but most of the results may be summarized as follows: 1—The critical shear stress for slip is increased markedly by very thin oxide films on zinc4 and cadmium crystals;" ' also, the entire stress-strain curve is raised by an oxide coating;3,4 both effects increase with decreasing specimen radius. 2—Surface active substances such as oleic acid destroy the hardening effect of oxide films.4, 6-8 The time required for the reduction of the hardening effect to begin increases with the viscosity of the solvent4~ 3—There has been dispute in recent years about the effects of surface active substances themselves on plastic deformation. Rehbinder, Lichtmann, and Maslenikov9 state that surface active substances reduce the strength of metal monocrystals markedly and increase the number of slip lines per unit length at a given strain. The effect varies with orientation for tin crystals being smallest for x, = 0, largest for x, = 45", and intermediate as x, + 90°.10 Several workers claim to have disproved the results of Rehbinder')"." but Rehbinder and Lichtmann9 have defended themselves and Masing with his collaborators" has confirmed some of Rehbinder's results. 4—Immersion of metal specimens in electrolytes such as sodium chloride reduces their strength12,13' and subsequent cathodic polarization increases the strength reduction. For noble metals, the strength reduction occurs for both anodic and cathodic polarization," increasing with increasing polarizing voltage. 5—Electroplated copper coatings effect the creep rates of zinc monocrystals. Poly crystals are not affected." 6—Immersion in dilute acid destroys the hardening effect of an hydroxide film. When an acid solution is applied to a steadily creeping crystal, the creep rate jumps to a very high value and then drops gradually to a steady value which is higher than the initial steady rate.*' " If the hydroxide film to be removed is very thin (about 100A thick), the creep rate does not change measurably, but a small rapid strain-increment is observed when dilute acid is applied." 7—Crystal plasticity depends on specimen dimensions. This was observed by Lichtmann and Ven-stroml" for monocrystalline tin and by Andrade and Kennedy" for polycrystalline lead. 8—Immerson of copper crystals in mercury has no effect on their stress-strain curves.'" 9—Bombardment of creeping cadmium crystals with a particles which penetrate the surface to only 0.005 mm may increase their creep rates as much as five times.''' The experiments which will be described in this paper pertain to: (a) the effect of a crystal's shape on its plastic deformation, and (b) the effect of metallic surface films on plastic deformation. For the work pertaining to (a), crystals of zinc, tin, and lead were grown which had various shapes. These were tested in simple tension. Macroscopic rotations about the tension axis which are unpredicted by the classical theory of slip were observed, and serrated edges were observed in tin crystals. For the work pertaining to (b), the effects of evaporated and electrodeposited films of copper, nickel, gold, and zinc on the mechanical properties of zinc, tin, and lead crystals were determined. The films influenced both slip and twinning of the crystals in creep and tensile tests. Experimental Details Chemically pure zinc, tin, and lead were used in this investigation. The zinc was 99.999t pct pure;

Jan 1, 1953

By W. A. Walker, G. W. Ostroot

An increase in the number of deep wells being drilled where extreme bottom-hole temperatures are encountered, and the anticipated drilling of wells where temperatures in the range of 500°F or higher may occur, has brought about a comprehensive investigation of cementing materials and of the techniques involved in their proper usage at these elevated temperatures. Included are developments in cements, retarders, weighting materials and other cement additives which make it possible to formulate a variety of compositions to help resolve the cementing problems of these extreme well conditions. The problems associated with the selection and testing of cements are discussed, and a resume of field results is included. Previous studies on strength retrogression indicate that caution should be exercised in the selection of a cementing composition for use in high-temperature wells. It now appears that, by the addition of silica flour as a stabilizing additive to certain cements, compositions covering a wide range of slurry densities can be designed to meet extreme well conditions without strength retrogression. Improvements in cement retarders make it possible to produce a four-hour thickening time at static temperatures up to 500°F. This temperature is considerably higher than conditions presently being encountered in drilling and completion work. Inert weighting materials, used to produce 20-lb/gal or heavier cement slurries, are reviewed. INTRODUCTION As the search for oil continues, the number of wells being drilled to depths of 10,000 ft or below has continued to increase over the past decade. During 1959, there were 50,893 wells drilled in the United States for a total of 209 million ft of hole. Of this total, approximately 4.45 per cent were drilled below 10,000 ft and 0.30 per cent below 15,000 ft, with a record depth of 25,340 ft recorded (Table 1). With the drilling of this record-shattering hole comes the question, "How deep can wells be drilled?".' Limitations with respect to equipment (surface and subsurface), cementing materials and techniques would cause some concern; however, a panel of experts have agreed that the U. S. oil industry has the equipment and know-how to drill a 50,000-ft hole. The greatest deterrent would be the 700°F temperature they would anticipate encountering at this depth. Although the number of wells drilled under what is generally considered hot-hole conditions (230°F and higher) has been relatively high, each well often presents distinctly individual problems and requires separate consideration and selection of special materials, techniques and equipment. At the present time there are over 30 additives and special cements available for use in "tailoring" a cementing composition for a specific well. It is not uncommon to use three or four different additives with the basic cement to produce the desired slurry characteristics. A generally accepted procedure is to make laboratory tests on selected cementing compositions prior to pumping them into wells deeper than 12,000 ft, or where static temperatures are in the range of 260°F or higher. In some instances, it is recommended that laboratory tests be made where less severe well conditions are encountered. Prior to selecting a cementing composition for a particular well, key personnel representing the operator, drilling contractor and service company should meet to determine that each individual is acquainted with every detail which might influence cementing results and to discuss the execution of the job. For cementing hot-holes, a number of factors must be evaluated before proceeding with the operation. Consideration must be given to the following. 1. Thickening time—The slurry must remain fluid for a sufficient length of time to allow displacement down the casing and up the annular space in a primary cementing job or, in squeeze-cementing, adequate thickening time for both the build-up of a satisfactory squeeze pressure and the reversing out of excess cement. The thickening time of standard retarded cements will often be too short under extreme temperature conditions, thus involving special formulations and requiring laboratory tests to determine the proper amount of additional retarder to obtain the desired pumpability.