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By J. D. Wood, J. T. Smith

ALLOYS with a dispersed second phase in a metallic matrix are generally much stronger than the matrix itself. Plastic deformation in dispersion-strengthened alloys is usually confined to the matrix phase when recovery processes are active, while in the absence of recovery both phases may yield.' The alloy system studied in the present research was WC-12 wt pct Co and consisted of noncoherent WC particles dispersed in the cobalt matrix. Some particle-to-particle contact existed but not enough to produce a continuous WC skeleton. The microstruc-ture of the WC particles was characterized by very straight edges, forming a trapezoidal shape in any plane of polish. Previous investigations with WC-Co alloys at room temperature have shown that fracture of the WC particles occurs in transverse rupture testing.' Room-temperature slip was reported for WC particles after indentation for hardness measurements.3 Elevated-temperature deformation of WC particles in a WC-12 pct Co alloy was suggested by recent electron microscope studies of specimens deformed at 900' to 1000°C.4 In highly deformed alloys, the WC edges were serrated in contrast to the usual straight or smooth appearance. WC-12 pct Co and WC-15 pct Co alloys have been previously studied under elevated-temperature com-pressive-creep conditions by the present authors. Electron microscope studies of two-stage replicas from deformed specimens showed no evidence of slip or fracture of the WC particles. These specimens were brought to temperature and allowed to equilibrate prior to the application of the creep load. It was believed that the load-application rate, a crosshead speed of 0.005 in. per min on an Instron universal testing machine, was sufficiently low that recovery within the cobalt matrix was sufficient to limit the deformation to this matrix. A series of experiments was performed to evaluate the influence of loading rate on the deformation of WC-Co alloys. A WC-12 pct Co alloy was selected for these determinations. The average WC particle size was 4.45 p with an average linear separation between particles of 0.59 p. The selected temperature was 800°C and was monitored with a Chromel-Alumel thermocouple attached to the specimen. Testing was conducted in an argon-atmosphere chamber to prevent oxidation of the WC-Co specimens. This chamber was mounted on an Instron universal testing machine equipped to apply the load at a fixed rate. Each specimen was loaded to 110,000 psi compression stress at 0.05 and 0.5 in. per min. The loading rate was monitored prior to insertion of the test chamber and was found to be almost precisely the nominal rate selected. The specimens were raised to temperature and held to equilibrate with the surroundings, and then the load was applied and held for 4 hr to duplicate the exposure time utilized for the creep specimens. The time to reach full load at a crosshead speed of 0.005 in. per min was some 500 sec and was reduced to 50 and 5 sec as the loading rate was increased to 0.05 and 0.5 in, per min, respectively. The model developed by Ansell,' when recovery processes do not occur, considers that fracture or deformation of the dispersed particles is necessary to relieve back stresses on dislocation sources and allow dislocations piled up against particles to sweep out in the matrix to cause plastic deformation; he further states that, even at elevated temperatures, the dispersed-particle deformation is necessary for yielding in the absence of recovery. For the case of straight dislocation segments piled up against a straight barrier, such as the straight-sided WC particles, the shear stress, 7 exerted on a particle is: where h is the spacing between particles (0.59 p), a is the applied stress (110,000 psi), p, is the shear modulus of the matrix (6.7 X lo6 psi at 80O°C), and b is the Burgers vector of the matrix dislocation. From Eq. [I], the shear stress, 7, exerted on the WC particles when no recovery occurs is of the order of 6 X lo6 psi at 800°C. The limiting stress, F, that will

Jan 1, 1969

By W. W. Eckles

This paper is presented as a suniniary report of the use of well gas composition correlations obtained from mass spectrometer recordings as a means of identification and determination of reservoir continuity. Conventional methods for detecting composition differences are en-pensive, elaborate, and difficult to obtain. This excludes the use of extensive composition data for most applications. During recent years the mass spec-trometer has come into general use as an analytical tool in petroleum refineries. The use of mass spectrometer composition patterns ir~ characterizing or "finger-printing" the produced gas from a reservoir, presents a novel method for correlatitlg gas samples from well to well. The mass spectrometer provides a trace similar to an electric log. having peaks which represent the abuntlnnce of certaitz hydrocarbons in the well gas sample. Without going further into the detailed analysis the idea has been advanced that these traces or patterns could be used as a means of identifying a particular natural gas. This theory has proven to be essentially correct. The mass spectrometer pattern method is simple and cheap as COi?7pared to other standard methods. It greatly facilitates the solution of reservoir and geological problems in which correlation of well gas Com-positions is a factor. Specific field applications have been made. This paper concerns the results obtained in 465 individual gas analyses from 35 fields and 77 res- ervoirs. In a number of cases it has been found that such data have been extremely valuable in the determination of reservoir continuity. In at least one case, the method was a valuable contribution in tracing a reservoir from sand to sand in a coinplex fnulted field involving n11rnerous gas reservoirs. Field applications are presented to illustrate the possibilities of the method at the present stage of developnient and to stimulate the ernployrnent of this new approach by geologists and petroleum engineers in the industry. INTRODUCTION Identification of producing horizons and the determination of reservoir continuity are often a problem in those areas where dome structures and highly faulted sands are encountered. To complicate the picture further, there may be numerous sands, one on top of the other which dip and diverge in different directions. Even though it may be possible to develop some solutions to the preceding problems on the basis of the geological and reservoir data on hand. it is readily recognized that substantiating data based on independent methods would he extremely valuable. It has been found through field studies that correiation of well gas composition can be used to advantage in the geological and engineering study of a complex reservoir identification problem. METHOD FOR COMPARISON AND CORRELATION OF WELL GAS SAMPLES Since a large number of well gas samples are required in a gas identification or reservoir con- tinuity problem, it is necessary that a method of analysis be employed which can detect differences in composition readily and inexpensively. Although detection would be possible by means of low temperature fractional distillation (POD) analpsis, the method requires a relatively large sample and is comparatively slow and expensive to run. The mass spectrometer affords an inexpensive and precise analysis of small well gas samples taken at the surface which are about 1/300 as large as a POD sample. These samples can be obtained by regular field personnei and shipped to the spectrometer for analysis. It is, therefore, a practical approach to the problem. A typical record from the mass spectrometer is illustrated in Fig. I. The peaks on the record represent the abundance of ions produced from the different hydrocarbon molecules making up the gas sam-ple. In well gas comparisons only five peaks are employed, since the other peaks are formed from the same gas molecules and furnish no additional information. The five peaks used represent the abundance of methane, propane, ethane, butane and heavier, and oxygen. They are referred to respectively as the 16, 29, 30, 43, and 32 peaks. The oxygen peak is used to correct for air content in

Jan 1, 1958

By W. E. Howlett, R. L. Slobod

Jan 1, 1965

By J. S., Douglas

IN 1880, the late James Douglas, LL.D., was superintendent of the Chemical Copper Co., operating the Hunt & Douglas process for the treatment of the siliceous copper ores of the Jones mine at Phoenixville, Pa. In conjunction with this a refining furnace was operated. After correspondence with Ben Williams, of Bisbee Ariz., who was commencing the mining and smelting of the Copper Queen ore. Dr. Douglas decided-to make his first journey to Arizona. While making plans for the trip lie visited a friend in Philadelphia, a .Mr. Logan, a prominent lawyer. He, through Professor Treadwell, brother of the discoverer of the great Treadwell gold mine of Douglas Island, Alaska, had an option with his min¬ing partner, Mr. Lenig, on the Eureka and Wade Hampton claims near the headwaters of the Verde River in Yavapai County, Arizona. These two claims, later the source of practically all of the wealth of the Jerome district, were then owned by C. P. Head and Co. Prescott merchants. Hugo Richards, banker, and formerly sutler at Camp Apache; Rod and John McKinnon prospectors and Nova Scotians; and perhaps one or two 'other owners of minor interests. Mr. Logan, hearing of Dr. Douglas intended trip to Bisbee, asked him to go first to see the two mining prospects, and this, Douglas, in need of the $500 fee, agreed to do as the time on the option on the claims was running short.

Jan 1, 1935

By Hermsdorf, Richard P. E.

AMONG the newer lead smelting and refining plants of the country is that of the United States Metals Refining Co., at Carteret, N. J. Not only is the technical practice here modern and efficient, but exceptional care is taken to assure healthful operating conditions for the employees. One large building houses the entire department, which includes, as its chief units, one reverberatory smelting furnace, one blast furnace, two softening furnaces, one utility reverberatory furnace, two Faber du Faur retorts, five 120-ton kettles, three 60-ton antimouial lead kettles, and one 5-ton kettle for experimental purposes. There is a 10-ton overhead crane in the kettle aisle, which is also used for shipping lead. Other auxiliary equipment is referred to in the outline of operations. A general plan of the plant is shown in Fig. 2, and a general view in Fig. 1. Primary smelting operations are conducted in a 30 by 11-ft. reverberatory furnace with bottom, roof, and upper walls of clay brick, magnesite

Jan 1, 1934

By J. Chipman, H. R. Larson

The data previously reported for the quantity as a function of oxygen pressure at 1550°C have been used to compute the activities of Fe, FeO, Fe2O3, and COO in slags of the ternary system. Activities of the first three have been obtained also for two quasi-ternaries involving fixed CaO:SiO2 ratios. IN a previous paper1' the authors reported the results of an investigation into the effects of oxygen pressure on the composition of various simple slags analogous to some of those which are found in steel-making practice. The ratio of ferric iron to total iron was studied at 1550°C in iron oxide melts to which lime, magnesia, lime-plus-silica, and other oxides were added. The oxygen pressures involved those of air, carbon dioxide, and carbon dioxide plus carbon monoxide in several proportions. Although very low oxygen pressures could not be used, the slag-metal equilibrium studies of Fetters and Chip-man' permitted extending the results to slags in equilibrium with iron. For the ternary system CaO-FeO-Fe,O, the oxygen pressure-composition relationship has been determined from zero percent lime to lime saturation over an oxygen pressure range from air to that represented by equilibrium with liquid iron. Lime and silica were added to iron oxide in the ratios 0.54, 1.306, and 2.235 to form three quasi-ternary systems which were also studied over the entire region of liquid melts at 1550°C. Ternary Gibbs-Duhem Equation Wagner" has developed a method by which the activities of two components of a ternary system can be calculated if' the activity of the third component is known throughout the composition range being considered. The fundamental form of the Gibbs-Duhem equation for ternary systems is N, d In a, + N2 d In a, + N3 d In a3 — 0. Wagner has developed a usable form of this equation by introducing the term y = N3/(N3 + N1) and rearranging the equation to give (? In a1/?N2)x = -y/(1-N2)2 (? In a2/?y) x2 N2/1-N2 (? In a2/?N2) y To apply this equation to the slag system CaO-FeO-Fe,O,, the activity of one of these components must be known. However, the only activity which is known from the experimental data is the activity or pressure of oxygen in the gas phase with which the slag is in equilibrium. The activity of oxygen in the slag can be defined as the square root of the oxygen pressure. In order to use oxygen as one component, the composition of the slags must be converted to the basis Fe-O-CaO. Oxygen, exclu-sive of that contained in CaO, then becomes com-ponent 2 in Eq. 1, Fe is selected as 1 and CaO as 3. Then y = Ncao/(Ncao + NFe). Eq. 1 then becomes (? In a fe/? Nx) y = -y/(1-No)2 (? In ao/?y) xp - No/ 1- No (? In ao/ O No) y. In order to evaluate Eq. 2, the boundary conditions must be known. The obvious choice of a standard state for iron is to assign slags in equilibrium with iron an activity of one. Then In a,., or log aFe, which is substituted for convenience, is determined by integrating along a line of constant y from the slag in equilibrium with iron to the composition at which log a,, is to be determined. Mathematically this can be expressed as log aFe (Nlog y) = log a'Fe + (? log a Fe/? No) dNo where the primes indicate equilibrium with liquid iron and a'Fe = 1. When Eq. 3 is integrated along a line of constant y, the following is obtained: log a Fe (No, y) = - y No ?/ No ? y ( log ao/(1- No)2) No d No - No/1-No d log ao. Lime-Iron Oxide Slags Iron Activity: The oxygen pressure of lime-iron oxide slags is shown in Fig. 1 as a function of j, de-fined as j = Fe+++/(Fe++ + Fe +++) for various constant mol percentages of lime. The j values at 0 to 60 pct lime were then determined for oxygen pressures of 1, 10-1, etc. to 10-" atm. For purposes of calculation, the line for zero percent lime was extrapolated to

Jan 1, 1955

By P. R. Sperry, P. A. Beck, J. Towers

Dahl and Pawlek1 found that electrolytic copper develops extremely coarse grains at 1000°C after about 90 pct reduction by rolling. This coarsening occurs only under conditions of penultimate grain size, deformation, and alloying which lead to the "cube" recrystallization texture.l,2,3,5 The peculiar angular shapes and straight grain boundaries of the coarse grains were noted by several investigator.1,4,5 On the other hand, coarsening in Fe-containing aluminum or in AI-Mn alloys8 does not depend on a "cube" (or any well developed) recrystallization texture. It is true that increasing deformation by rolling, and, therefore, an increasingly well developed re-crystallization texture, are associated with decreasing incubation periods of coarsening.6-7-8 Nevertheless, coarsening readily develops in aluminum even after only 30 pct reduction by rolling, where the recrystallization texture is very weak.6,8 Also, coarsening was observed by Jeffries9 many years ago in sintered thoriated tungsten, which presumably has no preferred orientation. In all these cases coarsening is associated with grain growth inhibition by a dispersed second phase.8,9 The annealing temperature has to be suficiently high to overcome the inhibition at a few locations. But if it is too high, growth starts at many points, and the resulting grain size becomes much smaller.9 Normally, the coarse grains are more or less equiaxed, and the boundaries have a typical ragged appearance.6.8 Cook and Macquarie4 demonstrated that, in addition to the texture-dependent coarsening previously found at 1000°C,l electrolytic tough pitch copper may also coarsen at 800°C after 50 pct reduction by cross rolling. The coarse grains formed under such conditions have rounded shapes and ragged boundaries, like those in aluminum. When the annealing temperature is higher, the tendency for their formation decreases. All these observations suggest that the coarsening at 800°C is associated with inhibition by a second phase. Actually, coarsening at 800°C after 50 pct reduction by cross rolling was observed only in tough pitch copper,4 which contains Cu2O particles. On the other hand, the texture-dependent 1000°C coarsening occurs in both tough pitch and oxygen-free copper;4 it does not appear to depend on the presence of a dispersed second phase. However, the interpretation of the 800°C coarsening in Cu after 50 pct rolling as an inhibition-dependent process, similar to the coarsening in A1-Mn alloys, is somewhat weakened by the fact that this coarsening was reported4 to occur only after cross rolling, and not after straight rolling. It was, therefore, decided to re-examine this question. A 1 in. diam electrolytic tough pitch copper rod, No. 2 hard drawn, was annealed for 20 min at 700°C, rolled to 0.5 in., annealed 10 min at 700°C, and straight rolled to 0.064 in. It was then given a penultimate anneal of 20 min at 500°C and it was cut into four sections, which were given final reductions by straight rolling as follows: A 30 pct reduction of area B 50 pct reduction of area C 70 pct reduction of area D 90 pct reduction of area Specimens cut from the four sections were finally annealed at 800°C in an oxidizing atmosphere. Strip A remained fine grained up to 10 hr, but the specimen annealed 12 hr consisted of only 2 large grains. Strip B had a few scattered large (1/2 to 3/4 mm) grains after 1 min, although the balance of the specimen consisted of fine grains of about 0.02 mm. After 5 min there were several 10 to 15 mm grains present, and after 1 hr strip B was completely coarsened. The coarse grains had the same characteristics (see Fig 1) as those obtained by Cook and Macquarie at 800°C after cross rolling. Strip C had several grains of 0.05 to 1 mm after 1 min, but it was still largely fine grained after 12 hr. After 48 hr it consisted entirely of grains of about 0.5 to 4 mm, with an extraordinarily large number of twin bands. Strip D remained com- pletely fine grained after 4 hr at 800°C. These results indicate that, in the deformation range of 30 to 70 pct reduction, the incubation period for coarsening as well as the rate of growth and the final size of the coarse grains decreases with increasing deformation. Similar

Jan 1, 1950

By William R. Chedsey

ALTHOUGH few changes can be reported in educational methods at the mineral technology schools during 1940, other events have taken place of direct interest to, and that will have a profound effect upon, mineral industry education on the United States. In the erection of new buildings, some notable changes have been made on the campus, The U.S. Bureau of Mines Experiment Station at the University of Utah, Salt Lake City, was dedicated on May 20. The construction and equipment of this building cost between $300,000 and $400,000. It will be the general headquarters for the metallurgical division of the Bureau and will have a pronounced effect upon the mineral industries research program at the University.

Jan 1, 1941

By John D. Weibmer

A strong step towards acknowledgement of mining as a skilled trade has been taken by the Tintic Division of Kennecott Copper Corp. In May 2 ceremonies, a pilot program was signed into force which, possibly for the first time in US history, recognizes mining as an apprentice trade and grants journeyman status to those who complete the prescribed training sequence. The apprenticeship standards, hammered out after months of consultation between management, labor, and government, are presently limited only to underground hard-rock mining and only within Kennecott's Tintic Division, but the architects of the program are confident that the concept will spread-first to other underground operations where miners desire documented evidence of their special skills, then to open-pit and soft-rock operations.

Jan 6, 1977

By Potts N. L., Eilerts C. K., Summer E. F.

The second-order, nonlinear, partial- dillerential equation representing the transient radial flow of gas - condensate fluids in reservoirs has been integmted by using finite -dillerence equations and electronic computers. Effect was given to pressure -dependent permeability, viscosity, and compressibility and to distance-dependent permeability. The influence of a second-degree velocity term in the Darcy equation was investigated. Implicit methods were used and practical, convergent solutions were obtained with material balance to less than 6 x 10-4 for recovery of one- hall the reserve at constant flow rate. Integration results provide the productive period of a reservoir for a given constant rate and the fraction of the fluid initially in place that can be recovered in that period. The properties of a lean and a rich fluid are represented in a set of integrations designed to demonstrate the effect of different constant-recovery rates and significant variables emphasized one at a time. INTRODUCTION A method is needed for computing the availability or reserve of a gas in a formation, making use of all technical and engineering information that is pertinent. As a step in this direction, a program for computing transient linear flow1 was developed that utilizes principles of earlier finite-difference computing for a similar purpose2-4 and gives effect to significant pressure- and distance-dependent properties of reservoirs and their contents. The present paper pertains to transient, radial flow of gas-condensate fluids. In the partial-differential equation of flow, effect is given to the variables viscosity and compressibility,5 and also to permeability. Included in the study is the quadratic form of the Darcy equation of flow that has been the subject of field tests6 and that has been applied to a gas with constant properties in transient flow computing. 7 Inclusion in the differential equation of variable coefficients to represent properties greatly complicates the higher derivatives of the equation. Because it was impractical to make a required improvement in certain of the finite-difference derivatives by Taylor-series expansion, five-term derivatives were used in the implicit computing. Related methods were developed that can improve general facility in manipulating finite-difference equations. BASIC EQUATIONS PARTIAL DIFFERENTIAL EQUATION The basic partial differential equation for transient radial flow in the direction of decreasing radius is l_ a(vpr) _? .......(1) r gr at where r = radius, L v = apparent velocity aIong radius, L/t p = density, M/L3 t = time, t Ø = porosity, dimensionless. In this equation, only the porosity is regarded constant. The density is where p = pressure, m/Lt2 M = molecular weight of fluid, m T = constant temperature of fluid, T z(p) = pressure-dependent compressibility factor, dimensionless R = gas Constant, mL2/t2t. The physical boundary of the reservoir is at re, and the producing boundary is defined to be at tq with rq = re /200. By definition,

Jan 1, 1966

By A. H. Cottrel

SINCE metallurgy exists to provide strong, tough, engineering materials it must inevitably be perpetually concerned with the problem of brittle-ness. The steel-making industry was created because chemically unrefined iron is brittle. The steel-working and alloy-steel industries exist partly because hardness and ductility are mutually exclusive qualities in structurally unrefined steel. Yet, in spite of these immense industries, the brittle failure of steel ships hulls, bridges, pressure vessels, and pipelines, is still a contemporary problem. In fact, brittleness is a normal property of most solids, including metals and alloys, at low temperatures. Only face-centred cubic metals are commonly ductile at the lowest temperatures, and even here exceptions are known.'," Expensive are-melting processes. under inert gases or vacuum, have been developed to overcome brittleness in titanium, zirconium, columbium. and molybdenum. Polycrystal-line zinc, magnesium, and uranium have little ductility at room temperature. Beryllium, chromium, and tungsten have even less; similarly for antimony, bismuth, germanium, silicon, intermetallic compounds, and metallic carbides, nitrides, silicides, and borides. Oxides and other ceramics would be ideal creep-resistant materials but for their extreme brittleness when cold, and great efforts have been made to overcome this problem by mixing ceramics with metals. In fact, the traditional use of the name metallurgy for what is really the science of engineering materials is a recognition that most non-metallic solids have so far been precluded, by their extreme brittleness, from use as major structural materials in mechanical engineering. Most of what we know about brittleness in metals has come from studies of structural steel, and it is this material that shall mainly be considered. The things we have learned from it have a wider application, at least to other body-centred cubic transition metals, although the extent to which similar ideas can be applied to hexagonal metals and other materials is not yet clear. Fracture and Plastic Deformation The theoretical breaking strength of an ideal solid, about E/10 where E is Young's modulus, has been approached reasonably closely in experiments on fibres. But large specimens break at much lower stresses. For example, brittle cracks in large structural steel assemblies have been observed to spread catastrophically at speeds of 6000 ft per sec under stresses of about 10,000 psi, i.e. about E/3000. There are two sources of weakness; stress concentrations, and chemical agents in grain boundaries or on crack faces that lower the surface energy y of the material. Although we shall deal mainly only with the first of these, much of the discussion is also applicable to the second when different values of y are used. Griffith's well-known formula p = [ 2Ey/p(1 - v2)c]1/2 [1] gives the smallest tensile stress p able to propagate an atomically sharp surface crack of length c or interior crack of length 2c through a thick plate (compared with c) of elastically isotropic material of

Jan 1, 1959

C. M. Bishop (Aircraft Div., A. V. Roe Cnnada Ltd., Toronto)—Reference is made to the coherency hardening of ß due to ß' and also to the hardening of ß due to a coherent precipitation of the a phase. What is meant by the term "ß'?" Would not the coherency hardening of ß during cooling be due to precipitation of a sub-microscopic dispersion of a, as is the case during reheating a ß-quenched alloy? I feel that these two phenomena are essentially the same mechanism, i.e., precipitation of a from unstable or supersaturated ß. From the text, it would appear that Tables III and IV are intended to show the effect of the martensite transformation on the properties of Ti-Mn alloys. Referring to Table 111, the ß-quenched alloys would contain a greater proportion of retained ß than would the equilibrated alloy. Could not the increase in properties be attributed to this, along with possible coherency hardening of the ß? In the same way, the structures referred to in Table IV contain ß from which they could derive their hardness. It would appear to be difficult, if not impossible, to assess the role of the martensite transformation from the data contained in these tables. F. C. Holden, H. R. Ogden, and R. I. Jaffee (authors' reply)—In reply to the first question raised by Mr. Bishop, the term "ß"' designates an apparent retained-p structure which is harder than it is after the most rapid cooling from the ß field. It is thought to be the result of a coherent precipitation of a. The mechanism by which hardening takes place is undoubtedly the same as that involved in the early stages of aging of a re-tained-p structure, and we are in agreement with Mr. Bishop on this point. The term "ß'" to designate the anomalously hard retained-ß structure was agreed upon at a symposium in 1953." The authors do not agree that the quantity of retained ß and/or ß' in the ß-quenched specimens of Table III is necessarily greater than that of the 600°C equilibrated specimens. The ß-quenched specimen has a structure which appears to be completely a' formed by the martensite reaction. It is recognized that there could be some retained ß (or ß') between the a plates if the M, temperature is below room temperature. This is unlikely for the 3.4 pct Mn alloy where the M, is known to be at about 640°C. Therefore, the increase in strength cannot be attributed to the presence of a greater quantity of coherency-hardened ß, but, rather, the increase in strength must be attributed to the transformation products formed, e.g., a' (martensiti-cally formed a). Concerning the data in Table IV, we agree that it is difficult to estimate the role of the martensite reaction, since the 6.4 pct Mn alloy was retained ß in the quenched condition.

Jan 1, 1955

By Richard R. Zupp, David A. Stevenson

A theoretical treatment of austenitic Fe-C-Z solutions (where Z represents a substitutional solute) was developed by extending a solution model which was previously applied to austenitic Fe-C alloys. The model assumes that each carbon atom in the solution excludes H interstitial sites from occupancy by other carbon atoms rmd thai otherwise the assembly of carbon atoms is equivalent to a modified Langmuir lattice gas. The results of the treatment provide theoretical justification for a phenomenological expression previously developed to represent the influence of substitutional solutes on the thernlodynamics of carbon in austenitic Fe-C-Z alloys. The expression is: where Kz is a parameter which characterizes the Fe-C-Z system at a given temperature and Xz and xj are the concentrations of substitutional solute and interstitial carbon expressed as atom ratios. The parameter, Kz, was associated with the differ- A number of investigators have measured the influence of various substitutional solutes on the carbon activity in austenitic Fe-C-Z alloys, where Z represents any substitutional solute. Experimental information is available for alloys containing the following substitutional solutes: Mn,' Si,' Ni,2-5 CO,'-', Cr,5, 9-11 CU,4 Al,= v,11, 12 and Ti.13 This information is essential for studying many phenomena in Fe-C-Z alloys, particularly diffusion and diffusion-controlled phase transformations For example the pffprts nf various ence between the interaction energies between carbon atorns and host-lattice atoms in the Fe-C-Z solution and the corresponding energies for the Fe-C solution. The carbon activity in a ternary Fe-C-Z alloy, ac, was obtained in terms of the composition of the alloy and the single parameter, Kz: where y° is the Henry's law constant for the binary Fe-C solution. Both y° and H crm be determined from the data for Fe-C solutions, without recourse to any experimental data for Fe-C-Z solutions. The theoretical treatment was extended to alloys containing several substitutional solutes. By neglecting interactions between the substitutional solutes, an expression for the carbon activity in multicomponent alloys was obtained in terms of the Kz values characterizing ternary Fe-C-Z systems. substitutional solutes on the rate of pearlite growth is explained14 on the basis of their effects on carbon diffusion, which result from their influences on the carbon activity in the solid phases. In addition, recent studies have emphasized the dependence of the pro-eutectoid ferrite transformation15 and other diffusional processes in austenite16 on the influence of substitutional solutes on the carbon activity in austenite. They have indicated the need for suitable analytical expressions to represent such influences. The present authors previously developed general phenomenological expressions to represent the equilibrium partitioning of carbon between Fe-C alloys and Fe-C-Z alloys.'2 The expressions are applicable to all the austenitic Fe-C-Z alloys for which experimental data are available, and are sufficiently simple to recommend their general application. For alloys in

Jan 1, 1969

By V. R. Garfias

EVEN before the arrival of Cortes in 1519, the history of Mexico was closely linked to that of its mineral production; the mining activities of the Aztecs being thus described by Clavigero, one of the early historians: In the mountains of Anahuac there abound in veins all kinds of metals, and infinite variety of other fossiliferous products. The Aztecs obtained gold from the countries of the Cohuijques, of the Mixtecas, of the Zapotecas and from various other places. They gathered that precious metal usually in grain, from the sands of the rivers, keeping a certain portion for the crown. They obtained silver from the mines of Taxco, of Tzumpango and others; however, this metal was not as esteemed by them as by other neighboring nations. They had two kinds of copper; one hard, which they used instead of iron to make sickles, lances and all kinds of rural and military instruments and another, soft, with which they made kettles, goblets and other receptacles. This metal abounds mainly in the province of Zacatula and in that of Cohuijques, which at present is in the kingdom of Michoacan. They obtained tin from the mines of Taxco and lead from those of Izmiquilpan, located in the country of the Otomiea. From the tin they made money, and the lead, we know they sold in the market places, but we ignore the uses to which they applied it. They also had mines of iron in Tlaxcala, in Taxco and in other places; but they either had not discovered them, or did not know how to make use of the metal they contained. In Chilapa there were mines of mercury, and in other places there were mines of sulphur, alum, vitriol, cinnabar, cohre and of a white earth which they esteemed highly. As regards mercury and vitriol, we do not know of what use they were to them; the other metals they used in paintings and dyes. There were then great abundance of amber and asphalt-bitu- men of Judea--on the coast of the seas and many cities of that territory paid tribute of one and the other to the kingdom of Mexico. They mounted the amber in gold, and it was only used as ornament and display by them. With the asphalt they made certain perfumes. The following extracts dealing with pre-Colonial gold are taken from the famous "Cartas de Relación" to Charles V, in which Cortea described the progress of the Conquest: . . . here there came two chiefs who owned lands in the valley; one four leagues down the valley and the other two leagues up the valley and they gave me some gold necklaces of little value and seven or eight girls. . . . . . . that day I left the City of Churultecal and traveled four leagues to some villages of the City of Cuasucingo where the natives received me very well and gave me women and pieces of gold. . . .

Jan 1, 1937

By Thomas A. Read

The internal friction of single crystals of metals is affected markedly by a variety of factors, which, according to the literature, are without influence on the internal friction of polycrystalline metals at strain amplitudes less than 10-5. For example, the internal friction of a suitably oriented zinc single crystal is multiplied by a factor of ioo when the vibration amplitude is slightly increased, and the internal friction of a copper crystal is markedly increased by the application of a stress of 30 lb. per sq. in. for one minute. In this paper it will be shown that these, together with the other phenomena associated with the damping of mechanical vibrations in metal single crystals, can be satisfactorily explained, at least qualitatively, in terms of an atomic mechanism that has been proposed for quite another purpose—that of accounting for slip in metals. The extraordinary agreement between the observed behavior of the metal single crystals and the behavior that could be predicted on the basis of this atomic mechanism leads to the belief that the latter is probably essentially correct. If this is so, it follows that valuable information about the slip process in metals may be gained from measurements of internal friction on single crystals. In this paper the results of the only study to date on the damping of mechanical vibrations in single crystals of metals mill be presented. Experimental Procedure The internal friction of a material is a measure of the rate at which the energy of mechanical vibrations is converted into heat. The measure of internal friction used in the present paper is the decrcment A, which is defined as the fraction of the energy of vibration lost per half cycle. The measurements described here were made for longitudinal oscillations at a frequency of 33.5 kc. for the copper crystals, and at a frequency of 39 kc. for the zinc crystals. The method used is that of the composite piezoelectric oscillator. The single-crystal specimens are in the form of circular cylinders M in. in diameter and I 1/2 to 21/2 in. long. The crystal rods are waxed in a V-block and sawed to the proper length with a fine jeweler's saw. The ends of the specimen are ground flat, and one end is cemented with a very thin film of beeswax and rosin to the end of an x-cut quartz rod which is 1/4 in. in diameter and 3 in. long. The length of the specimen rod is determined by the condition that its natural frequency for longitudinal vibration be the same as that of the quartz rod, within 0.5 per cent. The lengths of the specimens are thus different for different orientations of the cylinder axis with respect to the crystallographic axes. When the specimen and quartz rods are thus matched in frequency the adhesive is at a node of stress, and it has been shown experimentally that under these circumstances

Jan 1, 1941

By Thomas A. Read

The internal friction of single crystals of metals is affected markedly by a variety of factors, which, according to the literature, are without influence on the internal friction of polycrystalline metals at strain amplitudes less than 10-5. For example, the internal friction of a suitably oriented zinc single crystal is multiplied by a factor of ioo when the vibration amplitude is slightly increased, and the internal friction of a copper crystal is markedly increased by the application of a stress of 30 lb. per sq. in. for one minute. In this paper it will be shown that these, together with the other phenomena associated with the damping of mechanical vibrations in metal single crystals, can be satisfactorily explained, at least qualitatively, in terms of an atomic mechanism that has been proposed for quite another purpose—that of accounting for slip in metals. The extraordinary agreement between the observed behavior of the metal single crystals and the behavior that could be predicted on the basis of this atomic mechanism leads to the belief that the latter is probably essentially correct. If this is so, it follows that valuable information about the slip process in metals may be gained from measurements of internal friction on single crystals. In this paper the results of the only study to date on the damping of mechanical vibrations in single crystals of metals mill be presented. Experimental Procedure The internal friction of a material is a measure of the rate at which the energy of mechanical vibrations is converted into heat. The measure of internal friction used in the present paper is the decrcment A, which is defined as the fraction of the energy of vibration lost per half cycle. The measurements described here were made for longitudinal oscillations at a frequency of 33.5 kc. for the copper crystals, and at a frequency of 39 kc. for the zinc crystals. The method used is that of the composite piezoelectric oscillator. The single-crystal specimens are in the form of circular cylinders M in. in diameter and I 1/2 to 21/2 in. long. The crystal rods are waxed in a V-block and sawed to the proper length with a fine jeweler's saw. The ends of the specimen are ground flat, and one end is cemented with a very thin film of beeswax and rosin to the end of an x-cut quartz rod which is 1/4 in. in diameter and 3 in. long. The length of the specimen rod is determined by the condition that its natural frequency for longitudinal vibration be the same as that of the quartz rod, within 0.5 per cent. The lengths of the specimens are thus different for different orientations of the cylinder axis with respect to the crystallographic axes. When the specimen and quartz rods are thus matched in frequency the adhesive is at a node of stress, and it has been shown experimentally that under these circumstances

Jan 1, 1941

By J. L. Lamon, W. Crafts

Study with the electron microscope of the carbides in vanadium-chromium-molybdenum steels after tempering up to 1000 hr at 600 teelsto 1400°F confirmed that alloy carbides are formed at the secondary hardening temperature by decomposition of the plate-like iron carbides. It was also demonstrated that vanadium carbide persists as much smaller particles than do chromium- or molybdenum-bearing carbides. Conditions conducive to the formation of fine vanadium carbides are indicated to be favorable for high temperature strength. IN order to determine the effects of long exposure at high temperatures on vanadium-bearing steels, a survey has been made of their hardness and carbide structure. The behavior of carbides in the tempering of martensite was studied by X ray diffraction and electron microscope examination of electrolytically extracted residues in greater detail than in an earlier investigation.' A group of steels with about 0.25 pct carbon and containing chromium up to 5 pct, molybdenum up to 1 pct, and vanadium up to 1 pct was either quenched or annealed and then tempered for peri- ods of up to 1000 hr at 600° to 1400°F. Their hardness after tempering agreed with the Hollomon and Jaffe2 relation of equivalency of time and temperature, and with the degree of secondary hardening predicted from composition. Further, the appearance of the carbides indicated that the time and temperature equivalency relation was also applicable to the degree of carbide development. The mechanism of the tempering of martensite was demonstrated more clearly than in the earlier study. It was confirmed that carbides develop from martensite as poorly defined plates of the Fe2C type carbide followed by thickening of the plates with a transition to Fe,C. Finally the plate structure deteriorates into a lacy mass from which alloy carbides emerge as chunky particles that grow slowly with further increase of tempering temperature. Vanadium carbide derived from tempered martensite was found in characteristically small particles that tended to grow very slowly. The addition of chromium or molybdenum to a steel with predominant vanadium carbide tended to introduce carbides of a chromium- or molybdenum-bearing type having a somewhat larger particle size. The Cr,C, type carbide particles were larger than those of the V,C, type and somewhat smaller than the carbides in steel containing both Cr,C and M,C. Similar observations were made of the carbides in the pear lite of annealed steels. It appeared that high temperature properties would be benefited by the retention of the fine carbide particles that result from a balance of composition and heat-treatment designed to produce a maximum amount of vanadium carbide. Procedure: The steels used in the investigation were made with a base of Armco iron in a high-frequency furnace at the Union Carbide and Carbon Research Labs., Inc. The steels were deoxidized

Jan 1, 1951

By L. D. Jaffe, D. C. Buffum, F. L. Carr, M. Goldman

IN an earlier investigation by Jaffe and Buffum,' a time-temperature transformation diagram for isothermal temper embrittlement of an SAE 3140 steel showed some suggestion of two overlapping "noses." These authors mentioned, as a possible explanation, that their specimens embrittled at 525°C were slightly softer than expected, perhaps because of some slight unknown variation in tempering treatment. Because a double nose would be of considerable importance in elucidating the mechanism of the temper brittleness reaction,' it appeared worthwhile to redetermine the portion of the diagram near 525°C. The same heat of SAE 3140 steel was used as in the previous work, and experimental procedures were the same.' Blanks were austenitized 1 hr at 900°C, water-quenched, tempered 1 hr at 675°C, and water-quenched. Groups were then embrittled isothermally at the temperatures and for the times shown in Table I, and water-quenched. For comparison, one series was tempered various times at 675°C, water-quenched, and given no further treatment. V-notched Charpy impact specimens machined from the blanks were tested over a range of temperatures. Energy absorption and percentage of fibrous fracture were plotted vs test temperature for each group, utilizing both the new data and those for the same treatment obtained in the earlier work,' except that the earlier data at 525°C were discarded.* are given in Table I. After a small correction (-0.28°C/BHN) for the variation of hardness" from Rockwell C24, the transition temperatures were plotted against embrittling time for each embrittling temperature, using the data of Table I for the range of treatments covered there, and using previous dataL for other treatments. By cross-plotting, Fig. 1 was finally obtained. It appears from this figure that if there are two noses to the time-temperature-transformation diagram for isothermal temper embrittlement of SAE 3140 steel, the upper nose must lie above 575°C. As interest had been primarily in the region near 525°C, a thorough study of embrittlement above 575°C was not made. Fig. 1 shows that such a study is needed, especially since preliminary measurements at the tempering temperature of 675°C reveal much more rapid embrittlement than would be expected from the data at lower temperatures. The transition temperature increased, approximately, as the logarithm of the tempering time (about 81/2°C, after correction for hardness, for each doubling of the time).& Since three tensile specimens tempered 240 hr at 675°C and water-quenched had an average yield-tensile ratio of 0.844, with 70.7 pct reduction of area, it is assumed1 that the tempering temperature is not above Ae1. The omission of a quench-and-reheat operation below Ae1, for the groups tempered at 675°C, as compared to groups embrittled at lower temperatures, should not affect results." More detailed investigation of embrittlement above 575°C is now under way.

Jan 1, 1954

By A. B. Dyes

In the operation of a water-driven reservoir, a free gas saturation can he established by maintaining production rates fast enough to cause the reservoir pressure to decline below the bubble point. The benefit of such a procedure on the displacement efficiency of the oil by water is illustrated for two types of sandstone samples from one reservoir. Those rock samples showing the poorest recovery to water drive in the absence of a free gas saturation give the most improvement in the presence of a free gas saturation. Employing one type of reservoir fluid (low .shrinkage) the benefit of evolved gas on oil displacement is calculated at several reservoir pressures below the bubble point. This demonstrates the presence of air optimum pressure for water displace,rnent which is .several hundred pounds per .square inch below the bubble point. Displacement of oil at the optimum pressures results in the removal of 7 to I? per cent more oil than would be obtained by displacement at the bubble point pressure. In some instance. this magnitude of increase may be worthy of consideration in establishing the maximum efficient rate of production from many water-driven fields. INTRODUCTION For many years the belief that the use of gas could be combined with that of water to result in a better recovery operation than the use of either of these phases alone has been prevalent in the oil industry. This is evidenced by the several field trials which have employed the injection of gas and water in combination.1.2.v In most cases the small amount of any additional recovery, in combination with the limited knowledge available on the older reservoirs, has left uncertain the influence of gas on recovery in these earlier attempts to combine the use of gas with water. Research studies in recent years have demonstrated through laboratory. floods the magnitude of the benefit of the presence of gas on the displacement of oil from sandstones by water.5,6,7,8,9,10,11 In addition, studies of the distributions of the phases which are present during the displacement process have developed concepts concerning the mechanism by which this benefit is obtained."." It now remains for the engineer to incorporate these findings into the general technology of reservoir operations. It is the purpose of this paper to illustrate that the production of a water-driven reservoir at pressures below the bubble point is one such operation whereby improved oil displacement efficiency can be obtained as a result of gas evolution. The establishment of a gas saturation by evolution of dissolved gas was one of the techniques used by several investigators in the laboratory studies. The possibility of obtaining the benefit of gas evolution in water-driven reservoirs by producing at pressures less than the bubble point has 3150 been pointed out previously." The desirability of this operation has undoubtedly

Jan 1, 1955