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By E. H. Griswold

PETROLEUM production engineering is essentially the application of the laws of 'physics and mechanics to the production of oil. A true production engineer is one who can apply the principles of mechanics, hydraulics, thermodynamics, and other physical arts so as to facilitate' the efficient and economical recovery of oil. To this end he must have wide knowledge of practical field operations and a working knowledge of geology, with a thorough understanding of the underground state of the fluids which he seeks to control. Generally speaking, there have been two schools of production engineering, the California development and the Mid-Continent development. In California the preliminary production engineering was done by representatives of the U. S. Bureau of Mines and the State Mining Bureau, and was essentially subsurface work in regard to casing patterns, water shut-offs, and completion procedure, inasmuch as these were problems of the utmost immediate importance to the California industry. The oil companies were impressed with the value of such work and placed engineers, or more strictly speaking, engineering geologists on their field staffs. These men, realizing the need of further technical study in the mechanical phases of producing operations, added mechanical and mining engineers, and out of the blend developed what is now known as a production engineer.

Jan 1, 1932

By R. H. Card

THE geophysicist is interested greatly in the resistivities of different formations or parts of the earth's crust; sometimes he is interested in a single figure in the nature of an average, or what may be called the effec-tive resistivity, of all the formations or parts of the crust at a given location to a certain depth. Conversely, the engineer whose work is concerned with the coordination of power and communication lines in such a manner as to avoid electrical interference between these two types of facilities is concerned very much with this effective resistivity; to a lesser extent he is interested in the resistivities of individual forma-tions. In connection with his work the coordination engineer often estimates the effective resistivity; of the earth from measurements of the mutual impedances of lines; in other instances he derives front measure-ments of a different type the resistivities of different layers of the crust and synthesizes these to obtain an estimate of the effective resistivity. In connection with problems concerning the coordination of power and telephone lines, measurements from which earth resistivities can be estimated have been made at many points in the United States by the American Telephone and Telegraph Co. and its associated companies, in cooperation with various power companies and with railroad companies having electrified sections of line. Other such measurements have been made by the Joint Subcommittee on Development and Research of the Edison Electric Institute and the Bell Telephone System. A study has been made of the correlation of these data with the ages and types of materials of the strata of the earth's crust at each of the test locations. The purpose of this study has been primarily to provide information to assist in predicting the order of magnitude of effective resistivities to be used in the preliminary consideration of problems con-cerned with the coordination of power and communication lines. To make the results of such a study, most useful, it is essential that the geo-logical data to be secured when such a prediction is wanted be readily obtainable. Data on the ages of strata and the types of materials com-posing geological structures can usually be obtained, therefore particu-

Jan 1, 1937

By L. E. Welch, C. W. MacGregor

THE past 20 years has seen a revolutionary change in the testing of materials at elevated temperatures. This has largely been brought about by the practical importance of the creep problem in the design of equipment for operation at high temperatures. The short-time tension test carried out at elevated temperatures has consequently assumed in many cases only a secondary role compared to the long-time creep test. In many problems, however, a knowledge of the short-time high-temperature properties of a metal is essential. To mention a few cases in which this is so, we might include all hot-forming operations such as rolling, forging, etc., the welding problem if residual stresses are to be computed, the development of new alloys for high-temperature service where, in the initial stages, time does not permit of long duration creep tests, and others. Usually a knowledge of the short-time properties is desired, if only to serve as a background against which to compare creep behavior. In conducting short-time high-temperature tests, it has been almost the universal practice to determine stresses based on the original area and strains referred to the original gauge length which is usually chosen as 2 in. It has been pointed out 1,2 that one of the reasons that the true significance of certain tensile properties is not understood, which perhaps curtails their usefulness, especially when the material is applied [ ] under conditions differing considerably from the test conditions, is that true physical stresses and strains are not determined. It is essential to have a knowledge of true stresses and strains in order to correlate the values present in the tension test with those existing, for example, in hot metal-forming operations and in the general problem of the yielding and failure of metals under combined stress conditions. Stress and strain values based on the original dimensions of the test piece do not exist physically except in the early stages of the test. Since at the present time there is very little information in the literature on true stress-strain relations in the short-time high-temperature tension test it is proposed in the following pages to describe a

Jan 1, 1942

By A. W. Dana, F. J. Shortsleeve, A. R. Troiano

The phenomenon of flake formation which may occur during cooling or room temperature aging of large steel sections is caused by a combination of hydrogen and stress. As such, the transformation characteristics of austenite in a particular section play a major role in the occurrence of flakes. Isothermal and continuous cooling studies demonstrated that flake formation is particularly sensitive to nature, distribution, and relative proportions of the microconstituents in the cooled sections. In the absence of transformation stresses, abnormally high hydrogen contents could be tolerated without flaking. Flakes were not found in the absence of transformation stresses. No correlation existed between average hydrogen content and flake formation where cooling stresses were low. Hydrogen in the steel was a necessary but not sufficient condition for flake formation. DESPITE a long history of laboratory and commercial investigations, the mechanics of hairline crack or flake formation in large steel sections remains unresolved. Recent publications1-5 emphasize that flaking is still a serious problem.* immediate cause of flaking is the pressure exerted by entrapped hydrogen gas that collects in small internal cavities or voids during cooling. When the temperature of steel is low enough to preclude extensive plastic flow, the cavity pressures cause local brittle fractures (flakes). 2—The microsegregation and transformation stress theory. According to this theory, flaking results from stresses arising from the volume changes accompanying the formation of martensite and/or low temperature bainite after a large portion of the section has transformed at elevated temperatures. The proponents of the hydrogen theory suggest that stresses from other sources may be additive to the stresses produced by hydrogen and may, in some cases, provide the increment of stress necessary to produce flakes. However, the critical aspect of this theory is that no stress, except that produced by hydrogen pressure, is believed to be sufficient in itself to produce flakes. For practical considerations, this means that the only safe thermal cycle is one that reduces the hydrogen content to some low critical level below which flaking will not occur regardless of the state of transformation. On the other hand, the transformation stress theory does not rule out the embrittling effect of hydrogen. This theory simply proposes that flaking does not occur in the absence of transformation stresses. With the transformation stress theory as a basis, commercial antiflaking cycles would be designed to complete the transformation of austenite prior to cooling to ambient temperatures. Both of these theories appear compatible with industrial experience. Antiflaking cycles, such as cooling at extremely conservative rates or soaking for periods of time in the upper subcritical temperature ranges prior to cooling to ambient temperatures, may allow the following: 1—diffusion of hydrogen out of the section and, as a consequence, reduction of the hydrogen content to some safe minimum value (hydrogen pressure theory) and 2—complete decomposition of the austenite at rela-

Jan 1, 1956

A. M. Strong, Bishop, Cal. (communication to the Secretary*) :—The paper of Mr. Eeyes gives us the most complete account of the geology of the borax-deposits in the Death Valley region that has yet been published, and is a valuable addition to the geological literature of a very interesting though little-known territory. During the past ten years I have been over much of the northern part of the country described by Mr. Eeyes and the region immediately north of it.

Jan 1, 1910

By T. L. CONDRON

DURING the past fourteen months, this committee has had under consideration and study the subject of the licensing or registration of engineers. The fifteen members of the committee as appointed by Council, were selected from thirteen states, viz., Connecticut, New York, Pennsylvania, Ohio, Georgia, Louisiana, Illinois, Missouri, Iowa, Minnesota, Colorado, California and Washington, and therefore represent practically all sections* of the United States, as well as mechanical, electrical, mining, hydraulic, municipal, sanitary, railway and' structural engineering, and also colleges of engineering. The first work of the committee was to investigate the general subject and to collect, so far as possible, avail¬able material bearing upon the subject in hand,. in-' eluding opinions from many engineers as to the need or desirability of legislation, as well as copies of all state laws passed and proposed, having to do with licensing or registration of engineers, architects and surveyor, This preliminary investigation disclosed that very pronounced views were held by engineers throughout the country, both for and against state licensing or registration. The general sentiment one year ago was more opposed to such measures than it is today. The older members of the profession did not as a rule favor licensing nor did they feel there was need for state regulation of engineering practice, while among the younger men there was a feeling that licensing or registration by the states would add prestige to professional engineers and in many ways benefit -the profession, as well as individual engineers.

Jan 1, 1920

By Willard H. Mutchler, Henry S. Rawdon

As part of the study of the scratch-hardness method for metals and alloys in coöperation with one of the technical committees of the American Society for Testing Materials, the authors devoted considerable time to the study of the changes in hardness of metals when cold rolled. Various seemingly contradictory statements in the technical literature concerning the applicability of the serateh-hardness method to such cold-rolled materials were among the principal reasons for the investigation, the results of which are in part presented here. The effect of cold deformation, at least in moderate degrees, on the hardness as measured by the Brinell aid other common methods is so well known that further reference is unnecessary. Concerning the scratch hardness of cold-worked metals, Martens and Heyn state1 that hardness, whether measured by the ball-pressure rnethod (Kugeldruck) or by the scratch method (Ritzverfahren), will be increased by cold deformation. Jeffries and Archer give a seemingly diametrically opposite view. They state,2 that "the sclerometer (indicating scratch hardness) shows little difference between cold-worked and annealed metals." Portevin3 has shown that the rate at which the Brinell hardness of cold-worked copper and brass increases is very high for relatively slight deformation after which it drops to zero, while Merica and Walten-berg4 have observed that monel metal shows evidence of softening on

Jan 1, 1924

By M. Hansen, C. C. Wang

THE 6 phase of the Cu-Sn system (Cu,Sn, = 32.53 pct Sn) was considered to be stable down to room temperature until Owen and Iball showed by X-ray analysis that it undergoes a eutectoid decomposition into the a and e phases at a temperature near 300°C. Their findings have been confirmed by other investigators,2-5 a11 of whom used X-ray diffraction studies of powdered specimens. The transformation temperature has been reported by various authors as lying between 300" and 380°C. To our knowledge, no microscopic evidence of the eutectoid reaction has been published, as yet. It was in connection with other work that longtime annealings at 300°C were conducted with alloys containing 22, 32.6, and 35 pct Sn (nominal compositions) prepared from electrolytic copper and electrolytic tin and cast in a small iron mold. The specimens were first annealed at 700°C for 5½ hr and quenched from 400 °C after cooling to, and annealing at, that temperature for 9½ hr. After this treatment the sample with 22 pct Sn was cold worked for 15 pct reduction in thickness; the 32.6 and 35 pct alloys were too brittle to be cold worked. The final anneal at 300°C was done in a sealed pyrex glass bulb for 63 days. Fig. 1 represents the microstructure of the annealed alloy with 22 pct Sn. It shows primary a solid-solution crystals in a matrix of the (a + 8)-eutectoid formed at 520 °C. The a crystals contain 8 inclusions which precipitated from a due to the decrease in the solubility of tin in copper with fall in temperature. The decomposition of the 8 phase into a (light) and E (dark) can clearly be seen in some areas of the (a + 8)-eutectoid matrix. However, the decomposition is far from being completed, even after the long-time anneal used. Figs. 2 and 3 depict the microstructures of the annealed 32.6 pct alloy, which was intended to consist entirely of 6 phase. Due to some loss in copper on melting, the composition is slightly higher in tin. Accordingly, the sample after quenching from 400 "C consisted of grains of the 8 phase with the e phase arranged along the grain boundaries and as a Wid-manstatten precipitate within the 6 grains. After an annealing time of 63 days, the eutectoid decomposition, 6 + a + e, is still in its early, initial state. The pearlitic type of eutectoid is nucleated at the c crystals, preferably those surrounding the grain boundaries of the 8 phase. The appearance of the (a + €)-eutectoid in the 35 pct Sn alloy was very similar to that shown in Figs. 2 and 3. Again, the eutectoid decomposition nucleated at the 6 phase and proceeded from there into the 8 areas. As is evident from Fig. 2, the rate of decomposition of the 6 phase at 300°C is extremely sluggish: In the representative area shown only about 5 pct of the 8 phase has undergone the eutectoid trans- formation after annealing for 63 days. On the other hand, Owen and Williams2 have reported that X-ray photographs of powdered samples with between 17.5 and 34 pct Sn did not show 8 lines after annealing for only five to seven days at 300°C. This is in accordance with the fact that the rate of reaction in powdered samples is higher than in massive specimens, as a result of the excess free energy of the powder (due to its greater surface area). The authors wish to express their thanks to C. A. Johnson for preparation of the micrographs.

Jan 1, 1952

By David L. Pentz

This paper discusses exploration objectives and requirements for a nuclear repository in the U.S.A. The importance of designing the exploration program to meet the system performance objectives is emphasized and some examples of the extent of exploration required before the License Application for Construction Authorization is granted are also discussed.

Jan 1, 1984

By B. N. Das, J. B. Darby, P. A. Beck, Y. Shimomura

A POWDER pattern for the (Cr, Mo, Co)R-phase annealed at 1200°C was reported in 1951.1 Through the courtesy of Prof. D. P. Shoemaker of Massachusetts Institute of Technology the authors became aware of inconsistencies of this pattern with that to be expected from the structure of the R-phase, as determined by Komura, Shoemaker, and Shoemaker.' In addition, Professor Shoemaker also noted that the powder pattern reported8 for the (Cr, Mo, Co)D-phase, annealed at 1300°C, agreed fairly well with that to be expected from the newly determined structure of the R-phase. In view of these findings, the previously reported powder patterns of the R-phase and of the D-phase were compared carefully. It was found that most lines in the two patterns (22 lines) were to a fair degree of accuracy related to each other by a single conversion ratio for the corresponding d spacings. The value of the conversion ratio was found to be 1.0058, so that the d spacings of the D-phase pattern previously reported were larger than the correspond- ing d spacings for the reported R-phase pattern by approximately that factor. The intensity relationships between corresponding lines in the two patterns were similar, as judged visually. The following exceptions were noted: Lines Nos. 2, 3, 19, 20, 21, 22, 23, 24 and 25 of the D-phase pattern,:' all designated weak or very weak, were missing in the apparently somewhat fainter R-phase pattern used.' Lines 1, 2, 3, 5, 15 and 34 (all weak or very weak), reported for the R-phase, were missing in the D-phase pattern, Table I. In view of these observations it seemed probable that the two patterns differed from each other mainly by a slight change in lattice parameters (with c/a remaining approximately constant), and that the pattern previously reported for the R-phase contained some weak extraneous lines, most likely to be attributed to impurities. In order to check the foregoing assumption, two specimens of an alloy containing 43 pct by weight of cobalt, 41 pct molybdenum, and 16 pct chromium, which was within the composition range of stability of both the D-phase at 1300°C and of the R-phase at 1200°C, were homogenized at 1200 and 1300°C, respectively, and quenched. X-ray diffraction patterns of the powders of these two specimens, obtained with CrK radiation, were compared carefully with each other and with the D-phase pattern previously

Jan 1, 1959

By M. C. Udy, F. W. Boulger

AN incomplete phase diagram for the U-Ti systern was determined earlier 1 and more recently, a tentative diagram was presented for the uranium-rich end of the system.' In the present re-examination of the whole system of U-Ti alloys, high purity materials were used. Melting stock for the alloys was high purity uranium, containing about 0.09 pct C as the only appreciable impurity, and high purity iodide-process titanium purchased from New Jersey Zinc Co. Both metals were cold rolled to about 1/6 in. thickness, sheared to about I/' in. squares, and cleaned by pickling. The alloys were arc melted under a helium atmosphere in a water-cooled copper crucible. A thoriated-tungsten electrode was used. The furnace chamber was evacuated, then flushed with helium, prior to each melting. It was finally filled with stagnant helium at one atmosphere pressure. Each alloy was remelted three times after the original melting, to insure homogeneity. The alloy button was turned bottom side up before each re-melting operation. Some 22 alloys were examined. Their compositions were spaced at appropriate intervals between 100 pct Ti and 100 pct U. Analyses were made on chips taken after fabrication. The major contaminant was carbon, which varied from 0.03 to 0.08 pct. It appeared in the microstructure as titanium carbide. Alloy compositions were calculated to a carbon-free basis for consideration on the diagram. Tungsten and copper, possible contaminants from the melting operation, were generally less than 100 parts per million each. Fabrication All alloys were forged and rolled to bars approximately V8 in. square. They were clad either in SAE 1020 steel or in a 5 pct Cr-3 pct Al-Ti-base alloy, depending on the fabrication temperature. A temperature of 1800°F (980°C) was used for alloys near the compound composition. This necessitated using the titanium-base alloy, since iron reacts with titanium at this temperature, producing a low melting alloy. Other alloys were fabricated at 1450°F (790°C), using steel jackets. No iron-titanium reaction occurred at this temperature. The jackets were welded in place in an argon atmosphere. Those alloys sheathed in steel were declad and then reclad between rolling and forging operations. On the other hand, those clad with the titanium alloy were cut to a roughly rectangular shape prior to clading and were then carried through both the forging and rolling operations without opening. Those alloys near the compound composition were found to be cracked when the clading was removed. The cracked materials had been plastically deformed, however, and at least some of the cracking had OCcurred during cooling. Heat Treatment The rolled bars, after being declad and shaped to remove surface contamination, were all given an homogenizing treatment of 160 hr at 2000°F. (Samples were taken for analysis following the declading and shaping operations.) All were heat treated at the same time in one furnace, but each was sealed in a purified argon atmosphere in an individual Vycor glass tube. Argon pressure was such that it was approximately atmospheric at temperature. One end of each tube contained titanium chips and this end was heated to 1200°F (650°C) for 10 min prior to the heat treatment. This purged the atmosphere of residual reactive gases. The balance of the tube was warmed during the purge to liberate adsorbed moisture and gases, which also reacted with the hot chips. The bars were furnace cooled from the homogenization treatment. Specimens of each alloy were water quenched after 2 hr heating at 1000°, 1200°, 1400°, 1600°, 1800°, and 2000°F (540°, 650°, 760°, 870°, 980°, and 1095°C). In addition, some were treated at intermediate temperatures of 1300°, 1500°, and 1700°F (705", 815", and 925°C) and at 2150°F (1175°C). Specimens, about '/s in. cubes, were cut from the bars, sealed in individual Vycor tubes, and heat treated as described. All specimens heat treated at the same temperature were processed together. Samples were quenched by breaking the Vycor tube rapidly under water. Metallographic Examination Specimens were mounted in bakelite and ground wet on 180 grit paper held on a 1750 rpm disk. They were then ground wet by hand, using 240, 400, and 600 grit papers. The rough grinding was continued long enough to get well below the surface. Specimens were mounted separately because of the variation in the rate of etching between alloys. The specimens were polished with rouge on a 4 in., 1725 rpm wheel covered with Miracloth. Alloys on the titanium side of the compound composition were etched with a solution of 2 pct hydrofluoric acid in water saturated with oxalic acid. A few crystals of ferric nitrate were added as a bright -ener. Specimens were immersed 5 sec, polished to remove the etch, then re-etched. With the higher titanium alloys, it was often necessary to start the etch on the polishing wheel, because of the formation of a passive film. In some instances, a plain 2 pct hydrofluoric etch was satisfactory. For the alloys on the uranium side of the compound, a distinction between the compound and the uranium phase developed after standing a short time in air. This could be hastened by the application of heat, such as obtained by placing the specimen on a radiator. A deep etch was necessary to develop details in the uranium-rich phase, such as the Widmanstaetten pattern sometimes obtained by quenching y uranium. A 2 pct hydrofluoric acid solution was used for this deep etching.

Jan 1, 1955

By Elmars Ence, Harold Margolin

The partial isothermal section of the Ti-Cu-Zr system at 750°C has been studied. The crystal structure of Ti2Cu has been detevnlined as tetragonal and when expressed as face-centered tetragonal, a = 4.164, c = 3.611 A, and c/a = 0.867. TITANIUM systems in which ß-titanium decomposes eutectoidally into a and a compound rather readily are known as active eutectoid systems. Recently in active eutectoid alloys based on the Ti-Cu system, and containing additionally A1 and Zr, it has been found that zirconium may produce both moderate weakening and marked strengthening at elevated temperatures.' In order to obtain some insight into the behavior of zirconium, an attempt was made to determine whether compounds of the type Tix Zr and TixZryCu exist. Only results of the Ti-Zr-Cu studies are reported here since data on Ti-Zr alloys will be presented in a separate publication. EXPERIMENTAL PROCEDURE Alloys were nonconsumably arc melted in argon atmosphere as 15-g buttons, prepared from materials of the following purities: Bureau of Mines electrolytic Ti — Bhn 63 OFHC copper Reactor Grade Zirconium Bhn .C 150 Prior to alloy preparation the electrolytic titanium was premelted to avoid excessive spattering of the charges during melting. The weight losses during melting did not exceed 1 pct of the total weight of the charge. The alloys prepared are listed in Table I together with analyzed compositions. in cases where they were determined. Specimens for heat treatment were wrapped in molybdenum sheet, annealed in argon-filled quartz capsules, and quenched by breaking the capsules under iced brine. About 1 to 2 sec are required to remove the capsule from the furnace, drop it into the iced brine and break it. Because of the insulation effect of the quartz, this delay in quenching should not produce a large drop in temperature. In many instances, no difference or little difference has been detected between the same specimen quartz capsule quenched or directly quenched in vacuum by an impinging stream of water.' Ti-Cu-Zr alloys were homogenized for 48 hr at 825°C prior to annealing at 750°C. Annealing for 5 days at 750°C produced the same results as 11 days at temperature and it was therefore assumed that 5 days was sufficient to establish equilibrium. Titanium-copper alloys were homogenized for 40 hr at 900°C prior to heat treatment for 6 days at 750°C. Specimens for metallographic examination were electropolished and etched with the following solution:

Jan 1, 1962

By Ernest R. Lilley

THOSE acquainted with the fundamental differences between the, educational .systems of Great Britain and. the United States would hardly expect .the training of men for the petroleum industry to proceed along the same lines- in the two countries. The youth entering an English university is, except in the case of-the student from some of the British colonies or from foreign countries, almost invariably a product of the English "public school." This is in no sense comparable .with the American public school. It is a private institution, frequently of considerable age and heavily endowed. In many aspects it seems to be comparable with the American preparatory school. Yet to some extent its activities remind one of the first years of the undergraduate: course in the American college. Certainly it takes to-itself many of the social, functions of the American college.

Jan 1, 1931

By Daniel Martin

THE dendritic pattern of steels shown by deep etching is used exten-sively as a guide in the inspection and control of the great majority of steels used in ordnance construction and for many applications in indus-try. The value of macroexamination lies in the fact that it enables a qualified observer to make deductions concerning the general quality of relatively large pieces of metal. Not the least of these general character-istics is the nature and extent of the dendritic segregation in the steel, either in the cast condition or after mechanical or thermal treatments. PREVIOUS WORK Detailed methods for macroetching were outlined by Yatsevitch1 and have been standardized by the American Society for Testing Materials. A comprehensive study of the occurrence, origin, and the chemical and crystalline nature of the dendritic structure was given by Keshian.2 The marked persistence of dendritic segregation in nickel steels is fully recognized and has been discussed by Sauveur and Reed' and by others. Very little, however, has been written on the effect of the various elements on the macrostructure of steels, and the purpose of this report is to submit evidence concerning the influence of chromium, nickel and molybdenum on the macrostructure of some 4 by 4-in. ingots cast and treated under closely controlled conditions.

Jan 1, 1939

By C. H. Li, J. A. Sarteli, W. Bonfield

The stress to cause a permanent micros train of 2 x 10-6 in. per in. (defined as the microscopic yield stress) in beryllium is found to be very sensitive to surface condition. The initiation of plastic flow in as-machined specimens, which contain a high density of twins and large residual compressive stresses to a depth of about 0.010 in. from the surface, occurs by the nucleation of slip from high stress fields around twin tips in the surface layers. Removal of the damaged surface layers by chemical polishing results in an appreciable increase in the microscopic yield stress, which is attributed to the remoual of stress raising twins rather than to the release of residual stresses. IT is well known that the macroscopic mechanical properties of beryllium are very sensitive to the nature of its surface finish. Matthews and coworkers1 have shown that the tensile strength of rolled beryllium sheet can be considerably enhanced by chemical polishing, presumably due to the removal of the high concentration of notches, twins, and large residual compressive stresses normally found in the surface layers of ground specimens. There was no similar effect of surface condition on the macroscopic yield stress (the stress to produce 0.2 pct strain), a point to which we shall return. To date there have been no extensive studies on the effect of surface condition on the microstrain region of beryllium. Recent measurements2 of the stress to cause a permanent microstrain of the order of 10-6 in. per in. on machined hot pressed QMV beryllium, gave values ranging from about 600 to 3500 psi. Reasons for the considerable scatter in the results were not discussed. In this paper, we wish to show that the microstrain region of polycrystalline beryllium is, in fact, very sensitive to surface condition. The stress to cause a permanent strain of 2 x 10-6 in. per in. which for convenience we will refer to as the microscopic yield stress, can be increased by almost a factor of 2 by minimizing the amount of surface damage. The variation of microscopic yield stress is attributed to the presence of twins in the surface layers, rather than to residual compressive stresses. EXPERIMENTAL PROCEDURE The starting material for this investigation was commercial purity hot pressed QMV beryllium, with an average grain size of about 20 . The effects of surface treatment were assessed by comparing the microstructure with residual stress measurements and microstrain properties. Specimens were prepared for metallographic examination by wet grinding with 240, 400, and 600 papers, followed by polishing with a suspension of fine alumina in warm 20 pct oxalic acid. After this treatment, the surface could be examined, either in the unetched condition with polarized light, or after etching in 2 pct hydrofluoric or boiling 20 pct oxalic acid with normal light. Thin films of beryllium suitable for electron transmission study were prepared from bulk specimens by chemical polishing with hydrofluoric and phosphoric acids, and were examined in an RCA-EMU-3 electron microscope. Residual stress measurements were made by a procedure similar to that of Treuting and Read.3 Specimens 4.0 in. by 0.625 in. by 0.080 in. were prepared by three different machining procedures, namely: turning, grinding, and sawing, in order to introduce different levels of residual stresses. An electrical resistance strain gage was attached to one side of the specimen, and that surface was then protected by an acid-resistant enamel, while the opposite surface was removed by etching approximately one thousandth of an inch at a time. The change in specimen curvature in the longitudinal direction was noted and the residual stress was calculated by the following formula:

Jan 1, 1963

By Chu-Phay Yap

This is the second of a series of theoretical papers on the iron-carbon system. The solidus lines obtained as a result of the different methods of investigation fall naturally into three classes; namely, (1) a straight solidus originally suggested by Roozeboom and supported in the main by the work of Carpenter and Keeling, Andrew and Binnie, et al., using the thermal analysis method; (2) a slightly curved solidus obtained by Honda and his associates, using the magnetic and resistance methods: and (3) Gutowsky's solidus obtained by the Heycock and Neville method of quenching at successively higher temperatures. Although the thermodynamic method as shown in the first paper indicates that the straight solidus line is correct, no explanation has ever been advanced to account adequately for the disagreement in the shape and position of the solidus lines obtained. The nature of the solute in austenit.e is theoretically discussed. On the basis of such discussions, the data obtained by Honda and Endo on the change of magnetic susceptibility of steels with respect to temperature were critically subjected to a graphical analysis, which is shown to yield a rational explanation of the curved solidus lines.

Jan 1, 1931

By George S. Rice

INTERESTING developments going on in European coal mines look to: (1) increasing mechanization; 12) concentration of mining; (3) improvement in safety appliances; and (4) studies in bettering roof supports and. incident to this, studies of the way a roof behaves under different conditions of support and methods of mining and back-filling. Mechanization and concentration off mining is being adapted to a considerable extent in Great Britain coincident with the development of conveyors. Belt conveyors, shaker-pan conveyors, and a combination of the two are used. For pitching seams not steep enough for the coal to run in station-chutes, the. shaker-pan conveyor is more popular. For large outputs from a contiguous face, the double-unit system of face layout-that is, the conveyors meeting at a common point-is favored.

Jan 1, 1934

By AIME AIME

To accompanying photograph of the plant of the Beattie Gold Miner Limited, at Duparquet,. Quebec, is taken from the east and shows the open-pit from which the ore is taken during the summer. The buildings on the level ground above the open-pit are the change rooms, steel shop hoist building, compressor building, machine shop, warehouse, carpenter shop, and oil house. The main five-compartment shaft is shown west of the hoist building. The crusher plant adjoins the headframe and a long conveyor transports the fine ore

Jan 1, 1936

By J. D. Alexander, W. L. Martin, J. N. Dew

This paper presents data obtained using a flood-pot technique to determine the fuel available and the corresponding theoretical air requirements for in situ combustion of crude oils. Since the technique is relatively quick and easy, it is a practical and convenient tool for evaluating reservoirs as fireflood prospects. It is also a research tool which facilitates systematic study of the variables affecting fuel availability and corresponding air requirements. The understanding of these variables is of prime importance to those concerned with the technical and economic development of in situ combustion as an oil-recovery process. The experimental results show conclusively that the fuel available for in situ combustion is not a constant but, rather, varies with crude-oil characteristics, porous-medium type, oil saturation, air flux and time-temperature relationships. Thus, the fuel availability for specified field applications should be determined using actual reservoir crude and core material and the process conditions expected during in situ combustion in the reservoir. INTRODUCTION In situ combustion is a thermal process for recovering crude oil from reservoirs. The thermal energy released during the combustion of a small amount of the oil in place aids in the displacement of the remaining oil. Numerous articles have been published describing the in situ combustion process giving detailed results of laboratory and field experiments.10 In order to engineer an in situ combustion project, a number of important factors must be considered and determined. These factors include: (1) the amount of fuel consumed per unit of reservoir volume swept by the combustion zone, (2) the composition of the fuel consumed, (3) the amount of air required to consume this fuel, (4) the portion of the reservoir swept by the combustion zone, (5) the appropriate air-injection rates and pressures, (6) the amount of oil that will be recovered, (7) the rate of oil production and (8) the operating costs. Nelson and McNiell1 recently have described a procedure which utilizes laboratory combustion-tube data as a basis for the calculation of some of these design factors. Various authors have attempted to describe the in situ combustion process mathematically, and considerable progress has been made. Analytical solutions to the problem of heat transfer from a moving combustion front have been obtained for linear and radial systems."-' All of the published results involve the assumptions that: (1) fuel concentration is constant throughout the reservoir, or that fuel concentration is inversely proportional to the velocity of the front for a given rate of oxygen consumption; and (2) the fuel reacts instantaneously with injected oxygen, while liberating a constant amount of heat per unit weight of fuel at all temperatures. It seems both desirable and reasonable to test the validity of these assumptions experimentally. This paper presents laboratory data which were obtained by means of a "fire flood-pot" method for determining fuel availability and composition, and the corresponding theoretical air requirements for in situ combustion of crude oils under variable conditions. The mechanics of the method are similar to a conventional tube-run experiment.' The important differences involve the size of the reservoir model used and the method for providing the experimental environment. The new method subjects conventionally-sized core samples or unconsolidated sands to a programmed environmental sequence similar to that experienced by a similar volume of rock during the approach and passage of a combustion front in a long tube or in an oil reservoir undergoing in situ combustion. Restored-state samples can also be used. The small samples and relatively simple techniques involved allow an experiment to be set up, run and calculated in about three 8-hour days. This is a considerable improvement over long-combustion-tube techniques which can require several days to run and several more work days to set up and calculate. All the runs presented were run at 40-psig injection pressure. Pressure was not considered as a variable for these experiments, since we previously had found that it had only a small effect on fuel availability up to 600 psig.APPARATUS AND MATERIALS APPARATUS The fire flood-pot apparatus consists of a consolidated

By H. C. Gatos, A. D. Kurtz

WITH the growing interest in the mechanism of self-diffusion of metals, the study of accurate and convenient methods for determining self-diffu-sion coefficients appears highly desirable. It was with this objective in mind that the present investigation was undertaken. Gatos and Azzam1 employed an autoradiographic technique for measuring self-diffusion coefficients of gold. This method involved sectioning of the specimen through the diffusion zone and recording the radioactivity directly on a photographic film. Because of the very short range of the emitted ß rays in gold, the activity recorded on the film was essentially the true surface activity. With proper choice of the sectioning angle, sufficient resolution could be obtained and the entire concentration-distance curve recorded in one measurement. For the boundary conditions of the experiment, where an infinitesimally thin layer of radioactive material diffuses in positive and negative directions into the end faces of a rod of infinite length, the solution of the diffusion equation is C/Cn = 1/v4pDt exp (-x2/4Dt) where C is the concentration of diffusing element (photographic density in this case), C,, is the constant (depending upon amount of radioactive material), x is the diffusion distance, D is the diffusion coefficient, and t is the time. Thus, by plotting the logarithm of the concentration vs the square of the diffusion distance, a straight line results and the slope contains the diffusion coefficient. In this manner, the self-diffusion coefficient of gold can be obtained as a function of temperature. In the present investigation the results reported by Gatos and Azzam1 have been verified, and the autoradiographic technique has been further developed and applied for the determination of the self-diffusion coefficient of gold at a number of temperatures. Furthermore, the energy of activation for the self-diffusion of gold has been conveniently determined. . Experimental Techniques Preparation of Specimens: The inert gold of high purity was received in the form of a rod from which cylinders were cut and machined to a diameter of 0.500 in. The specimens were annealed to a suitably large grain size and the faces were surface ground prior to the deposition of the radioactive layer. The radioactive isotope Au198 was chosen. It was produced in the Brookhaven pile by means of the reaction Au197 + n ? Au108. It decays by ß emission (0.96 mev) to Hg108 with the subsequent emission of a y ray (0.41 mev). 70Au 108 ? 80Hg 108 + -1e°. The half life of the Au108 is 2.7 days so that a strict time schedule had to be maintained in order to secure sufficient activity until the end of the experiments. For this reason, initial activities as high as 10,000 millicuries per gram were used. The gold arrived in the form of foil and was evaporated onto one face of each gold specimen cylinder to a thickness of about 100A. A sandwich-type specimen was formed by welding two such cylinders together. Evaporation of Gold: The gold was evaporated under vacuum from heated tantalum strips which were bent in such a way as to limit the solid angle through which the gold was allowed to vaporize, thus insuring a more efficient utilization of the gold. The specimens rested on flat brass rings which had an inner diameter of 0.475 in. The entire specimen-holding assembly could be manipulated from outside the vacuum system by means of a magnet which attracted a slug of soft iron attached to the assembly. By evaporating inert gold on glass slides under conditions identical to those employed for the radioactive gold, it was found that the thickness of the films was about 100A. Welding: The welding was performed by hot pressing in a stainless steel cylinder. The inside of the cylinder was threaded and fitted for two plugs. The specimens to be welded were placed in the middle of the cylinder and two pressing disks, one at each end, were inserted to avoid shearing stresses as the plugs were tightened. Mica disks were placed between the pressing disks and the specimens to prevent them from welding. The plugs were then tightened with a hand wrench and the entire unit was placed in an argon stream for about an hour to remove the oxygen. The unit was then inserted in the center of an argon atmosphere furnace maintained at about 700°C and left there for about an hour. Because of the difference in the temperature coefficient of expansion of the two metals, as the temperature rose. the pressure on the specimen-rollple increased and a weld resulted Welding was generally satisfactory under the conditions described.

Jan 1, 1955