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By A. A. Venghiattis

A rational approach to the selection of the appropriate perforator to use in each specific zone of an oil well is presented. The criteria presently in use for this choice bear little resemblance with actual down-hole condilions. These environmental conditions affect the elastic properties of rocks. One of these elastic properties, acoustic velocity, is suggested as the leading parameter to adopt for the choice of a perforator because, being currently measured in the natural location of the formation, it takes into account all of the effects of compaction, saturation, temperature, etc., which are overlooked in the laboratory. Equations and curves in relation with this suggestion are given to allow the prediction of the depth of perforation of bullets and shaped charges when an acoustic log has been run in the zone to be perforated. INTRODUCTION When an oil company has to decide on the perforator to choose for a completion job, I wonder if it is really understood that, to date, there is no rational way of selecting the right perforator on the basis of what it will do down-hole. This situation stems from the fact that the many varieties of existing perforators, bullets or shaped charges, are promoted on the basis of their performance in the laboratory, but very little is said on how this performance will be affected by subsurface conditions such as the combination of high overburden pressure and high temperature, for example. The purpose of this paper is to show the limitations of the existing ways of evaluating the performance of perforators, to show that performances obtained in laboratories cannot be extended to down-hole conditions because the elastic properties of rocks are affected by these conditions and, finally, to suggest and justify the use of the acoustic velocity of rocks, as the parameter to utilize for the anticipation of the performance of a perforator in true down-hole environment. EVALUATING THE PERFORMANCE OF A PERFORATOR It is natural, of course, to judge the performance of a perforator from the size of the hole it makes in a predetermined target. Considering that the ultimate target for an oilwell perforator is the oil-bearing formation preceded in most cases by a layer of cement and by the wall of a steel casing, the difficulties begin with the choice of an adequate experimental target material. For obvious reasons of convenience, the first choice that came to the mind of perforator designers was mild steel. This is a reasonable choice for the comparison of two perforators in first approximation. Mild steel is commercially available in a rather consistent state and quality, and is comparatively inexpensive. The trouble with mild steel is that it represents a yardstick very much contracted; minute variations in depth of penetration or hole diameter and shape may be significant though difficult to measure. The penetration of projectiles in steel being a function of the Brinell hardness of the steel (Gabeaud, O'Neill, Grun-wood, Poboril, et al), it is often difficult to decide whether to attribute a small difference in penetration to a variation on the target hardness or to an actual variation on the efficiency of the projectile. Another target material which has been widely used for testing the efficiency of bullets or shaped charges in an effort to represent a formation—a mineral target as opposed to an all-steel target—is cement cast in steel containers. This type of target, although offering a larger scale for measuring penetrations, proved so unreliable because of its poor repeatability that it had to be abandoned by most designers. The drawbacks of these target materials, and particularly their complete lack of similarity with an oil-bearing formation, became so evident that a more realistic target arrangement was sought until a tacit agreement was reached between customers and designers of oilwell perforators on a testing target of the type shown on Fig. 1. This became almost a necessity about seven years ago because of the introduction of a new parameter in the evaluation of the efficiency of a perforator, the well flow index (WFI). The WFI is the ratio (under predetermined and constant conditions of ambiance, pressure and temperature) of the permeability to a ceitain grade of kerosene of the target core (usually Berea sandstone) after verforation. to its vermeabilitv before perforation. The value of this index ;or the present state if the perforation technique varies from 0 to 2.5, the good perforators presently available rating somewhere around 2.0 and the poor ones around 0.8, There is no doubt that, to date, the WFI type of test is by far the most significant one for comparing perforators. It is obvious that a demonstration of a perforator

By A. Muan

Liquidus data are presented for mixtures in the ternary system FeO-Fe2O3-SiO2 in equilibrium with a gas phase with O2 pressures ranging from 10-10.9 to 1 atm. Data obtained are combined with previously published data to construct lines of equal 02 pressures and lines of equal CO2/H2 mixing ratios along the liquidus surface. Courses of crystallization of selected mixtures under conditions of constant total composition, constant O2 pressures, and constant CO2/H2 mixing ratios are discussed. PHASE equilibrium studies of silicate systems where iron is one component are complicated by the fact that iron readily occurs in three different states of oxidation: Fe3+, Fe2+, and Fe0. Success or failure in work with iron silicate systems is to a large extent dependent on control of the oxidation state of iron and all investigations therefore must be carried out under carefully controlled atmospheric conditions. Silicate systems containing only strongly electropositive metals (like Na+, Ca2+, Mg", etc.) can, for simplicity, be treated as condensed systems, that is, the gas phase can be neglected and the phase relationships discussed in terms of the phase rule written in the well known simplified form P + F = C + 1. In the case of iron silicate systems, however, the composition of the condensed phases varies with the gas composition, and a complete picture of phase relationships can be obtained only by varying the gas composition over a wide range. In order to understand the phase relationships in the more complicated multicomponent silicate systems with iron oxide as one of the constituents, a knowledge of the ternary system FeO-Fe2O3-SiO2 is essential, since it constitutes a bounding portion of all such systems. It was with this in mind that the present study was undertaken. Previous Work A considerable amount of work has been done on various aspects of the chemistry and metallurgy of systems containing silica and iron oxides. The two bounding binary systems FeO-Fe2O3 and FeO-SiO2" The first attempt to obtain information on phase relationships of iron oxide-SiO, mixtures at different 0, pressures was made by Greig.' Darken" determined the melting points of iron oxide on solid silica under various atmospheric conditions. Darken did not determine experimentally the composition of the melts at liquidus temperatures but discussed very ably the principles involved in applying the phase rule to the system. In a recent study Schuhmann, Powell, and Michal8 determined experimentally the liquidus surface of a portion of the ternary system and combined the new information with data in the literature to construct a phase diagram. Their method was briefly as follows: Homogeneous mixtures with various contents of SiO2, FeO, and Fe2O3 were made up by melting together stock mixtures in various proportions. Samples of the homogeneous mixtures, the compositions of which were determined by chemical analysis, were then heated in platinum crucibles in an inert atmosphere until equilibrium among the condensed phases was achieved. The samples were quenched to room temperature and the phases present determined by microscopic examination. Assuming that no change in composition takes place during the equilibration run in inert atmosphere, the liquidus surface can be determined, but no information is obtained regarding the partial pressures of 0, of the gas phase in equilibrium with the condensed phases. The author's method, to be described in the next section, permitted the location of points at the liquidus surface as well as a calculation of the corresponding partial pressures of O2. Experimental Method General Procedure: The standard quenching technique was adapted for a study under controlled variable atmospheric conditions. Premelted mixtures of silica and iron oxides in platinum envelopes were held at constant temperature under chosen atmospheric conditions until equilibrium was reached among solid, liquid, and gas phases. The sample was then quenched to room temperature, the phases present identified, and, for the most significant runs, the composition was determined by chemical analysis. The corresponding partial pressure of 0, was calculated from known equilibrium constants of the gas reactions occuring in the furnace atmosphere. Materials: Starting materials were oxides of commercially highest available purity; cp silicic acid was dehydrated by heating to 1350°C for 6 hr and cp Fe2O3 was dried at 400° C for the same length of time. Samples of 10 g were made up by mixing

Jan 1, 1956

By G. T. Davis, C. C. Mattax, M. O. Denekas

The effects of crude oil cornponents on the wellabil-ities of sandstone and limestone were investigated. Fractions containing cornponents differing in molecular weight and molecular structure were obtained from crude oils by distillation, extration and chromatography. Individual fractions were then tested for their effects on rock wettability. Tests indicate that sundstone wetta-bility may he changed by a complex variety of surfactants varying both in molecular structure and molecular weight. Limestone appears to be particularly sensitive to basic, nitrogenous surfactants. INTRODUCTION Investigations in recent years have shown that petroleum reservoir rock wettability can exert a significant influence on the efficiency with which oil can be produced by water flooding. While most reservoirs are presumably water-wet, they niay range in their degree of water-wettability from near-neutral to strongly water-wet.'" Reservoir wettabilities other than strongly water-wet are likely to be induced by adsorption of surface-active components froni the crude oil on the pore walls of reservoir rock.:' Little is known, however, about the nature of the surface-active materials which are likely to be adsorbed by the reservoir rock. Due to the complexity of crude oils. attempts made in the past90 isolate these surface-active components have met with only limited success. It is probable that many different types of surface-active materials arc indigenous to crude oils and that many of these may be adsorbed to varying degrees by reservoir rock. This was cxolored in the studies discussed in this paper. The over-all objective in these studies is to ascertain whether the wettability of a given reservoir can be determined by examining the surfactant content of the reservoir crude. To this end, crude oils were examined to determine the variability of indigeneous surfactants with regard to chemical type and molecular weight. Crude oils were separated by distillation into fractions differing principally in molecular weight, by chroma-tography into fractions containing compounds differing in polarity, and by solvent extraction into nitrogenous and non-nitrogenous fractions. Individual fractions were then tested for their effects on the wettabilities of sandstone or limestone rock samples. EXPERIMENTAL PROCEDURES Fractionation of the Crude Oils Samples of Miocene, Eocene and Jurassic crudes were distilled at temperatures not exceeding 200°C. The final stages of distillation were completed in a molecular still at pressures down to three microns of mercury. Fifteen to 30 fractions were obtained from each crude oil. These cuts were sufficiently broad that separation can be considered to have been effected principally on the basis of the molecular weights of the constituents of the crude oil. A considerable portion (20 to 40 per cent) of the crudes would not distill under these conditions. The residues were recovered and tested with the other fractions. Fractions differing in polarity were separated from a crude of Pennsylvanian age and an extracted sample of Miocene oil by chromatography, using a solid adsorbent. Since surfactants are, for the most part, polar compounds, chromatography should separate many of the surfactants from the crude oil. Such a separation should provide fractions containing compounds differing in molecular structure. Nitrogeneous compounds were extracted from Miocene crude oil with a solution of sulfuric acid in meth-anol. The residual oil was further processed by chonia-tography. Each of the fractions obtained by thesc procedures was dissolved in a non-polar solvent (xylene) and diluted to its original concentration in the crude oil. No attempt was made to maintain an anaerobic atmosphere above the samples while they were being dissolved. These solutions of the fractions were then tested for their effects on the wettability of sandstone and limcstone as discussed in the next section. Measurement of the Effects of Crude Oil Fractions on Rock Wettability No entirely satisfactory method for measuring rock wettability has yet been developed. All methods used are empirical. The imbibition test was used in these studies. This test is based on the tendency of a rock to imbibe the wetting phase spontaneously. For example, if a strongly water-wet rock is first saturated with oil and then placed in water, the water will quickly invade the rock by capillarity and much of the oil will be displaced. If the rock is slightly water-wet, water irnbibition will proceed more slowly and, in many instances, considerably less oil will be displaced. A water-saturated, oil-wet rock will imbibe oil. The initial rate with which water (or oil) imbibition takes place indicates, qualitatively, the degree of water (or oil) wettability of the rock.

By C. S. Finney, John Mitchell

The decline of the beehive coking industry was inevitable, but it had filled the needs and economy of its day. A beehive plant required neither large capital investment to construct nor an elaborate and expensive organization to run. The ovens were built near mines from which large quantities of easily-won coking coal of excellent quality could be taken, and handling and preparation costs were thus at a minimum. The beehive process undoubtedly produced fine metallurgical coke, and low yields were considered to be the price that had to be paid for a superior product. Few could have foreseen that the time would come when lack of satisfactory coking coal would force most of the beehive plants in the Connellsville district, for example, to stay idle; and if there were those like Belden who cried out against the enormous waste which was leading to exhaustion of the country's best coking coals, there were many more to whom conservation was almost the negation of what has since become popularly known as the spirit of free enterprise. As for the recovery of such by-products as tar, light oil, and ammonia compounds, throughout much of the beehive era there was little economic incentive to move away from a tried and trusted carbonization method simply to produce materials for which no great market existed anyway. With the twentieth century came changes that were to bring an end to the predominance of beehive coking. Large new steel-producing corporations were formed whose operations were integrated to include not only the making and marketing of iron or steel but also the mining of coal and ore from their own properties, the quarrying of their own limestone and dolomite, and the production of coke at or near their blast furnaces. As the steel industry expanded so did the geographic center of production move westward. By 1893 it had moved from east-central to western Pennsylvania, and by 1923 was located to the north and center of Ohio. This western movement led, of course, to the utilization of the poorer quality coking coals of Illinois, Indiana and Ohio. These coals could not be carbonized to produce an acceptable metallurgical coke in the beehive oven, but could be so treated in the by-product oven. By World War I the technological and economic limitations of the beehive oven as a coke producer were being widely recognized. After the war the number of beehive ovens in existence dropped steadily to a low of 10,816 in 1938, in which year the industry produced only some 800,000 tons of coke out of a total US production of 32.5 million tons. The demands of the second World War led to the rehabilitation of many ovens which had not been used for years, and in 1941, for the first time since 1929, beehive ovens produced more than 10 pet of the country's total coke output. Production fell off again after 1945, but the war in Korea made it necessary once more to utilize all available carbonizing capacity so that by 1951 there were 20,458 ovens with an annual coke capacity of 13.9 million tons in existence. Since that time the iron and steel industry has expanded and modernized its by-product coking facilities, and by the end of 1958 only 64 pet of the 8682 beehive ovens still left were capable of being operated. Because beehive ovens are cheap and easy to build and can be closed down and started up with no great damage to brickwork or refractory, it is likely that they will always have a place, albeit a minor one, in the coking industry. The future role of the beehive oven would seem to be precisely that predicted forty years ago by R. S. McBride of the US Geological Survey. Writing with considerable prescience, McBride declared: "A by-product coke-oven plant requires an elaborate organization and a large investment per unit of coke produced per day. Operators of such plants cannot afford to close them down and start them up with every minor change in market conditions. It is not altogether a question whether beehive coke or by-product coke can be produced at a lower price at any particular time. Often by-product coke will be produced and sold at less than cost simply in order to maintain an organization and give some measure of financial return upon the large investment, which would otherwise

Jan 1, 1961

By E. C. Hopkinson, A. H. Youmans, R. A. Bergan, H. I. Oshry

A new log has been developed for quantitative formation evaluation which is based on a measurement of the length of time slow neutrons survive before they are captured in the rocks and fluids. The logging instrument employs a cyclically pulsed neutron generator and a gated scintillation counter which is synchronized with the source. The source emits short, intense bursts of 14 mev neutrons once every 1,000 microsec and is quiescent between bursts. During the period the source is quiescent, the detector is electronically actuated for two independent preselected intervals. A comparison of the counting rates during these two intervals gives a measure of the rate of decay of the slow neutrons and of the associated gamma radiation. The average neutron lifetime in most earth formations is in the range from 50 to 500 microsec. It can be measured during a continuous logging operation at conventional logging speeds. The design of the logging instrument is described and the results of tests are compared with theoretical predictiom. Formulas are developed which give the relationship between log response and formation properties. It is shown that the method is particularly sensitive to formation fluid salinity, and that salt water saturation can be measured accurately in either cased or open hole. The measurement can be made independent of borehole size, fluid type, casing and tool position in the hole by properly selecting the intervals during which the measurements are made. The results of tests with a prototype logging tool are given. INTRODUCTION A new nuclear logging system has been developed which employs the Accelatron,* an accelerator-type neutron source, and accurately measures formation brine saturation in an entirely new way. It has produced a type of formation log with better sensitivity, greater sampling depth and simpler quantitative interpretation than any other nuclear log thus far suggested. The new Neutron Lifetime Log* employs a pulsed electromechanical neutron source and a synchronously gated radiation detector. A prototype instrument has been field tested during recent months to demonstrate the operability of the apparatus and the feasibility of the method. Tests in wells and simulated boreholes have confirmed theoretical predictions and have shown that formation param ters can be measured independent of casing and other borehole parameters. Preliminary results of field tests have indicated that the log may have important and widespread applications. BASIC PRINCIPLE OF NEUTRON LIFETIME LOG The Neutron Lifetime Log is based on the fact that neutrons emitted by a source in a well are rapidly but not instantly captured by the material around the source. Their capture is a matter of statistical probability; the greater the number of capturing nuclei and the greater the "capture cross section", the greater is the probability that a neutron will be captured quickly. The average life of a thermal neutron in vacuum is about 13 minutes, but in common earth materials, the average neutron life ranges between extremes of about 5 rnicrosec for rock salt and perhaps 900 microsec for quartzite. The Neutron Lifetime Log responds to variations in this average neutron life. The theoretical basis for a log of this general type has been well understood by nuclear logging experts in many laboratories both in America and in Russia, and develop mental work along these lines has been in progress for many years. The Russian literature has reported both theoretical and experimental work1,2 but in this country there have been no published reports of progress toward a practical logging instrument. The logging instrument is designed to measure radiation produced by slow neutrons during selected intervals when no neutrons are being emitted by the source. The source is arranged to emit neutrons in bursts or pulses. During the quiescent interval between the pulses, it is possible to observe the exponential "decay" of the neutrons and the neutron-induced radiation as the individual neutrons progressively disappear due to capture by atoms in the formation or the borehole. When a short pulse of 14 mev neutrons is emitted by a source in a borehole, the individual neutrons are slowed to thermal energy within a few microsec. Thus, a cloud of "slow" neutrons is formed around the source within 10 to 50 microsec after the pulse. This cloud is most dense within a few inches of the source, and is progressively less dense out to a radius of about 3 ft, where radiation from the source is practically undetectable.

Jan 1, 1965

By Rex O. McLellan

The formation of a monovacancy in a metal is simulated in an elastic model by the displacement of the surface of a small spherical cavity in a large elastic continuum. The application of linear elasticity to this distortion results in a well- known formula for the energy and an expression for the concomitant entropy change due both to the shear strain in the continuum and also to the dilation of the solid resulting from the boundary conditions at the surface of the solid. Elastic data (the sliear modulus and its temperature coelficient) are used to calculate the entropy and energy of formation for many metals. Despite the simplicity of the assumptions involved, the agreement between the calculated entropies and energies and experimental values is remarkably good. In recent years there has been a large increase in measurements of the absolute concentration of mono-vacancies in metals as a function of temperature. Hence new data for both the energy and the noncon-figurational entropy of formation of monovacancies has become available. Recent measurements&apos; of the anomalous (non-Arrhenius) self-diffusion in many bcc metals has also focused interest on the prediction of the thermodynamic parameters of mono- and multi-vacancies in those metals for which no data are available. Damask and Dienes&apos; have discussed the various theoretical calculations of the energy of formation EL, of a monovacancy. These include simple models involving the breaking of atomic bonds on moving atoms from the interior of a crystal to the surface, models combining elastic calculations with surface-energy terms and detailed quantum mechanical calculations. The simler models give the correct order of magnitude of &, but tend to overestimate it by a factor of about two. The quantum mechanical calculations4"7 have been carried out for the noble and alkali metals with generally reasonably good agreement with the available Ef data. The calculation of entropy of formation Sfv14 lnvolves a fundamental calculation of the perturbation of the phonon spectrum caused by the creation of a vacancy. Huntington, Shirn. and wajda8 have given an approximate evaluation of sJV by considering an Einstein model for the localized vibrations in the immediate neighborhood of the defect and then using elastic theory to calculate the entropy associated with the shear stress field in the distorted crystal (as originally proposed by Zenerg). They also included a term due to the dilation of the crystal. They obtained a value of 1.47k for copper, in good agreement with the experimental value (1.50k). However, Nardelli and Tetta- manzi1° have recently shown that neglecting the coupling between atoms (Einstein Model) may lead to a serious error so the agreement may be somewhat fortuitous. In this work simple linear elastic theory is used to calculate the entropy and energy of formation of mono-vacancies. Despite the simplicity of some of the assumptions involved, the agreement with the available experimental data is remarkable. However. the reasonable degree of success in the application of linear elastic calculations to the excess entropy of a solute atom in a dilute solid solution1&apos; indicates that the application of elastic theory to vacancies. where the interaction of different atomic species is not involved, may not be inappropriate. THE ELASTIC MODEL The metal is assumed to be a spherical elastic continuum. A small spherical cavity of volume V = 4i;v:&apos;/3 is cut from the center. removed. and dissolved rever-sibly in the bulk of the material. TO a good approximation no net atomic bonds are broken and the material does not undergo a volume change although the externally measured volume of the body would increase by V. The radius of the sphere of metal is much larger than r Next a negative pressure is applied to the cavity causing its surface to be displaced inward by an amount simulating the relaxation of the lattice around a monovacancy. In this model the energy and entropy accompanying the distortion are taken as 4, and <. As a first approximation the equation of state for the solid is taken as: r = ro(i + *~D LiJ where K is the bulk modulus. P the hydrostatic pressure. Vo the volume of the material at 0°K and zero pressure. and d+/dT = 30. where 0 is the linear thermal expansion coefficient. The variation of entropy with hydrostatic pressure is given by the Maxwell equation: These equations give the entropy change resulting from increasing the hydrostatic pressure from 0 to P as: and since • we have: This is the entropy arising from the dilation resulting

Jan 1, 1970

By L. T. Lloyd, H. H. Chiswik

The operative deformation elements in a-uranium single crystals under compression at room temperature have been determined as a function of the compression directions. The deformation mechanisms noted may be arranged with respect to their frequency of occurrence and ease of operation in the following order: 1 — (010)-[I001 slip, 2—{130} twinning, 3—{~172} twinning, and 4bunder special conditions of stress application, kinking, cross-slip, {.-176) twinning, and (011) slip. The composition planes of the (172) and (176) systems were found to be irrational. Cross-slip was shown to be associated with the major (010) slip system, coupled with localized interaction of slip on the (001) planes. The mechanism of kinking was found to be similar to that observed in other metals in that it occurred chiefly when the compression direction was, nearly parallel to the principal slip direction [loo] and was associated with a lattice rotation about an axis contained in the slip plane and normal to the slip direction: the [001] in the uranium lattice. The resolved critical shear stress for slip on the (010)-[100] system was found to be 0.34 kg per mm2 In a single test it was shown that under compression in suitable directions twinning on the (130) also occurs at 600°C. DEFORMATION mechanisms of large grained polycrystalline orthorhombic a-uranium have been studied by Cahn.1 A major slip system identified as the (010) with a probable [loo] slip direction and a minor slip system on the (110) planes were reported; the slip direction of the minor system was not determined. The twinning systems that were identified experimentally included the (130) and the irrational (172) composition planes; observations of other traces which were not as frequent and which did not lend themselves to positive experimental identification led Cahn to postulate on the basis of indirect evidence that twinning also occurred on (112) and (121) planes. In addition to the foregoing slip and twinning mechanisms, Cahn also observed kinking and cross-slip in conjunction with the major (010) system; the cooperative cross-slip plane was not identified. The availability of single crystals to the present authors has enabled them to check these results, particularly with reference to the doubtful mechanisms and the preference of operation of any one mechanism in relation to the direction of stress application. The tests were confined to compression only, primarily because of experimental limitations imposed by the size and shape of the available crystals. The tests were performed at room temperature except for one crystal compressed at 600°C. The compression directions were chosen to obtain a representative coverage of one quadrant of the stereo-graphic projection. To test the existence of some of the deformation elements that were reported by Cahn, but were not found in the present study, several additional crystals were compressed in specifically chosen directions considered most ideal for their operation. Experimental Techniques The single crystals were obtained by the grain coarsening technique described by Fisher? They grinding and polishing on rotating laps, with final surface preparation performed in a H3PO4-HNO3 electropolishing bath. A typical crystal readied for compression is shown in Fig. 1; their dimensions were rather small and depended upon the testing direction. Crystals isolated for compression in a direction close to the [010] axis, which lay roughly parallel to the longitudinal axis of the polycrystalline rod, were about 3 to 4 mm long and 5 mm2 in cross-section, while those prepared for compression in other directions were smaller. Most of the crystals were free from twin markings and showed no evidence of Laue asterism. Several crystals, however, contained twin traces prior to compression; these were identified prior to compression so as to clearly distinguish them from those initiated during deformation. The origin of the twin markings prior to deformation may be ascribed to two sources: thermal stresses and specimen handling during isolation and preparation. Two other types of imperfections in the crystals should be mentioned: inclusions, which were probably oxides or carbides. and three of the crystals contained a small number of spherical included grains (<0.01 mm diam), which were remnants of unabsorbed grains from the coarsening treatment. The volume represented by these imperfections was small, and their presence presented no difficulties in the interpretation of the macrodeformation processes during subsequent compression. Two compression fixtures were employed: crystals A, B, C, E, and G were compressed in a hand-operated screw-driven jig whose compression platens were designed to minimize axial rotation;

Jan 1, 1956

By J. R. Cahoon, H. W. Paxton

The mechanical properties of unidirectionally solidified A1(rich)-Mg and A1(rich)-Cu castings containing up to 15 wt pct solute have been determined with re -spect to the volume fraction of interdendritic eutectic. Pioperties were determined in the directions pumllel and Perpendicular to that of solidification; the volume fraction of eutectic was varied between the "as-cast" and equilibrizcm amounts by approperiate heat treatment following solidification. The principles of fiber strengthened composites and dispersion strengthened materials are adapted to explain the mechanical properties of these castings. It is generally accepted that castings often have inferior mechanical properties when con~pared to wrought products. However, there is little quantitative data available concerning the factors which make apparently sound castings weak and/or brittle. The relative ease and inexpensiveness of the casting process have always been attractive and, therefore, an understanding of the factors which contribute to the mechanical properties of castings would seem desirable. Such an understanding may lead to an improvement in the mechanical properties to an extent where castings would become competitive in applications where presently only wrought products are considered to have the requisite properties. Such an understanding could also improve the reliability of present cast products. Much of the recent research on castings has centered about determining the extent of segregation in cast alloys. Macrosegregation, particularly inverse segregation, has been studied in some detail 1-8 and a considerable understanding of microsegregation has been obtained.9&apos;10 The effect of solidification rate on dendrite spacing and on the amount of interdendritic eutectic in binary alloys has been established, particularly for Al(rich)-Cu alloys.""0 However, the extension of these ideas to relate the amount of interdendritic eutectic, concentration gradients, micro-segregation, dendrite spacings, and so forth, to the rnechanical properties has been limited. Dean and spear" have related the mechanical properties of an Al-Si-Mg alloy, A356-T62, to the dendrite spacing and have shown that the mechanical properties improve with decreasing dendrite spacing. Passmore et al.12 have shown that annealing at high temperature improves the mechanical properties of Al(rich)-Cu al- loys and Archer and Kempf 13 have shown that an Al-1 pct Mg-1.75 pct Si alloy behaves in a similar manner. Ahearn and Quigley 14 have shown that high temperature homogenization also enhances the mechanical properties of an SAE 4330 steel. However, in the above investigations, no underlying reasons were suggested for the improvement in mechanical properties. The purpose of the present investigation is to relate the mechanical properties of castings to some of the solichfication variables and to derive some equations by which calculations of the mechanical properties may be attempted. In particular, the effect of the amount of interdendritic eutectic and the effect of stress direction with respect to that of solidification on the mechanical properties will be considered. The Al(rich)-Mg and Al(rich)-Cu binary alloy systems were chosen for study. The A1-Mg system was chosen because its constitutional relationships are such that large volunles of eutectic (up to 24 vol pct) may be obtained in the as-cast condition and then be completely dissolved by subsequent heat treatment at about 440°C. This allows a comprehensive study relating the mechanical properties of castings to the amount of interdendritic eutectic. Also the Al(rich)-Mg eutectic is almost a single phase 15 which should make the experimental results more amenable to theoretical interpretation and calculation. The A1-Cu system was chosen for study because of the large amount of related information available concerning segregation, dendrite spacing, and so forth. Unidirectionally solidified castings were used throughout the investigation so that the effect of solidification direction with respect to the direction of applied stress could be determined. THEORETICAL It is well known that upon solidification of binary alloy castings, the nonequilibrium amount of eutectic which forms is given by 10 where fe o is the weight fraction of eutectic, Cs is the solid solubility of solute at the eutectic temperature, k is the equilibrium partition coefficient, and C, is the average composition of the alloy. In the development of Eq. [I], it is assumed that the effects of inverse segregation and diffusion in the solid are negligible, and that no porosity is present. If the casting is homogenized at a high temperature for a long period of time, some (or all) of the eutectic is dissolved and the amount of eutectic for this "equilibrium" condition may be calculated directly from the constitutional diagram. By appropriate intermediate annealing, the

Jan 1, 1970

By K. J. Gill, P. Chiotti, J. T. Mason

Thermal, metallographic, and vapor pressure data were obtained to establish the pkase boundaries and the standard free energy, enthalpy, and entropy of formation for the compounds in the Y-Zn system. Three coinpounds with stoichiometric formulas of YZn, YZn2, and Y2Zn17 melt congruently at 1105", 1080°, and 890°C, respectively. Four compounds with stoiclziometric formulas of YZn3, YZn4, YZn5, and YZn,, undergo perztectic reactions at 905", 895", 870º, and 685ºC, respectively. Three eutec-tics exisl in this system with the .following eutectic temperatures and zinc contents in wtpct: 875ºC, 23.2 Zn; 1015ºC, 51 Zn; 865ºC, 82 Zn. The YZn, pkase undergoes an allotropic transformation. In the two phase YZn2 -YZn alloys the trans.formation gives a weak thermal arrest at 750°C, whereas in the two phase YZn2-YZn3 alloys no thermal arrest is observed and the transformation occurs over a temperature range below 750°C. At 500°C the free mzergies of formation per lnole vavy from —18,090 for YZn to —53,430 fov YZr11 and corresponding enthalpies vary from -24,050 to -92,080. The free energies and enthalpies per g atom as a function of composition show a maximum for the YZn2 phase; the 500°C values are -9580 and -13,180, vespectively. 1 HE only information found in the literature on Y-Zn alloys was the observation reported by Carlson, Schmidt. and speddingl that Y-20 wt pct Zn forms a low melting alloy. The alloy was produced by the bomb-reduction of YF3 and ZnF2 with calcium in an investigation of methods for producing yttrium metal. The solubility of yttrium in zinc has been determined by P. F. woerner2 and reported by Chiotti, Woerner, and Parry.3 In the temperature range 495" to 685°C the solubility may be represented by the relation In these equations N represents atom fraction of yttrium and T is the temperature in degrees Kelvin. The purpose of the present investigation was to establish the phase diagram for the Y-Zn system and to determine the standard free energy, enthalpy, and entropy of formation for the compounds formed. MATERIALS AND EXPERIMENTAL PROCEDURES The metals used in the preparation of alloys were Bunker Hill slab zinc, 99.99 pct pure, and Ames Laboratory yttrium sponge. Arc-melted yttrium buttons contained the following impurities in parts per million: C-129, N-12, 0-307, Fe-209, Ni-126, Mg-13, Ca < 10, F-105, and Ti < 50. Some of the alloys containing 70 wt pct or more of Zn were prepared from yttrium containing 5000 ppm Ti as a major impurity. Tantalum containers were found to be suitable for all alloys studied and were used throughout the investigation. The pure metals, total weight about 30 g, were sealed in 1 in. diam tantalum crucibles by welding on preformed tantalum covers. A 1/8 in. diam tantalum tube was welded in the base of each crucible for use as a thermocouple well. Welding was done with a heli-arc in a glove box which was initially evacuated and filled with argon. The sealed crucibles were enclosed in stainless steel jackets and heated in an oscillating furnace at temperatures up to 1150°C. Homogeneous liquid alloys were obtained within a half hr at these temperatures except for alloys containing less than 20 pct zinc. The latter alloys were held at 1000º to 1100°C for 2 to 3 days in order to obtain equilibrium. After the initial equilibrations the tantalum crucibles containing the alloys were removed from the steel containers and used directly for differential thermal analyses. Further annealing heat treatments for alloys in which peritectic reactions were involved were carried out in the thermal analyses furnace. After thermal analyses the tantalum crucibles were opened and the alloys sectioned and polished for metallographic examination. In the following discussion alloys referred to as "quenched" were obtained by quenching the sealed stainless steel jacket containing the tantalum crucible and alloy in water. The differential thermal analyses apparatus used was a modified version of the one described in an earlier paper., The graphite crucible was replaced by an inconel crucible, the nickel standard and sampie container were separated by a 1/8 in. MgO plate, no getter was used, and provisions were made to

Jan 1, 1963

By H. D. Hodges, J. K. Godbey

In order to better understand the fracturing process, bottom-hole pressures were measured during a number of typical fracturing operations. A recently developed system was used that allows simultaneous surface recording of both the bottom-hole and wellhead pressures on the same chart. The results from six fracruring treatments are summarized on the basis of the pressure data obtained. Al-though no complete analysis is attempted, the value of accurate pressure measurements is emphasized. Important characteristics of the bottom-hole pressure record do not appear at the wellhead because of the damping effect of the fluid-filled column. In four of the six treatments described, the formations apparently fractured during the initial surge of pressure with only crude oil in the well. The properties of the fluids used during the treatments are given and the fluid friction losses are obtained directly from the pressure records. This technique is also shown to be adequate for determining when various fluids, used during the process, enter the formation. INTRODUCTION Hydraulic fracturing for the purpose of increasing well productivity is now accepted in many areas as a regular completion and workover practice. Numerous articles have appeared in the literature discussing the various techniques and theories of hydraulic fracturing&apos;. In general, three basic types of formation fractures are recognized today. These are the horizontal fracture, the vertical fracture, and fractures along natural planes of weakness in the formation&apos;. Any one or all three of these fracture types may be present in a fracturing operation. However, with only the wellhead pressure record as a guide, it is difficult at best to determine if the formation actually fractured, and is almost impossible to determine the type of fracture induced. These difficulties arise in part because the wellhead pressure record, especially when fracturing through tubing, does not accurately reflect the pressure variations occurring at the formation. Several factors contribute to this effect and preclude the possibility of using the wellhead pressure as a basis for accurately calculating the bottom-hole pressure. These factors are: 1. The compressibilities of the fluids which damp the pressure variations. 2. The changes in the densities of the fluids or apparent densities of the sand-laden fluids. 3. The flowing friction of the various fluids and mixtures, which is dependent on the flow rates and the condition of the tubing, casing, or wellbore. 4. The non-Newtonian characteristics of a sand-oil mixture and its dependence upon the fluid properties, the concentration of sand, and the mesh size used. 5. The unknown and variable temperatures throughout the fluid column. Because of these reasons it was determined that in order to obtain a more accurate knowledge of the nature of fracturing, the bottom-hole pressure must be measured along with the pressure at the surface during a fracturing treatment. Even with accurate pressure data, a reliable estimate of the nature of fracturing is still dependent upon knowledge of the tectonic conditions. However, the hydraulic pressure on the formation is basic to any approach to a complete analysis. In order to accomplish this objective a system was developed to record the wellhead and bottom-hole pressures simultaneously at the surface. By recording both pressures on a dual pen strip-chart recorder, it was possible to greatly expand the time scale so that rapid pressure variations would be faithfully recorded. By such simultaneous recording, time discrepancies inherent in separate records are eliminated, thus overcoming one of the most difficult problems associated with bottom-hole recording systems. This paper illustrates the results obtained by using this system during six typical fracturing operations. All of these tests were taken in wells that were treated through tubing. By a direct comparison of the wellhead and bottom-hole pressures, the importance of obtaining complete pressure information during a fracturing treatment is emphasized. THE INSTRUMENTATION AND PROCEDURES The bottom-hole pressure measuring instrument consisted of a pressure-sensing element, a telemetering section, and a lead-filled weight or sinker bar. The pressure-sensing element used was an isoelastic Amerada pressure-gauge element. By using an isoelastic element, no temperature compensation was necessary in the tests described, since the temperature was believed to be well below the maximum temperature limit of 270°F. The rotary output shaft of this helical Bourdon tube element was coupled to a precision miniature potentiometer. The rotation of the pressure-gauge shaft thus changed the resistance presented by the potentiometer

By R. H. Kolb

A long term project at Shell Development Co.&apos;s Exploration and Production Research Laboratory has been the improvement of the accuracy and the ease of BHP measurements. As a result of these efforts, two complete and separate systems have now been built for the automatic logging of BHP variations. The first of these is a small-diameter instrument suitable for running through production tubing on a single-conductor well cable. During the development of this instrument, as much emphasis was placed on providing a high degree of usable sensitivity and repeatable accuracy as on obtaining the advantages of surface recording. The second system combines the benefits of automatic, unattended recording with the convenience of a permanently installed Maihak BHP transmitter.&apos; THE CABLE INSTRUMENT For many years the standard instrument for BHP determination has been the wireline-operated Amerada recording pressure gauge or one of several other similar devices. This gauge records on a small clock-driven chart carried within the instrument, and although relatively precise readings from the chart are possible, they are difficult to ob-tain. a Both the maximum recording time and the resolution of the time measurements are limited by chart size, and when a slow clock is required for long tests, the precision of the time measurement is often inadequate. Since it is impossible to determine the data being recorded until the gauge has been returned to the surface, wasted time often results when a test is protracted beyond the necessary time or when it is terminated too soon and must be re-run. Clock stoppage or other malfunctions which would be immediately apparent with surface recording remains undetected with down-hole recording; the test is continued for its full term with a consequent loss in production time. As new uses for subsurface pressure data evolved, the shortcomings of the wireline instrument became increasingly apparent, and the concurrent development of a surface-recording pressure gauge and the associated high-pressure well cable service unit&apos; was undertaken. Description of the Instrument Because of its ready availability and advanced degree of development, the Amerada bourdon-tube element was chosen as the basic pressure-sensing device. This element converts a given pressure into a proportional angular displacement of its output shaft, and a suitable telemetering system was designed to measure accurately the extent of this displacement and to transmit the measurement to the surface and record it. The telemetering system furnishes a digital record printed on paper tape by an adding machine-type printer. The present arrangement provides a resolution of one part in 42,000 over the angular equivalent of full-scale deflection, giving a usable sensitivity of better than 0.0025 per cent of full scale. An additional refinement simultaneously records on the tape the time or the depth of the measurement, also in digital form. When the instrument is placed in operation, an adjustable programer can be set to initiate a read-out cycle automatically at selected time intervals. When subsurface pressures are changing rapidly, readings may be recorded as frequently as once every 10 seconds; when pressures are more nearly stabilized, the period between readings may be extended to as much as 30 minutes. Because the instrument is surface-powered as well as surface-recording, the maximum period of continuous logging is (for all prac. tical purposes) unlimited. The subsurface instrument is a tubular tool, 1 1/4-in. in diameter and 6.5 ft in length, operating on 12,000 ft of conventional 3/16-in. IHO logging cable. The transmitting section, mounted above the bourdon-tube element in place of the regular recording mechanism, contains no fragile vacuum tubes or temperature-sensitive transistors. This unit has been laboratory-tested to 1 0,000 psi and 300°F and has performed dependably during a number of field operations. The down-hole transmitting arrangement can be fitted to any standard Amerada pressure element, regardless of range and with no modification of the element itself. Calibration To obtain a repeatability commensurate with the sensitivity and resolution of the instrument, it was necessary to develop a special calibrating technique. The manufacturers of the Amerada recording pressure gauge claim an accuracy of only 0.25 per cent of full scale, which is a realistic figure for normal calibrating and operating procedures. An exhaustive investigation was made of the errors inherent in the bourdon-tube element, itself, independent

By R. W. Fountain, W. D. Forgeng, G. M. Faulring

Precipitation of the phase Co3Ti (Cu3Au type) from a Co-5 pct Ti a11oy has been investigated using single-crystal X-ray diffraction techniques. Oscillation and transmission Laue patterns of specimens aged for short-time periods at 600" C indicate the formation of titanium-rich and titanium-poor zones coherent with the {100} matrix planes. Longer aging times at 600° C establish that the equilibrium phase also forms on the {100} matrix planes as platelets. These observations are corroborated by electron metallography; electron diffraction studies show the phase Co3Ti to be ordered. A probable sequence of the precipitation reaction is discussed. A previous publication by two of the present authors reported on the phase relations and precipitation in Co-Ti alloys containing up to 30 pct Ti.1 The results of this investigation established the existence of a new face-centered cubic inter metallic phase, ranging in composition from about 17.0 to 21.7 pct Ti at temperatures below 1000° C The decomposition of the fcc supersaturated solid solution was studied employing hardness and electrical resistivity measurements. The changes in hardness upon precipitation in alloys containing 3, 6, and 9 pct* Ti were found to be associated with an initial increase in hardness followed by a plateau and then a second, more pronounced hardness increase. Investigation of this behavior by electrical resistivity measurements suggested that two different kinetic processes were involved, which, when interpreted in terms of the kinetic relation,2-4 indicated that initial precipitation was in the form of thin plates. On continued aging, the plates impinged during the growth process. The general features of these findings have been confirmed by Bibring and Manenc,5 while, in addition, they report the phase to be ordered. The present investigation was undertaken to provide more definite information on the structural relationships between the precipitate and the matrix. EXPERIMENTAL PROCEDURE Single crystals of a (20-5 pct Ti alloy were prepared from the melt employing the Bridgman technique. Poly crystalline rod, 1/2 in. in diam, prepared from vacuum-melted material, was machined to 3/8- in. diam to remove any surface contamination that may have resulted from hot-working. The crystals were grown under a purified hydrogen atmosphere in high-purity alumina crucibles heated by induction. Considerable difficulty was encountered in attempting to grow monocrystals because of the high melting point of the alloy and the high solute concentration. However, one crystal about 6 in. long was obtained which was essentially a single crystal except for one or two very small grains around the periphery. The as-grown crystal was solution heat-treated for 24 hr at 1200°Cin a purified argon atmosphere and water-quenched. One-quarter-in. slices were taken from each end of the solution heat-treated crystal for chemical analyses, and the remainder of the crystal was mounted and oriented by the back reflection Laue Method. The chemical analysis of the crystal was as follows: Pct Ti Pct 0 Pct C Pct N Pct H Pet CO 5.29 0.08 0.004 0.002 0.0003 Balance By proper tilting of the crystal, it was possible to obtain slices 1/32 in. thick of [loo] and [110] orientation. The solution heat-treated crystal slices were sealed in silica capsules for the aging treatments, with titanium sponge placed at one end of the capsule to act as a getter. All slices were water-quenched from the aging temperatures, the capsules being broken under the water to ensure a rapid quench. Thinning of the slices for transmission X-ray studies was accomplished by a combination of mechanical and electrolytic techniques, the final thickness being about 0.1 mm. Laue patterns of the solution heat-treated crystal indicated that no strain was introduced by the thinning technique. ELECTRON METALLOGRAPHY After X-ray examination, the structural changes attending the precipitation were followed by examination of direct carbon replicas of polished and etched surfaces of the single-crystal slices and extracted phases. The earliest indication of significant structural change was observed after aging at 600°C The structure of a heavily etched, solution-treated crystal is shown in Fig. l(a). Aside from the etch pit pattern, no regularity of background structure is observed. On the other hand, in the background of the specimen heated for 500 hr at 600°C, the etching pattern shows a directionality indicating the influence of minute precipitate particles, Fig. l(b). On electrolytic dissolution of this specimen in 10 pct HC1 in alcohol, a large volume of very small, flattened cubes

Jan 1, 1962

By W. C. Hagel, H. J. Beattie

Seven austenitic alloys of varions base compositions and minor-alloy additions were solution-treated, aged systematically between 1200oand 1800oF, and examined by X-ray and electron metallography. Intragranular preczpitations of µ, Laves, s, ?&apos;, Ni3Ti, and x phases were observed as a function of composition and aging time and temperatwre. Phase solubility limits were detevtnitzed within 100Fo intervals. These inter metallic compounds fall into two distinct general classes, and whichever class predomznates depends on base composition. It has become increasingly evident that multicom-ponent austenitic alloys are well characterized by their precipitation processes. Since certain groups of elements act as one, the relationships among these processes are reasonably simple; complete identification of such processes is usually attainable by a systematic aging study with a combination of techniques centered on microscopy and diffraction. Several nickel- and cobalt-base alloys illustrating cellular precipitation and its interaction with general precipitation were reported previously.1 The group of alloys covered in the present paper demonstrates precipitation-hardening reactions involving two distinct classes of intermetallic compounds where the predominating class appears to depend on base composition. This dependency ties in with a crystal-chemistry regularity first observed some twenty years ago by Laves and Wallbaum but never amplified to our knowledge. Results of electron-microscope and X-ray diffraction studies on systematically aged hot-rolled alloys known commercially as S-816, S-590, Rene-41, Incoloy-901, M-308, and M-647 are reported here. Some of these alloys have previously undergone minor-phase analyses by other investiators. Alloy S-816 was investigated by Rosenbaum, Lane and Grant,3 and Weeton and Signorelli.4 Rosenbaum found only CbC in hot-rolled bars. Lane and Grant found CbC and a small amount of M6C in the cast structure and stated that both carbides form during aging, most of the precipitation being CbC. Weeton and Signorelli found CbC, M23C6 and a weak indication of a phase after a slow step-down cooling cycle from 2250°F. Rosenbaum also investigated hot-rolled samples of S-590 and identified CbC and M6C. Preliminary information on Rene-41, gained partly from the present work, was reported by Morris.5 Long-time precipitation phenomena in Incoloy-901 at 1350°Fwere investigated by Clark and Iwanski.B heir raw data re- semble those of our present heat with 0.1 pct B, while their interpretation of these data resembles our interpretation of data from another heat with only 0.001 pct B; they made no statement as to boron content. No previous minor-phase studies of alloys M-308 or M-647 have been reported. EXPERIMENTAL METHODS Table I gives alloy compositions in both weight and atomic percent. Specimens were solution-treated from 1700º to 2200ºF, aged at logarithmic-time intervals up to 1000 hours between 1200 and 1800 F, and examined in accordance with procedures previously described in detail. &apos; &apos; Phase extractions were carried out in electrolytic cells containing 800 ml of either 7 pct HC1 in denatured ethanol or 20 pct H3PO4 in water. After electrolysis for 48 hr at 0.1 to 0.2 amp per sq inch, residues were separated by filtration or centrifuging. X-ray powder patterns of residues were recorded on a diffractometer for accuracy and on film for sensitivity. Lattice parameters were calculated by least-squares analyses of indexed sin 8 values, and relative abundances were estimated from intensities of strongest lines of each phase. These phase abundances denote relative amounts with respect to each other rather than to the alloy. Mechanically polished specimens were etched in a freshly mixed solution of 92 pct HC1, 5 pct H2SO4, and 3 pct HNO3. Parlodion replicas for the electron microscope were chromium-shadowed in high vacuum at a glancing angle of 20deg. All electron micrographs are reproduced here with the shadowing source above. The correspondence betweenelectronmicrostructures and phases identified by X-rays was established by a high redundancy of correlation between relative amounts at different stages of aging and examination above and below critical transformation or solubility temperatures. EXPERIMENTAL RESULTS S-816 and S-590—The phases found in S-816 and S-590 after various aging and solutioning treatments are listed in Table 11. These data and the observed

Jan 1, 1962

By A. U. Seybolt

In part 111&apos; of this series it was shown that during the selective oxidation of a 3 1/4 pct Si-Fe alloy in damp hydrogen, only silica, (observed at room temperature) as low cristobalite or low tridy-mite or both, was formed as an oxidation product. In some in- „ stances where the film was fairly thin (probably well under 100A) there was some suggestion of an amorphous form of SiO2. The present investigation of oxidation rate showed that the selective oxidation of silicon-iron can be rather complicated, and apparently impossible to rationalize in an unequivocal manner. In some temperature regions, notably near 800" and 1000°C, the data seem to obey the familiar parabolic rate law. However, at intermediate temperatures complications were noted, some of which are possibly due to the order-disorder reaction in the silicon-iron solid solution. IN an earlier report&apos; it was shown that during the oxidation of 3 1/4 pct Si-Fe alloys in H2O-H2 atmospheres only silica films were formed in the temperature range from 400° to 1000°C in hydrogen nearly saturated with water at room temperatures, or at dew points as low as -45°C. In the work to be reported here, some observations are made on the rate of oxide film formation. As in the earlier investigation, electron diffraction patterns generally showed either low tridymite or low cristobalite or both, except for some very thin films. These sometimes showed diffuse rings, presumably due to a very small crystallite size, or in a few cases, diffuse bands probably caused by an amorphous film. EXPERIMENTAL PROCEDURE Vacuum-melted silicon iron made of high-purity materials was rolled into strips 0.014 in. thick, and cut into samples 1/2 in. wide by 1 in. long. Chemical analysis showed 3.2 pct Si and 0.002 pct C. All samples were surface abraded with 600-grit paper, were solvent cleaned, and then placed in an paper,apparatus containing a "Gulbransen type"2 micro-balance. Here the gain in weight of the samples of about 5 sq cm area could be followed as a function of time during the oxidation caused by the water in atmospheres of various controlled water-hydrogen ratios. The water-hydrogen ratios can most easily be described as varying from a dew point of 0°C (PH2O-p^2 = 6.2 x 10-3 , to K (P j -40°C (PH2O/PH^= 1.3 X 10-* Most of the experiments were conducted at the 0°C dew-point atmosphere because drier atmospheres caused so little gain in weight that the accuracy of measurement was poor. Because of this, only the data obtained at PH2O,/P,,,= 6.2 x X3 will be reported. The temperature range extended from 800" to 1000°C; and most of the oxidation runs lasted for about 24 hr. The reproducibility of any reading was about ± 1 ?, but the sensitivity of the balance was about 0.2 ?. The atmosphere, flowing at 200 cm per-min, was preheated to the furnace temperature before contacting the specimen. While the gas flow caused a measurable lift on the sample, it was ordinarily sufficiently constant so that it was not an appreciable source of error. X-ray and electron diffraction checks of the samples before and after oxidation showed no evidence of preferred orientation, either on the metal samples or on the silica films formed. EXPERIMENTAL RESULTS The data obtained are summarized in Table I, and some are given in detail in Figs. 1 to 4. In the fourth column of Table I, kp refers to the parabolic rate constant in the expression (?/cm2)2 = kpt + c [1] where ? = micrograms gain in weight kp = parabolic rate constant in units r2 /cm4 t = time in minutes c = constant It will be noted that in many cases no value for kp is given; this is because in these instances the data did not obey the parabolic rate law. The silica film thicknesses given in the last columns are values calculated from the weight gain, an average tridy-mite-crystobalite density, and by assuming a perfectly plane surface. Fig. 1 shows the data plotted in the form of Eq. [I], hence a linear plot indicates parabolic behavior. It has been frequently observed in the literature that

Jan 1, 1960

By D. G. Westlake

The electrical resistance of a series of V-H alloys (0 to 3.5 at. pct H) has been measured over the temperature range G° to 360°. Interstitial impurities made contributions to the residual resistivity, but not the ideal resistivity. The contribution of hydrogen in solid solution is expressed by Ap = 1.12 microhm-cm per at. pct H; but the contribution of precipitated hydride was negligible. A portion of the so1vu.s for the V-H phase diagram is presented. The solubility limit is given by In N (at. pct H) = (5.828 i 0.009) - (2933 i 44)/RT. Comparison of critical temperatures joy hydride precipitation and published critical temperatures for hydrogen embrittlement suggests the two are related. ThiS study was initiated as part of an investigation of the mechanism by which small concentrations of hydrogen embrittle the hydride-forming metals at low temperatures. It has already been shown that, in the case of hcp zirconium, a reduction in ductility accompanies the strengthening resulting from precipitation of a finely dispersed hydride phase.&apos;&apos;&apos; Our attempts to detect a similar precipitation of a second phase at low temperatures in V-H alloys by transmission electron microscopy have been thwarted because we have been unable to prepare thin foils that are representative of the bulk material with respect to hydrogen concentrati~n.~&apos;~ The present investigation establishes the solvus of the V-H system at subambient temperatures. Subsequently, we hope to be able to determine whether the embrittlement temperature is related to the critical temperature for precipitation of the hydride in a given V-H alloy. veleckis5 has proposed a partial phase diagram for the V-H system based on extrapolations of the pressure-composition relations he measured at higher temperatures. Kofstad and wallace&apos; conducted a similar study of single-phase alloys but did not attempt to establish the phase diagram. Zanowick and wallace&apos; and ~aeland&apos; have studied a portion of the phase diagram by X-ray diffraction, but they investigated no alloys in the hydrogen concentration range 0 to 3 at. pct, the range of interest to us. EXPERIMENTAL PROCEDURE The vanadium was obtained from the Bureau of Mines, Boulder City, Nev., in the form of electrolytic crystals. The analyses supplied with them listed 230 ppm by weight metallic impurities, 20 ppm C, 100 ppm N, and 290 ppm 0. The crystals were electron-beam-melted into an ingot that was rolled to 0.64 mm. Strips, 60 mm long and 4.2 mm wide, were cut from the sheet, and both rolled surfaces were ground on wet 600-grit Sic paper to produce specimens 0.4 mm thick. They were wrapped in molybdenum foil, vacuum-encapsulated in quartz, and annealed 4 hr at 1273°K. The specimens were annealed in a dynamic vacuum of 2X lo-&apos; Torr for 30 min at 1073°K for dehydrogenation, and charged with the desired quantity of hydrogen by allowing reaction with hydrogen gas at 1073°K for 2 hr and cooling at 100°K per hr. Purified hydrogen was obtained by thermal decomposition of UH3. Sixteen specimens were studied: two contained no hydrogen and the others had hydrogen concentrations between 0.5 and 3.5 at. pct (hydrogen analyses were done by vacuum extraction at 1073°K). Electrical resistances were measured by the four-terminal-resistor method on an apparatus similar to the one described by Horak.~ The specimen holder was designed so that both current and potential leads made spring-loaded mechanical contact with the specimen. The potential leads were 30 mm apart, and the current leads were 55 mm apart. The current was 0.10000 amp. We used the following baths for the indicated temperature ranges: liquid nitrogen, 77°K; Freon 12, 120" to 230°K; Freon 11, 230" to 290°K; and ethanol, 290" to 340°K. Temperatures lower than 77°K were achieved by allowing the specimen to warm up after removal from liquid helium. Temperatures above 77°K were measured by a calibrated copper-constantan thermocouple (soldered to the specimen holder) and below 77°K by a calibrated carbon resistor. The temperature of the bath changed less than 0.l0K between duplicate measurements of the resistance. RESULTS AND DISCUSSION Typical plots of resistivity p vs temperature T are shown in Fig. 1. In the interest of clarity, only five curves are presented and the data points have been

Jan 1, 1968

By Larry Kaufman, Martin Schatz

The effect of pressure on the extension of the ? loop in the FeSi system has been determined by means of metallogvaphic studies and hardness measurements performed on a series of high-purity Fe-Si alloys containing 7.5, 11.0, and 13.9 at. pct Si, respectively. These mensurements, performed at 42 kbar and temperatures up to 1200oC, indicate that the ? loop is expanded to about 10 at. pct Si at 42 kbar as opposed to a maximum extension of 4 at. pct Si at 1 atm. Comparison of the experimental results with thermodynamic predictions of the pressure shifts yields satisfnctory results. DURING the past few years, several studies have been performed in our laboratory1-&apos; in order to determine the effect of high pressure on phase equilibrium in pure iron and iron-base alloys. The purpose of these studies has been to elucidate the effects of high pressure experimentally and to compare the observed results with predicted pressure effects derived on the basis of known thermody-namic and volumetric data at 1 atm. These studies have included work on pure iron2,5,7 as well as Fe-Ni,1,5 Fe-cr,l,5 and Fe-c4-6 alloys. In addition, Tanner and Kulin3 have reported results of pressure studies on two Fe-Si alloys containing 2.0 and 6.25 at. pct Si. At the time of this latter study, no detailed information was available concerning the difference in volume between the a (bcc) and ? (fcc) phases in the Fe-Si system as a function of silicon content. In order to compare their observations with calculated pressure shifts, Tanner and Kulin were forced to assume that silicon had no effect on the difference in volume between a and ? iron. The resulting discrepancy between their calculation of the a/? phase boundary at 42 kbar and the observed results led them to the conclusion that silicon additions probably decrease the difference in volume between a and ? iron. Recently: Cockett and Davis8,9 have reported de- tailed studies of the lattice parameters of a series of Fe-Si alloys at temperatures ranging from 20" to 1150°C. These measurements, performed on alloys in the bcc and fcc range, show that silicon does indeed decrease the difference in volume between a and ? iron. By correcting the calculations of Tanner and Kulin in line with the observed effect of silicon they were able to show improved agreement between computed and observed pressure shifts.&apos; The present measurements were undertaken to provide additional corroboration of this effect, by extending the range of composition, in addition to exploring a situation where large extensions of a ? loop could result in impingement of the ? field with an ordered bcc phase (based on Feo.75Sio.25). I) EXPERIMENTAL PROCEDURES AND RESULTS The alloys investigated were obtained from Dr. F. Kayser of M.I.T. They were prepared at the Ford Scientific Laboratory by vacuum melting electrolytic iron and high-purity silicon. The melts were poured under an argon atmosphere into hot-topped steel molds. Subsequently the ingots were hot-worked down to 1/2-in.-diam rods. Three alloys containing 7.5, 11.0, and 13.9 pct Si were studied. Carbon, regarded as the principal impurity, analyzed at, or below, 0.001 wt pct for all of the alloys. Prior to pressure-temperature treatment, the rod was annealed for 24 hr in vacuum at 1000°C, water-quenched, and subsequently machined into 0.100-in.-diam by 0.100-in.-long specimens. Subsequent to machining, the specimens were again annealed and then examined metallographically. They were found to exhibit a clear coarse-grained ferrite similar to Figs. 10 and 110 of Ref. 1 and Fig. 2 of Ref. 3. Subsequently, specimens of each alloy were equilibrated at 42 kbar at various temperatures in supported piston apparatus.1,3,4,6 Three specimens, one of each alloy, were wrapped in platinum and exposed simultaneously. The pressure-temperature cycle consisted of increasing the pressure from ambient to 42 kbar at 25oC, heating rapidly to the desired temperature, holding for 15 min, and quenching to 100°C, followed by slower cooling to 25°C and pressure release. The temperature was measured with a Pt/Pt-13 pct Rh thermocouple which was not corrected for pressure effects. Subsequently, specimens were examined metallographically and by

Jan 1, 1964

By Charles H. Grant

Some of the limiting factors relative to explosive crushing of rock and ways to overcome a few of these problems are presented. Relationships between borehole diameters, bench heights, and spacings, along with a review of the influence geometry has on energy as these are changed, are discussed. Efficiency in use of explosives and the decay of energy as it moves through rock and is absorbed and dissipated, is described, along with fragmentation as a function of spacings and energy zoning, etc. Communications are one of the major problems encountered. In an effort to provide a better understanding of the use of explosives, it is necessary to take a little different view of what explosives are, how to look at them as tools to fragment rock, and some of the problems encountered in doing so. First, take the explosive: although there are many factors involved, consider these as being reduced to only two — shock-strain imparted to the rock by the high early development of energy, and the gas effect which is a combination of heat, moles of gas formed, rate of formation of these gases which develop pressures, etc. First, consider shock energy by itself and assume there is no gas effect in the reaction. Fig. 1 illustrates a block or cube of rock, in the center of which is detonated an explosive charge which is 100% shock energy. Tensile slabbing would be seen on the surface and probably the cube of rock would generally hang together even though microcracks were formed. If the situation is reversed and an explosive whch has no shock energy and only gas effect (Fig. 2) is considered, the cube of rock would act as a pressure vessel and contain the pressure from the gas effect until it exceeded the rock-vessel strength; then the rock would break in a few large pieces. If these two kinds of energy are put together and the area of shock-strain around the explosive (Fig. 3) is considered, the two energies will be seen working together to furnish broken rock. The gas effect applies pressure to the microcracks formed from the shock energy to weaken the rock-pressure vessel and propagate these cracks to break the rock apart. It not only will be broken more finely, but will break apart at a lower pressure than the gaseffect case, since the shock energy has first weakened the rock vessel. Although tensile spalling from the shock-strain imparts momentum to the rock, the main source of displacement comes from the gas effect. The term "rock" is being used to mean any material to be blasted. These energies are absorbed by the rock in different ways. First, classify rock into two main categories: "elastic" and "plastic-acting." Elastic rock should be thought of as rock which can transmit a shock wave and is high in compressive strength, such as granite or quartzite. Since this elastic rock transmits a shock wave well, it makes good use of the shock energy from the explosive-forming cracks, etc., for the gas effect to work on. Plastic-acting rocks are rock masses which are relatively low in compressive strength and absorb shock energy at a much faster rate, thereby making poor use of the shock energy by not developing as extensive a cracked zone for the gas effect to work on. Rocks of this type are generally softer materials such as some limestones, sandstones, and porphyries. For the most part, the shockenergy part of the explosive reaction is wasted in plastic-acting rock, leaving most of the work to the gas effect. Since the ratio of gas effect to shock energy is different in different explosives, it is easy to understand why some explosives perform well in elastic rock and poorly in plastic-acting rock, and vice versa. Some of the most difficult blasting situations arise when mixtures of plastic-acting and elastic rock are encountered (Fig. 4). Fig. 4 shows an example of granite boulders cemented together with something like a decomposed quartz monzonite which is plastic-acting. The elastic granite boulders will transmit the shock-strain within itself, but when this shock tries to move through the monzonite to the next boulder, its intensity is absorbed by the monzonite and little shock-strain is placed on the adjoining boulder. In addition to this loss by absorbtion, shock reflection at the surface of the boulder will effect tensile spalling. The net effect is poor breakage of the boulders which do not have drillholes in them as they simply will be popped out with the muck. The same is true (Fig. 5) when layers and joints make

Jan 1, 1970

By J. Philibert, A. G. Guy

Diffusion in the presence of deformation was studied by the method of vacuum dezincification of copper-rich and silver-rich solid solutions containing 7 to 30 pct Zn. The specimens were designed to permit the study of diffusion in separate portions of a given specimen characterized by strain rates ranging from essentially zero to approximately 10 sec-. No effect of deformation on diffusion was observed. BEGINNING with the work of Buffington and Cohen: interest in the question of the effect of stress or strain on diffusion has largely been concentrated on the enhancement of diffusion in specimens subjected to Continuous plastic deformation. The present research is a contribution to this limited area. However, as a preliminary to focusing attention on this special topic, it will be desirable to make a broad survey of the larger question, especially since there has been considerable foreign work in areas outside those of current interest in the United States. Since most of the topics referred to in the following section are both complex and imperfectly understood at present, it has been expedient in most instances to offer only a guide to the general nature of the work rather than a critical evaluation. PREVIOUS WORK The effect of elastic stress on diffusion has received considerable attention, especially with regard to the thermodynamic driving force for diffusion. The thermodynamic treatments have been based on the work of Gibb, Voigt, Planck, and Leontovich.&apos; Konobeevskii and Selisski6 made a first attempt at treating the problem in 1933, and Gorskii7 a few years later gave a solution applicable to single crystals as well as to polycrystalline specimens. In 1943 Konobeevski8 published treatments that have been the basis of much Russian work up to the present. For example, Aleksandrov and Lyubov used his work in explaining the velocity of lateral growth of pearlite. Early work in the United States was that of Mooradian and Norton, which showed that lattice distortion tends to be relieved before it can significantly affect the diffusion process. Druyvesteyn and Berghoutl1 observed a slight effect of elastic strain on self-diffusion in copper, while de Kazinczy12 found that both elastic and plastic deformation increased the rate of diffusion of hydrogen in steel. On the other hand, Grimes58 observed no effect of either elastic or plastic straining on the diffusion of hydrogen in nickel. High-frequency alternating stresses have been reported by various investigator s13-l5 to increase the rate of diffusion. A special form of elastic stressing is the imposition of hydrostatic pressure, a condition that is amenable to Conventional thermodvnamic analysis. Most of the experimental results in this area are consistent in showing a slight decrease in diffusion rates at high pressures.16-l8 Although Geguzinl reported a pronounced effect of relatively small pressures, Barnes and Mazey20 failed to Corroborate this finding, while Guy and Spinelli21 advanced an explanation of the phenomenon observed by Geguzin. It has been recognized that the thermodynamic treatment of diffusion phenomena in an arbitrarily stressed body is complicated by the fact that the desired state of quasi-equilibrium of the shear stresses cannot be maintained during a general diffusion process. However, attempts have been made by Meix-ner22-24 and Fasto to treat certain restricted cases, such as relaxation. FastovZ7 has also incorporated the general stress tensor into the thermodynamics of irreversible processes. The lattice strain that accompanies the formation of a solid solution has been the subject of much study,28-s0 and indirectly it has entered into many recent theories of diffusion. However, some Russian investigators31&apos;32 have taken other views of this matter and have predicted large effects on diffusion rates because of concentration stresses.o In completing this brief resume of previous work involving elastic strains and before proceeding to a consideration of the effect of continuous plastic deformation, it should be pointed out that deformation of various additional types may also influence diffusion. The effect of cold-working on subsequent diffusion has been studied directly by AndreevaS and by Schumann and Erdmann-Jesnitzer, while indirect evidence has been obtained by Miller and Guarnieri and by Vitman.38 Thermal stresses may also influence diffusion, contributions to this subject having been made by Fastovs7 and by Aleksandrov and Lyubv. The work of Johnson and Martin,o Dienes and Damask,3Band DamaskS considered the question of radiation-enhanced diffusion. In considering previous work on the subject of plastic deformation and diffusion, attention will be directed to those studies concerned primarily with diffusion rather than with its relation to Creep, e.g., the work of Dorn, or to the acceleration of diffusion -controlled reactions. Observations of the effect of

Jan 1, 1962

By D. M. Waldorf

Steam permeability measurements have been made in the laboratory on several samples of natural reservoir materials. The steam temperatures and pressures were selected to simulate conditions which might exist in a reservoir during the injection of steam. For each sample tested, the experimental permeability to superheated steam was comparable to that measured with air and no evidence of plugging was detected. Some samples were exposed to water at various temperatures and plugging was found to occur in materials which contained significant quantities of monmorillonite clay. Temperature had little effect on the degree of plug-ning between 75 and 325 F. The measured pemeabilities tended to increase slightly with temperature, but the changes were small compared with the initial loss of per~neability on wetting. Sequential pemzeability measurements were made on two samples using air, water, steam, water and air, in that order. Both samples were water-sensitive and plugged extensively after the initial injection of water. Upon exposure to superheated steatm the samples dehydrated and their permenbilities to superheated steam were comparable to those initially measured with air. The remaining measuretnetzts with water and air confirmed that the water plugging was reversible and that the samples were not seriorrsly damaged during the tests. INTRODUCTION The swelling of water-sensitive clays during water floods has long been recognized as a potential source of reservoir damage. The recent extensive application of steam injection and stimulation has compounded this problem since both hot water and steam (as well as fresh water at reservoir temperatures) are, at sume time, in contact with the producing zone adjacent to the bore of a steam injection well. The purpose of this paper is to present data which compare the sensitivity of some natural sedimentary rock samples to water at various temperatures, and to super-heated steam. Some properties of montmorillonite clay are briefly reviewed, and comparisons are drawn between empirical data and the predicted behavior of the montmorillonite known to be present in the samples. PROPERTIES OF MONTMORILLONIT E CLAY Water initially adsorbs on dry N a -montmorillonite clay in discrete layers in the interlaminar space between clal platelets. The platelet spacing, which is 9.6 A (angstroms) for a dehydrated clay, has been observed to expand in discrete steps to 12.4, 15.5, 18.4 and 21.4 A spacings, indicating the formation of four discrete layers of regularly oriented water molecules.&apos; The first two layers are easily formed by hydrating a dry sample to equilibrium in an atmosphere with carefully controlled humidity. The formation of the higher layers is more difficult. The usual X-ray diffraction patterns of the more highly hydrated samples indicate a gradual increase in the average spacing betwcen 15.5 and 19.2 A, followed by a discontinuous expansion to 31 A when the weight ratio of water to dry clay is between 0.5 and 1.2.&apos; Platelet expansion above 31 A proceeds monotonically as the moisture is increased and no regular arrangement of the platelets ib observed. Water-sensitivity in sedimentary rocks is usually associated with Na-montmorillonite clay when it is in the noncrystal-line state. Mering3 found that the average lattice spacing of sodium montmorillonite hydrated at 68 F and 70 per cent relative humidity was 15.5 A, and that the spacing, at 92 per cent humidity was 16.5 A. The water adsorbed at the higher humidity has the same free energy as liquid water at 65.6 F. Kolaian and Low&apos; used a tensiometer to measure the thermodynamic properties of water in diffuse suspensions of montmorillonite clays relative to pure water. They observed that water in suspensions as dilute as 6 per cent clay became partially oriented when left undisturbed. The bonding associated with this orientation was not extensive because the free energy difference between the water in suspension and pure water was only a few millicalories per mole. They also found that the measured free energy difference decreased rapidly with temperature and became negligible above 100 F. This evidence indicates that montmorillonites contained in sedimentary rocks would dehydrate to a crystalline structure when exposed to superheated steam, and that the rock permeability measured with steam would be equivalent to that measured with air. The effect of elevated temperatures on the swelline of montmorillonite clays in aqueous suspensions has not been investigated. The Gouy-Chapman diffuse-ion-layer theory has been used to predict the swelling pressure of clay suspensions in dilute salt solutions at room temperature with reasonable success. theory also correctly predicts the direction of the thermal response of Na-mont-morillonite swelling pressures in dilute salt suspensions, 9 Over the temperature range of 33 to 68 F, an increase in

Jan 1, 1966

By A. H. Daane, J. F. Haefling

The limits of miscibility in some of the uranium rare-earth alloy systems have been determined in the temperature range 1000°to 1250°C. The solubilities of lanthanum and cerium in uranium are greater than those of the remaining rare earths by a factor of more than two. The solubility of uranium is greater in cerium, braseodymium, and neodymium than in the other rare-earth metals studied. The values found in this study are in qualitative agreement with those which might be expected if the solubility rules of Hildebrand and Scott are applicable. AS interest in nuclear reactors intensifies, many new types of fuels are being suggested in attempts to improve the economics of some of the proposed reactor schemes. To remove some of the difficulties inherent in the use of solid-fuel elements and their reprocessing, many types of liquid-metal reactors have been suggested. One of the more attractive features of several of these reactor concepts is that they include a continuous or semicontinuous process for the extraction of fission products and "bred" fissionable materials from the fuel, utilizing immiscible metal extractants. This would enable a much higher burn-up of fissionable material to be achieved and would present a very attractive economic picture. Several studies have been reported on equilibrium systems in which there exists a high degree of immiscibility between uranium and another metal that might be used as an extractant in such a processing scheme.&apos; Two of these systems in which a high degree of immiscibility exists are those of uranium with the two rare-earth metals, lanthanum, and cerium. Since the rare earths constitute a significant fraction of the fission products, their removal is of prime importance. It is reasonable to believe that this might be accomplished by equilibrating a rare-earth phase with the contaminated uranium fuel in the liquid state. In order to make a more complete study of those systems which would be of interest either as extractants in a liquid-liquid extraction process, or as fission products formed in the fuel, the alloy systems of uranium with lanthanum, cerium, praseodymium, neodymium, and samarium were studied in some detail in the temperature range 1000" to 1250°C; less detailed studies were made with the other rare earths. In addition to being of value to the reactor program, the data obtained in this study should be of help in making a study of the role played by the electronic structures of metals in determining the nature of metallic solutions. The unique electronic structures of the rare-earth elements make them particularly interesting in this respect. EXPERIMENTAL The usual procedure for a solubility determination was to seal equal volumes of uranium and the particular rare earth in a tantalum crucible under an atmosphere of helium; this crucible was then sealed in a stainless steel jacket in an atmosphere of helium. These samples were equilibrated by repeated inverting of the crucibles in a furnace for 15 min at the desired temperature, left in an upright position for 15 min to permit separation of the two phases, and then quenched under a stream of water. In some runs the temperature of the furnace was held 50&apos; to 100°C above the desired quenching temperature while inverting in order to insure good mixing. However, it was found that above 1200°C the crucibles were subject to failure and for these runs the furnace temperature was not raised above the desired quenching temperature. A small amount of tantalum was dissolved in the uranium and the rare earths in these runs, a maximum of 3 wt pct in the uranium phase at 1250°C and up to 1 wt pct in the rare-earth phase at this temperature. On cooling, the major portion of this tantalum precipitated as primary tantalum crystals. Any residual tantalum would probably have a negligible effect on the mutual solubility of uranium and the rare earths in each other. Samples for analysis were cut from each phase with an abrasive cutting wheel; the region near the interface between the two metals was carefully avoided. In the case of the rare earths with melting points above 1250°C no solubility data were taken on the rare-earth phase since this phase could not have achieved equilibrium in a reasonable length of time. (For the same reason no data were taken on the uranium phase below its melting point of 1132°C.) Equilibrium appeared to have been reached in the uranium phase in these cases although the rare-earth phase had not melted. To verify this, samples were melted together in an arc furnace similar to that described by Kroll.2 These samples were sub-

Jan 1, 1960