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By K. G. Davis, P. Fryzuk

WEINBERG,' in an investigation of solute distributions along unidirectionally solidified rods of dilute silver in tin alloys, concluded that, for rods of 2 mm diameter or larger, a very high degree of convective mixing takes place in the melt ahead of the advancing interface. Suggested causes for this mixing were (1) thermal convection, (2) convective currents set up by changes in fluid density associated with concentration gradients, and (3) currents set up by changes in volume during freezing. Evidence was presented that causes (2) and (3) were not primarily responsible for solute mixing. Further, convective mixing took place in horizontal rods with temperature gradients as low as 6°C per cm, putting doubt on the importance of thermal convection. In the present work, convective mixing is examined in melts with known initial solute distributions. The absence of a moving solid-liquid interface eliminates convection caused by volume changes on solidification. In the first experiment, three silica tubes 2 mm in diameter and approximately 20 cm long were filled with an alloy of 100 ppm Ag in tin (99.999 pct purity), placed adjacent to one another in a horizontal furnace, and solidified from one end at 1 cm per hr. A110, a Y emitter, was used as tracer. A few centimeters at the start of the rod were left un-melted, to determine the initial concentration Go. After solidification was complete, one sample was removed. The hot furnace was then brought back to

Jan 1, 1965

By Klaus Detert, Joseph Ziebs

The progress of primary recrystallization has been studied in alloys of nickel with 0.01, 0.1, and 1.0 at. pct Ta. It can be shown that in one particular alloy containing 0.1 Ta and 0.02 02 a large grain structure (0.5 mm diameter) could he developed by primary recrystallization in thin sheetswhich were cold-worked to 97.5 pct reduction in thickness. Electron transmission microscopy demmstrates that this phenonmena was linked to the occurrence of fine precipitates at subboundaries. A discussion of how fine precipitates can exert under certain circumstances a rather distinct retarding effect on nucleation rate is included. IN a former investigation, the recrystallization and texture formation in a series of binary nickel alloys including pure nickel were studied.' Nickel alloys with some tantalum added displayed peculiarities during primary recrystallization which have been studied in more detail by the present investigation. The alloys were prepared by powder-metallurgy techniques similar to the previous investigations.' Pure metal powders were thoroughly mixed to obtain alloys with 0.01, 0.1, and 1.0 at. pct Ta, referred to as alloy A, B, or C, respectively. The metallic impurities were kept below 0.005 at. pct except iron and cobalt with slightly higher values, by using highly purified nickel and tantalum powders.' The mixtures were pressed into bars and sintered at a temperature of 1300°C. The sintering atmosphere was either high vacuum, dry hydrogen, or wet hydrogen. The sintered rectangular bars were rolled 60 pct to a thickness of 3 mm, then annealed under dry hydrogen at 1000°C and again cold-rolled to about 97.5 pct reduction of thickness. After this heavy deformation, the sheet thickness was 0.07 5 mm. In order to study primary recrystallization which occurred in measurable times between 400" and 600°C, samples of the rolled sheets were annealed in a salt bath at various temperatures for various times in order to study the progress of recrystallization by hardness measurements and microscopic examination. For optical observation, the samples were polished and etched in a solution of 65 cu cm HNO3, and 18 cu cm acetic acid in 20 to 100 cu cm H2O, depending of the particular sample and annealing treatment. For transmission electron microscopic observations, thin foils were prepared by the combined window and Bollmann techniques.273 The electrolyte consisted of 23 pct perchloric acid and 77 pct acetic acid held at 15°C. Polishing was performed at a current density of 0.4 amp per sq cm. The examination of the thin films was accomplished by a Siemens Elmiscope-type electron microscope operating with an accelerating voltage of 100 kv. RESULTS The progress of primary recrystallization was very accurately followed in these alloys in order to establish the time when recrystallization started and ended during an isothermal anneal. Table I gives some of the data obtained by these studies. The time for beginning is a mean value, where somewhat less than 5 pct of the matrix has recrys-tallized. The end of recrystallization is the time when at least 95 pct of the matrix has recrystal-lized. Listed are three different B alloys, with 0.1 Ta distinguished by different sintering atmospheres which influence the rate of primary recrystallization. By the above method, the state of the grain structure when primary recrystallization was completed, but grain coarsening had not started, could be clearly defined. The average grain diameter of the recrystallized matrix is also given in Table I. The grain size did not depend on the temperature of the isothermal anneal. With one exception the grain size of the primary recrystallized matrix was found to be rather fine

Jan 1, 1965

By W. M. Wojcik, R. F. Kowal

Oxygen and sulfur distributions in commercial, 5-ton ingots of killed, medium carbon steel are described. Oxygen distribution is found to vary with deoxidation practice. Irregular distribution of oxygen within ingots makes necessary special precautions in sampling of rolled products for analysis of oxygen. Oxygen distribution is discussed in terms of recently published solidification concepts which had been successfully applied to simpler cases of segregation. These concepts have been found inadequate to explain observed oxygen distributions. Convective movements of the liquid metal, as determined by tracer elements, are shown to be capable of accounting for the observed distributions of oxygen. IN an effort to explore the origin of surface and subsurface imperfections in pierced steel products, a study of oxygen and sulfur segregation was made on ingots cast in open-top and hot-top molds. The results of our previous investigations1"3 have indicated the importance of the location and amount of oxide inclusions in an ingot. Inclusions close to the surface of the ingot have been found to contribute greatly to the formation of imperfections in the surface of finished products. This study of the effects of deoxidation and casting practice on segregation and the resulting oxygen distribution in ingots was initiated to determine the parameters controlling the location of inclusions in an ingot. Segregation of solute elements during solidification of low-melting binary alloys has been studied in the past.1, 5 Formation and growth of inclusions in iron melts have been studied under specific conditions."- In spite of these and other recent studies,10-12 segregation during solidification of commercial, killed steel ingots is not well understood. Consideration of solidification rates, of segregation during solidification of the chill, dendritic, and central zones, and of material balances for the segregated elements has indicated that a simplified theoretical solidification model is not adequate. However, the observed high oxygen contents in localized volumes of the dendritic zone can be rationalized if additional effects of convection currents in the ingots, precipitation, and rapid growth of new phases are considered. EXPERIMENTAL PROCEDURE Steelmaking and Processing. A group of nine killed. medium carbon steel heats having compositions listed in Table I have been studied. The deoxidation and mold practices used were varied to give a wide range of steel oxygen contents. The amounts of aluminum added to the ladle and the ingot casting practices (hot top and open top) were the main variables. The steel was made by a duplex practice in 160-ton tilting basic open-hearth furnaces. All nine heats were top-cast into 24 by 24 in. big end down, fluted molds, to a height between 60 and 76 in., using both open tops and exothermic hot tops. The deoxidation practice and the tapping and teeming details for each heat and ingot studied are given in Tables II and III, respectively. Hot-top practice is indicated by the letter H following the heat designation. Furnace and ladle temperatures were measured by standard disposable-tip, Pt/10 pet PtRh thermocouples. Teeming-stream temperatures were obtained as described by Samways et al.,13 by immersing a Pt/10 pet PtRh thermocouple, covered by a silica sheath, into the teeming stream under the nozzle. The output of this thermocouple was recorded with Leeds & Northrup Speedomax potentiometer. Calibration of the latter thermocouples was based on the freezing point of a pure iron/oxygen alloy (2795°F). The accumulated errors of measurements were within ±10°F. The thermocouple measurements were supplemented in this investigation by continuous recording of a ratioing, two-color pyrometer (Shawmeter), protected from smoke by a blast of clean air within the sighting tube, and calibrated to read with better than ±10°F accuracy. Following teeming of three heats, P, R, and T, tracer elements were added to the steel in the molds to obtain a record of the progress of solidification. As soon as the teeming stream was shut off, a 0.010-in.-thick steel can containing a mixture of crushed standard ferro-titanium and ferro-vanadium (0.05 pet of each alloy element) was plunged into the middle of the steel pool to a depth of 6 in. In about 30 sec no indication of the can or its contents remained. The surface of the open-top ingots solidified in 20 to 30 sec. A study of liquid metal movement and the precipitation of oxides was facilitated materially by use of the tracer technique as titanium has a low distribution coefficient between solid and liquid steel while vanadium has a high distribution coefficient.

Jan 1, 1965

By G. W. Tuffnell, S. Floreen

Three Fe-18 pet Ni-base ternary alloys cortaining carbon, molybdenum, or cobalt were aged to pgroduce hardening by carbides, Ni3Mo, or ordering, respectively. Each alloy was tested in rotating-beam fatigue tests after hardening to 195 ksi ultim.ate tensile strength. The fatigue properties of three alloys soere essentially identical. The results suggest that tlze fatigae strength of an igron-base alloy should depend primarily upon the tensile strength, and not iipon the hardening mecaanisrn. THE fatigue strengths of many steels have been found to be directly related to their ultimate tensile strengths. For ultimate strengths below 200 ksi, the fatigue endurance limit at 107 to 10' cycles is approximately half the ultimate strength. Essen- tially all of the steels that have been tested, however, have been hardened by carbon. Thus it is not certain whether the relationship between fatigue and ultimate strengths is a general one that applies for other types of hardening in steels, or whether it applies only for the particular method of hardening by carbon. The recent developments of the martensitic precipitation-hardened stainless steels and the marag-ing steels have shown that it is possible to achieve high strengths in iron-base alloys by other hardening mechanisms than that due to carbon. Therefore, it is now possible to compare the fatigue properties of several iron alloys having the same ultimate strengths but hardened by different mechanisms. For the present study a binary Fe-18 pct Ni martensite was selected as a base material. This particular matrix composition has been found to have deformation characteristics very similar to pure iron.' Three ternary alloys were prepared by adding either 0.23 wt pct C, 5 wt pct Mo, or 20 wt pct Co to this base composition. After suitable heat

Jan 1, 1965

By I. R. Kramer

In his technical note entitled "Specimen Temperature During Electropolishing of Aluminum Crystals" Dr. Nakada 1 reported that the temperature of his aluminum specimen increased 65°C when it was polished electrolytically; he therefore concluded that the temperature of the specimens used in our work (Nakada's Refs. 4 and 5) similarly increased by this amount. From this he further concluded that the change in plastic-flow characteristics associated with the removal of the surface by electropolishing during deformation were due solely to a temperature effect. We believe that these conclusions are not justified. It is, of course, realized that the amount of the temperature increase is associated with the total heat transfer of the entire system, including the efficiency of the stirring action, heat flow through the grip on the specimen, and so forth. Of great importance is the type of electrolyte used for polishing. For example, in Dr. Nakada's experiment a methanol-perchloric acid solution was used. This solution has a relatively low conductivity and for the current densities of 0.3 amp per sq cm the electrical input appears to be 30 to 40 w. In our research a methanol-nitric acid solution was employed and for the size of specimen employed by Dr. Nakada the power never exceeded 7 w. There are major differences in his test apparatus which also appear to affect the temperature rise. In our case a steel specimen holder was used in order to provide a good heat sink and conduction path. The cathode was a coiled aluminum tube through which a coolant was passed to provide additional cooling capacity near the specimen. The stirrer was placed in a position to obtain maximum cooling efficiency. On the other hand Dr. Nakada used a teflon grip to support the specimen at one end only and a flat copper cathode, thus providing a much less efficient heat sink as compared with the metal grips used in our studies. The relationship between current density and the temperature increase for two sizes of aluminum crystals tested in our tensile apparatus is shown in Fig. 1. The temperature measurements were determined from a thermocouple spot-welded to the surface of the specimen. The thermocouple was protected from the acid by means of a thin film of stop-off lacquer which covered just the outer surface of the thermocouple junction. Whereas Dr. Nakada reported a time delay of 8 min to obtain thermal equilibrium, our temperature change occurred in less than 15 sec. This emphasizes the importance of the heat transfer in the entire system. However, readings were taken over a period in excess of 5 min to assure thermal equilibrium. The data given in Fig. 1 were also confirmed by determining the thermal expansion of the specimen in a sensitive creep apparatus. As may be seen, for the smaller specimen the increase in temperature is 8°C for a current density of 1.5 amp per sq in. The majority of our data on the effects of surface removal on the plastic behavior of aluminum were obtained at current-density values at or below this value. Above 1.5 amp per sq in. the temperature rise is uncertain; for under certain conditions, especially in the vicinity of 2 amp per sq in., the anode can become polarized and the current flow is erratic. This condition is easily recognized because wide fluctuations in the creep and stress-strain data make it virtually impossible to obtain reliable results, a condition which was always avoided in our work. The temperature increase for the larger specimen is less than that of the smaller specimens below a current density of 1.5 amp per sq in. At a current density of 2 amp per sq in. the temperature increased 13°C. The above temperature data are in accord with observations made during creep experiments. Using a transducer capable of detecting strain in the order of 10 -6, temperature changes in excess of 8°C were not detected when the current density was increased to as high a value as 1.5 amp per sq in. Finally we wish to point out that our experimental data cannot be explained solely in terms of a heat effect. For example in Fig. 1 of Ref. 2, the change in slope of Stage II as a function of the rate of re-

Jan 1, 1965

By R. H. Martinson, E. Teghtsoonian

The effects of magnesium-rich surface layers of varying thickness on the mechanical properties of LiF have been studied. The yield stress, critical tensile stress, and work-hardening slope increase linearly with layer thickness, while the strain to fracture decreases yapidly with increasing layev thickness. The results of tension tests, supplemente!d by metallogvaphic evidence, indicate that the yield stress and the critical tensile stress of LiF are not determined by surface-source operation. DESPITE many studies, which have recently been reviewed,' certain fundamental aspects of the plasticity of LiF remain obscure. For instance, one question still awaiting a satisfactory resolution is: does the nucleation and/or multiplication of dislocations in LiF occur preferentially near the crystal surface? It is now well-known that mobile dislocation half-loops are readily introduced into LiF by surface damage and that the mechanical properties of LiF are affected by such surface damage. However, it is not known whether, in the absence of large numbers of artificially introduced surface sources, dislocations arise more easily near the crystal surface than in the crystal interior. It is also not known to what extent existing mobile dislocations multiply preferentially near the crystal surface. The present work is an experimental study of these matters. A short investigation of the effect of magnesium -alloyed surface layers on the deformation of LiF single crystals has been published by Westwood.~ Some relevant conclusions drawn from this research are restated here for purposes of comparison with the present work. 1) Tests on as-cleaved, chemically polished, coated and uncoated specimens show that a coating can restrict the operation of artificially introduced surface sources, thereby raising the yield stress. A coating also decreases the initial rate of work hardening of as-cleaved LiF crystals by reducing the number of active dislocation sources, so that mutual dislocation interference effects are lessened. 2) A catastrophic breakthrough of groups of edge dislocations piled up against the coating can cause stress drops in the stress-strain curve. 3) A surface coating can act as a stable barrier to the egress of dislocations, causing pile-up,

Jan 1, 1965

By S. J. Aresco, A. A. Orning

This paper presents a scheme for determination of precision for the various steps in coal sampling, sample preparation, and analysis. It also gives resulting data on 100 American coals from various coal fields covering all ranks from lignite to low-volatile bituminous. The variances of the various steps of the process are studied as a function of the coal preparation, the top size of coal as sampled, and the ash content. To understand the variables that control the overall precision of coal sampling, sample reduction and analysis, there is need for measurement of the variances due to successive steps of the process. The purpose of the present paper is to study these variances as a function of the coal preparation, the top size of the coal as sampled, and the ash content. The sampling data included are from 100 individual preparation plants from various coal fields in the United States and cover all ranks of coal from lignite to low-volatile bituminous; various types of preparation; deep, strip and auger mines; and, in many instances, mixtures from two or more coal beds or types of mines. The collection of data is continuing, but sufficient work has been done to draw some important conclusions. Since the work was done in conjunction with the Bureau of Mines sampling of coal, the time that could be spent at any one preparation plant was limited. This, in turn, limited the number of samples that could be taken for each coal. For this reason, the results from various plants were pooled in blocks. Two major blocking categories were "raw" and "washed!' The raw coals had no preparation other than crushing and/or screening. The washed coals were prepared further by some combination of wet or dry mechanical cleaning. The categories were further divided ac-

Jan 1, 1965

By T. S. Lovering

The use of geothermal energy is expanding very rapidly. This type of energy has proven commercially profitable for generation of electricity, for space heating, process heating, auxiliary heating of water in conventional steam power plants and for recovery of chemicals contained in natural hot water and steam. Two types of geothermal energy sources are recognized: 1) hot springs in regions of nearly normal heat flow that tap a deep reservoir through which water moves slowly to a hot springs conduit and then rapidly to the surface; 2) hyperthermal areas in which the water is heated by a relatively concentrated heat source related to volcanicity. If there is a geologic trap that provides a geologic analog to a steam boiler, as at Larderello, Italy, the hyperthermal area will have a convecting system that develops superheated water at relatively shallow depth and may provide natural steam in large quantities. If a hyperthermal area is to be productive for a long time, the underflow into the reservoir should be slow enough to allow the heat source and convective system to heat the underflow to the working temperature, and the production rate must not exceed this rate of underflow. A model based on a typical aquifer suggests that the rate of movement of water through the reservoir be such that a few years are spent in transit between isotherms that are spaced about 2°F apart. The possibility of finding blind geothermal areas is illustrated by discussion of the techniques developed in evaluating the subsurface temperatures in the East Tintic district of Utah where a map of isotherms at water level (2000 to 2000 ft below the surface) shows that a hyperthermal area may exist a short distance southeast of the mining district. Very nearly all of the energy that man currently uses comes ultimately from the sun's radiation. This includes water power, fuels such as wood, peat, coal and petroleum, the wind and all our animal power. In the paper summarizing a conference on solar energyl6 the average amount of solar energy received daily on the earth is taken at about 1 cal per m2 per min or slightly less than 2 pcal per cm2 per sec; this is almost exactly the amount of energy on the average that the earth liberates in regions of normal geothermal gradient due to its own internal heating. In many places, however, the energy released is many times the average and in some of these hyperthermal areas, geothermal steam is used for generation of electricity, and hot springs are used for heating buildings and private dwellings, process heating, auxiliary heating of water in conventional steam power plants, and chemicals may be recoverable from both hot water and steam. The use of hot springs waters for heating houses goes back hundreds of years but until recently was confined to a few dwellings close to the hot springs. In Korea, some houses had hot spring water channeled through conduits in the floor centuries ago and thus the Koreans can be credited with pioneer development of radiant heating. In Iceland at present nearly a third of the population uses natural thermal water for domestic heating." The Reykjavik system pipes hot spring water at about 94°C throughout the city and has devised insulated double pipes that allow the water to be piped for some 25 km with a drop of only 1°C for every 5 km. The actual cost to the Icelandic consumer is only one-third the cost of heating by imported coal and yet the industry is one of the most profitable in Iceland. The most profitable use of geothermal energy has been its conversion into electricity which can be transmitted economically much greater distances than hot water. The largest installation at the present time is that at Larderello, Italy, where the Count of Larderello began to experiment in the production of electricity from geothermal steam 60 years ago — in 1904. He installed his first steam turbine, with a capacity of only 250 kw, in 1912 as the result of a local quarrel with the power company which furnished the current required in the Larderello chemical industry - an industry that then dated back nearly a century. As experience was gained in drilling deep holes to tap geothermal steam and in converting it to electric power, the capacity of the installation of Larderello gradually increased, but was all destroyed by the Germans during their retreat from Italy in the closing

Jan 1, 1965

By R. P. Agarwala, M. S. Anand, S. P. Murarka

Using the residual radioactivity technique, grain boundary diffusion of chromium in the a phase and lattice diffusion in the a and the ß phases of zirconiuttz have heen studied. The diffusivities (in units of square centiwzeters per second) have heen expressed as THE rate of diffusion of atoms along surfaces and grain boundaries is much greater than that through the bulk of metals. These processes dominate over volume diffusion at lower temperatures and can be isolated from it by virtue of the fact they yield high diffusion coefficients and low activation energies. In zirconium, however, it has been difficult to identify the volume-diffusion parameters separately. The values of diffusivities in a and ß phases yield activation energies which are much lower than those for volume diffusion calculated from semiempirical or empirical relations.1-3 Diffusion frequencies (Do) are also low, especially in the a phase, and they yield negative entropies of activation.4 Similar results have been obtained for the diffusion studies in titanium.5 It is yet to be investigated whether in these cases diffusion is controlled by some extraneous factor like high reactivity of these metals, or some different mechanism is operative as suggested by Pound et a1.6 However, Federer and Lundy,7 in a recent study of the diffusion of zir-conium-95 and columbium-95 in zirconium, and Gibbs et al.,8 in the study of impurity diffusion in titanium, found a temperature dependence of D, Do, and Q. Their study yielded low values of Do and Q at low temperatures, as has been found by other workers, and fairly high values at high tempera-

Jan 1, 1965

By A. Bar-Or, G. Hirsch, P. Wynblatt

The rates of nucleation and growth at various temperatures in U-0.45, 1.1, and 1.35 at. pet Cr alloys were determined, providing a qliantitative proo-f that the transformation under study is controlled by nucleation and diffusional growth. The linear relationship, log NV US 1/T2, suggests that the nucleation of the alloys is probably controlled by the classical law of nuclei formation. The established linear relationship, log G us 1/T, in the three alloys suggests that growth in each alloy is controlled by bidimensional nuclei formation. It was established that both the rate of growth and the rale of nucleation are retarded by chromium. Using classical formulae, sApN and sApG (the ß/a apparent interphase energies) were calculated from the nucleation and growth results, respectively. It was found that their values are of the same order of magnitude. However, why sApN is unaffected by the addition of chromium, while sApG increases with it, needs further clarification. In the last few years, an appreciable amount of research has been conducted into the mechanism of the ß —a transformation in uranium and uranium alloys. Recently, Burke1 reviewed the published work concerned with this transformation. He concluded that in dilute uranium alloys it is firmly established that, at high temperatures in the a range, ß decomposes by a diffusional process. However, he based his conclusion either on experiments dealing with the over-all transformation,2 or on qualitative results.3 It was felt that a better insight into the mechanism of transformation could be gained from the quantitative measurements of growth and nucleation rates, which we aimed to carry out in the present work. In a previous work,4 we have shown that a rim of a grains is formed during isothermal transformation at the high a range. The thickening of this rim with time was used for measuring the rate of growth in U-0.45 at. pet Cr.5 This method was also used in the present work for two more alloys. For meas- uring the rate of nucleation, a method of quenching and macroetching was developed to allow differentiation between a grains formed during isothermal transformation at the upper a range and those formed after quenching, Fig. 1. EXPERIMENTAL WORK AND RESULTS The Rate of Nucleation. The measurement of the rate of nucleation was effected by counting the num-

Jan 1, 1965

By P. M. Blair

This paper presents numerical solutions of the equations describing the imbibition of water and the countercurrent flow of oil in porous rocks. The imbibition process is of practical importance in recovering oil from heterogeneous formations and has been studied principally by experimental means. Calculations were made for imbibition of water into both linear and radial systems. Imbibition in the linear systems was allowed to take place through one open, or permeable, face of the porous medium studied. ln the radial system, water was imbibed inward from the outer radius. The effects on rate of imbibition of varying the capillary pressure and relative permeability curves, oil viscosity and the initial water saturation were computed. For each case studied, the rate of water imbibition and the saturation and pressure profiles were calculated as functions of time. The results of these calculations indicate that, for the porous medium studied, the time required to imbibe a fixed volume of water of a certain viscosity is approximately proportional to the square root of the viscosity of the reservoir oil whenever the oil viscosity is greater than the water viscosity. Results are also presented illustrating the effects on rate of imbibition of the other variables studied. INTRODUCTION The process of imbibition, or spontaneous flow of fluids in porous media under the influence of capillary pressure gradients, occurs wherever there exist in permeable rock capillary pressure gradients which are not exactly balanced by opposing pressure gradients (such as those resulting from the influence of gravity). The importance of such capillary movement in the displacement of oil by water or gas was recognized in early investigations and described by Leverett, Lewis and True in 1942.1 Methods advanced by these authors for studying the process using dynamically scaled models were rendered more general and flexible by the research of later workers.2"4 The influence of capillary forces in laboratory water floods has also been discussed by several authors.6 While imbibition plays a very important role in the recovery of oil from normal reservoirs, Brownscombe and Dyes7 pointed out that imbibition might be the dominant displacement process in water flooding reservoirs characterized by drastic variations in permeability, such as in fractured-matrix reservoirs. In water-wet, fractured-matrix reservoirs, water will be imbibed from fractures into the matrix with a countercurrent expulsion of oil into the fractures. If the imbibition occurs at a sufficiently rapid rate, a very successful water flood can result; if the imbibition proceeds slowly the project might not be economically attractive. Scaled-model studies have demonstrated the vital importance of imbibition in secondary recovery in fractured reservoirs. It is therefore important in the evaluation of waterflooding prospects to develop a thorough understanding of the quantitative relationships of the factors which control the rapidity of capillary imbibition. The imbibition process serves reservoir engineers in still another important way by providing a technique for studying the wettability of reservoir core samples. Such experiments are usually conducted by observing the rate of expulsion of oil or water from core samples submerged in the appropriate fluid. Several papers have been 'published on the experimental techniques involved?? Although Handy has recently published a method for calculating capillary pressures from experiments with gas-saturated cores, it has not yet been possible to deduce quantitative information regarding water-oil relative permeability and capillary pressure characteristics of the rock from the experimental results. Thus a technique is needed for studying the quantitative dependence of imbibition rate on oil and water viscosity, initial water saturation, relative permeability-saturation, and capillary pressure-saturation relations. The development of such information, including saturation and pressure profiles by laboratory experiments, would be very difficult. However, recent advances in the applica-

Jan 1, 1965

By W. D. Potter, E. T. Peters

Single-crystal films of boron have been prepared by thermal reduction of purified BCl3 with subsequent deposition of boron onto oriented silicon substrates. The structural properties of films prepared under various experimental conditions have been determined by light microscopy and "grazing-" angle electron-diffraction analysis. An orientation relationship is proposed for silicon and a-rhombo-heclral boron which can be described as: (111)Si (112)B; [110]Si [110]B ElEMENTAL boron is known to have a high energy gap* and consequently has potential value as a high- *Energy gaps of 1.35 to 1.55 ev1-3 and 2.0 ev4 are reported for ß-rhombohedra1 and a-rhombohedra1 boron, respectively. temperature semiconductor.5 Exploitation of this property, however, has been hampered by the unavailability of pure, single-crystal boron. Two recent symposia6,7 have treated the subjects of purification, crystallization, crystal structure, and properties of boron and provide a foundation for further investigation in these areas. The structure, lattice constants, and observed ranges of stability for four reported polymorphs of boron8°12 are listed in Table I. The rhombohedral structures are generally accepted as true polymorphs whereas the tetragonal structures are believed to be transient. Most currently available single crystals of boron are grown by Czochralski and floating-zone methods. Unfortunately, these crystals generally contain numerous defects and are highly strained as a result of the polymorphic transformations and thermal contractions which accompany cooling. Recently, Armington, Potter, and Tanner13 have reported the formation of single-crystal films of boron on silicon single-crystal substrates through thermal reduction of BC13 by hydrogen. Boron E. T. PETERS, Junior Member AIME, is Senior Scientist, ManLabs, Inc., Cambridge, Mass., and W. D. POTTER is Research Chemist, Air Force Cambridge Research Loboratories, L. G. Hanscom Field, Bedford, Mass. formed at temperatures near 1000°C was identified as the a-rhombohedra1 form, consistent with expectation. The present paper describes a continuation of this work with specific reference to characterization of the structural features of the deposited boron films. The ultimate goal of this program is to produce large-area single-crystal films of a-rhombohedra1 boron to be used for electrical and optical studies by other workers. PREPARATION OF BORON FILMS The experimental procedure of preparing thin films is similar to that described by Theuerer14 for the epitaxial growth of silicon by thermal reduction of SiC14 and has already been described as applied to the preparation of boron thin films.13 Briefly, a mixture of BCl3 vapor and dry hydrogen is passed over a heated silicon substrate. Thermal reduction of the BC13 allows deposition of boron on the substrate. The substrates were {111}-oriented, single-crystal silicon. Square specimens 0.63 cm on an edge by 0.1 to 3.0 mm thickness were prepared by

Jan 1, 1965

By C. J. Leith

A study of the effect of rock composition, rock structure and degree of weathering on the stability of cut slopes is being sponsored jointly by the U.S. Bureau of Public Roads and the North Carolina Highway Commission. In 58 mountain and piedmont counties of North Carolina the percentage of failed cut slopes is greatest in micaceous metasediments, gneisses, and schists, and in saprolite and soil derived from these rock types. Soil slope failures outnumber rock slope failures by two to one. Joints and similar planes of separation exert a strong influence on size and shape of the sliding mass. They may or may not act as failure surfaces, depending on their orientation with respect to the active forces. Climatological data, though indicative of weathering conditions, do not correlate well with slope failure frequency. Because of the presence of joints and similar planes of weakness in soil and rock materials, conventional methods for analyzing slope stabilities are not directly applicable. Empirically derived modifications of these methods are being investigated. A study of the stability of highway cut slopes, sponsored by the U. S. Bureau of Public Roads and the North Carolina State Highway Commission, began in 1962 at North Carolina State of the University of North Carolina at Raleigh. As part of this study all slides, rockfalls and other types of cut slope failures on Federal and State highways in the 58 mountain and piedmont counties in North Carolina were located and described, and the data catalogued in a punched card file system. A major objective of the project is to relate slope failures to properties and physical conditions of the geological units in which the slopes were constructed, and to correlate soil type and/or geological unit with type and frequency of slope failure. The complexities of the problem of slope stability and the limitations which these complexities impose on methods for analyzing slopes have been recognized for many years. A great variety of factors and processes may lead to slides, often making it almost impossible to analyze theoretically the conditions required for stability of slopes. One of the principal factors determining maximum safe slopes is the shear strength of the material in which the slopes are cut, but unfortunately there are very few data available concerning shear strengths of residual soils. Vargasl tested clay derived from gneiss and granite in southern Brazil; the properties of decomposed granite occurring near Hong Kong were determined by Lumb.2 These data are being used, when applicable, to supplement the test data obtained in the present study by Yorke.3 The locations of the North Carolina slope failures, more than 400 in number, are shown on Fig. 1. This map, adapted from the Geological Map of North Carolina,4 suggests the possibility of a relationship between frequency of slide occurrence and rock type. However, the evaluation of this possibility requires consideration not only of the type of rock, but also of its large and small scale structural features, its susceptibility to and degree of weathering, and the composition and structure of the weathering products. Soil slope failures in thoroughly weathered soil material and saprolite outnumber rock slope failures two to one. INFLUENCE OF ROCK TYPE The agricultural soil type involved in each soil slope failure was identified and each failure was catalogued in terms of the parent material from which the soil was derived. These data indicate that most of the slope failures, whether in the rock or in the derived soil, are associated with metamorphic rocks (see Fig. 2a). The data may be skewed somewhat because of the relative sizes of the total areas underlain by the various rock types, but Leith and Gupton5 have demonstrated that the preponderance of failures in metamorphic rocks is of much greater magnitude than could be accounted for by the areal factor alone. The dominance of metamorphic rocks is emphasized when soil slope failures are considered in terms of the specific rock types from which the soils were derived (see Fig. 2b). In particular, mica schists and mica gneisses account for more slides

Jan 1, 1965

By G. B. Gibbs

Equations are developed which describe the growth of graphite nodules during a-phase graphi-tization under various limiting conditions. When carbon diffusion is rate controlling a modified Birchenall-Mead equation predicts an "apparent activation energy" of 31 keal mole-1 for the parabolic growth constant. If reaction at the graphite or cementite interface is the slowest process growth kinetics are linear and the activation energy must be >50 keal mole-1. A comparison with the experimental observations confirms that carbon diffusion determines the rate of reaction. Apart from minor approximations associated with Lou! carbon solubility in a iron the theory may also be applied to I-phase graphitization. It is well-known that cementite exists as a meta-stable phase in Fe-C alloys and that if a specimen is held at diffusion temperatures for a sufficiently long period the cementite decomposes and graphite is precipitated. Graphitization occurs by the nu-cleation and growth of carbon nodules and several impurity elements catalyze the reaction when present in minute amounts.1 The growth kinetics of nodules have been extensively investigated in the fee ? range2, 3 and recently in the bee a range.1, 4 Experimental growth curves (nodule radius vs time) have slopes which decrease with time, but the scatter of points is generally too great for the precise shape of the curves to be determined with certainty. Birchenall and Mead5 concluded that they are parabolic in ? iron, indicating diffusion control, and that the migration of carbon atoms determines the growth rate. Harris et al, 1 advanced similar ideas for a-phase graphitization. On the other hand, Burke3 chose to regard the initial growth region as a transient followed by a linear law indicating interface control (i.e., cementite decomposition or graphite crystallization as the slowest process). Jeminson and Higgins6 originally considered the initial steep portions of "parabolie" to represent a linear growth law also. In support of these ideas it has been noted that the apparent ac-

Jan 1, 1965

By Nicholas J. Grant, Michio Yamazaki

Cast copper alloys containing 10, 20, and 30 pct Ni and 0.75 to 0.80 pct Al were machine-milled into chips, then comminuted in a rod mill to fine flake powder utilizing a number of processing variables. The powders here internally oxidized, mostly at 800°C, in a low-pressure oxygen atmosphere. The consolidated powders were hot-extruded into bar stock. Room-tenmperature tension tests, stress-rupture tests mostly at 650°C, but also at 450° and 850°C, and hardness measurements after various annealing temperature treatments to study alloy stability were perfomted. Excellent room-temperature strength, high rupture strength at 650°C, and resistance to recrystallization at 1050°C were obtained. Problems in optimizing conditions for internal oxidation of Cu-Ni base alloys are discussed. THE interesting high-temperature properties of SAP' have stimulated considerable effort in the study of more refractory alloy systems where the potential for high-strength alloys at high temperature is great.2-13 A number of methods have been utilized to produce the desired fine, hard particle dispersions, of which internal oxidation2,9,7 of dilute solid-solution systems offers considerable promise by virtue of the potential for producing ultrafine, well-dispersed oxides. While most of the published works are concerned with pure metal matrices, a number of investigators have studied the effects of solid-solution strengthening.10,19 Use of more complex alloy matrices (for example, aging systems) has been unsuccessful because overaging still occurs at high temperatures in the oxide-containing alloys.14.15 Solid-solution strengthening is, however, effective at very high temperatures9,10 and might be expected to contribute importantly to the strength of oxide-dispersion strengthened alloys. For this study, internal oxidation of solid-solution alloys of copper and nickel, containing small amounts of aluminum, was chosen as the method of alloy preparation. PREPARATION OF ALLOYS Three copper alloys containing about 10, 20, and 30 pct Ni and each containing 0.75 to 0.80 pct A1 (enough to yield about 3.5 vol pct alumina) were prepared as air-cast ingots measuring 2.5 in. diameter by 6 in. high (see Table I for the analyses). Processing steps for all the alloys were as follows (also see Table 11): 1) Homogenization of the ingot at 982°C (1800°F) for 45 hr in an argon atmosphere. 2) Machine milling of ingots into fine chips. Average thickness was about 0.1 to 0.2 mm. 3) Hydrogen reduction of chips at 593°C (1100° F) for 1 hr to reduce copper and nickel oxides. 4) Rod milling of chips to finer powders. 5) Hydrogen treatment of powders as in step 3. 6) Internal oxidation of the powders. 7) Hydrogen treatment of oxidized powders as in step 3. 8) Hydrostatic compression of evacuated powders. 9) Sintering of compacts in hydrogen. 10) Hot extrusion. Variations in processing among the alloys were made in steps 4, 5, and 10 (see Table 11). In the past, two methods were utilized to internally oxidize alloy powders. Preston and Grant3 surface-oxidized dilute Cu-Al powders to obtain the necessary amount of oxygen to oxidize the solute metal (aluminum and silicon), and then permitted the formed copper oxide to diffuse and react with the solute in an argon atmosphere. Bonis and Grant4 exposed Ni-A1 and other nickel alloys to an oxygen pressure derived from the decomposition of nickel oxide at a preselected temperature, in an argon atmosphere. Both methods are applicable and can be modified to generate a range of oxygen pressures for oxidation of the solute but not the solvent metals. Procedure I: Surface Oxidation of Alloy A3, Cu-10Ni-0.76A1. Powders of -20 to +28 mesh were surface-oxidized at 500°C (932°F) to obtain the desired amount of oxygen for oxidation of the aluminum to alumina; the powder was then sealed in Vycor and heated at 900°C (1652°F) for various times up to

Jan 1, 1965

By R. Schuhmann

Jan 1, 1965

By Claude H. P. Lupis, John F. Elliott

IN the recent past, extensive use has been made of the interaction coefficient in treating the thermody-namic behavior of components in solutions at dilute concentrations. The development of this concept by wagner1 and chipman2,3 is sufficiently well-known so that its elucidation here is not necessary. The present discussion is focused on the relationship between E and e. The usual forms of the coefficients are:

Jan 1, 1965

By H. C. Chao, L. H. Van Vlack

Small MnS inclusions with known crystallographic orientations were placed inside powder compacts of low-carbon steel. After the metal was axially campressed with negligible end friction, the deformstions for the metal and the inclusions were compared. The MnS inclusions deformed more when the [100] direction was aligned with the compression axis than when the [111] direction was parallel to this axis. The deformations of the inclusions in the two principal radial directions were equal for each of the above orientations. Inclusions with [110] compression alignments did not deform with radial symmetry. The relative deformation of the inclusion and metal was closely dependent upon the relatiue hardness of the two phases. The relative deformation of the two phases was not sensitive to the rate of deformation. RECENT studies by the authors1.&apos; suggested that the plastic deformation of MnS in steel would probably be highly sensitive to the orientation of the inclusions and to the temperature of the metal. This paper reports an investigation of these factors upon MnS behavior within steel. Manganese sulfide (MnS) possesses an NaCl-type structure and typically has extensive (l10) {110} slip as a separate (noninclusion) crystal.&apos; A secondary slip system, ( l 10) { l l l}, has also been observed where the major slip system is restricted. In general, MnS inclusions must be rated as a highly deformable second phase.3 The amount of sulfide deformation varies, however, with several composition and processing factors. Some of these have been only partially assigned. For example, it is known that minor amounts (<0.01 pct) of silicon within free-machining steels will increase the amount of MnS deformation,4 but the mechanism of the added deformation can only be surmised at the present. Manganese sulfide and steel have sufficiently comparable deformation characteristics so that slip which is started in steel may be continued through the sulfide inclusions and back into the steel if the crystal orientations are favorable.5 A more detailed discussion of previous work on the plastic deformation of NaC1-type crystals and on the plastic deformation of inclusions within a metal is given in Chao&apos;s work.6 EXPERIMENTAL PROCEDURE The manganese sulfide which was used in this study was prepared by previously described methods.&apos; Single crystals of MnS, both as cleavage cubes and as spheres, were oriented within steel powder compacts so that the desired crystal directions were parallel to the direction of axial compression. A four-stage hydrostatic compaction procedure was used and involved the following steps. In the first stage part of the powder was placed in a metal die 1 in. in diameter with a thick (1 in. OD, 5/8 in. ID) rubber liner which had one end plugged. The steel powder was hand-rammed, making it as dense as possible before placing a carefully sized MnS crystal (either as a sphere or as a cube) near the center. The crystal was oriented with the chosen direction vertical; viz., [001], [011], or [111], with the aid of a X10 microscope. A pair of tungsten wire threads 0.020 in. in diameter was inserted along the side of this (&apos;core compact" to locate the desired plane after the compression tests. After the crystal was positioned in the center of the die, more powder was added and carefully rammed by hand. The die was then capped with a rubber plug of the same hardness and thickness as that of the liner. The whole assembly as shown in Fig. 1 was compacted by a ram load of 54,000 lb (about 70,000 psi). In the second stage a smaller, 3/4-in, rubber-lined die was used to give a stress of approximately 120,000 psi. The above process was repeated with the initial compact serving as a core for a larger compact. The final product after sintering was a cylinder 1 cm long and 1 cm in diameter, having a density of 7.54 g per cu cm. This was close to the theoretical density since the metal contained a non-metallic phase. There was no evidence of MnS deformation during the hydrostatic compaction or subsequent sintering. Elevated-temperature hardness data were obtained by procedures previously described.2 Compression tests for inclusion deformation utilized the cylinders which were described above. The critical problem in these tests was the lubri-

Jan 1, 1965

By G. C. Gould

The nature of freckle segregation is determined by chemical analyses, microradiograplzy, and electron microprobe. In addition, the influence of chemistry variation on freckle formation is studied in laboratory experiments. It is demonstrated that controlled laboratory solidification can produce freckle segregation. "FRECKLE" segregation is a discontinuous, rod-like segregation observed in vacuum consumable -electrode ingots. Figs. 1 and 2 show this segregation as pencil-like, dark etching areas of segregation found in longitudinally and transversely sectioned consumable-electrode ingots. In most cases the segregation is perpendicular with respect to gravity and parallel to the longitudinal axis of the ingot. The segregation is most often encountered at or around midradius in large vacuum consumable-electrode ingots. The austenitic alloys, A-286* and 718, † have shown particular susceptibility to this mode of segregation. The origin of these rodlike pockets of segregation is not well-understood. It has been suggested that the segregation is a result of a portion of the electrode dropping off and passing through the arc un-melted. On the other hand, smeltzer1 has demonstrated the freckle segregation in A-286 to be comprised of low-melting constituents relative to the matrix melting point. The discontinuous nature of the segregation is the most perplexing facet of the freckle-segregation problem. It was to gain some insight into this part of the problem that this work was undertaken. I) PROCEDURE At the outset of the work identification of the phases present and the elements concentrated in the freckle areas was undertaken. The electron micro-probe was used to determine general enrichment as well as -specific phase chemistry. Microradiography using cobalt and chromium radiation was also used on the freckle areas again to determine the elements involved in the segregation. X-ray diffraction patterns were made from portions of these same foil specimens. Optical microscopy and Knoop hardness measurements were also used

Jan 1, 1965

By J. C. Uy, T. E. Davidson, A. P. Lee

The effects of hydrostatic pressures to 26 kbars on the micro structure of poly crystalline Cd, Zn, Bi, Sn, Zr, Mg, Cu, and Fe were examined. Pressure-induced microscopic plastic flow in the form of boundary migration, slip, multiple slip, or twinning has been observed either singly or in combination in cadmium, zinc, bismuth, and tin. No deformation was observed in zirconium, magnesium, copper, and iron. The occurrence of such deformation, in terms of its intensity and initiation pressure, relates directly to the degree of anisotropy in the linear compressibility. Pressure cycling does not significantly affect the residual tensile properties of a zinc alloy which exhibits pressure-induced deformation similar to that of Pure zinc. CLASSICAL elastic theory predicts that a superim-posed hydrostatic pressure will not induce shear stresses in an ideal material. Thus, in the case of homogeneous and isotropic materials plastic flow will not occur as a result of pressure exposure, regardless of the magnitude of the pressure, unless some form of phase transformation or permanent density change takes place. However, real materials, such as metals, often vary considerably from the condition of homogeneity and isotropy. For this reason, Vu and johannin&apos; examined and observed hydrostatic pressure-induced microscopic deformation in polycrystalline cadmium and, zinc, but not in aluminum, to pressures of 9 kbars and indicated that its occurrence is related to anisotropy in the linear compressibility. In addition, Davidson and omaan&apos; have reported pressure-induced deformation in polycrystalline bismuth below the 1-11 transformation pressure. It becomes apparent then that under certain conditions hydrostatic pressure can induce localized internal shear stresses in some metals of sufficient magnitude to cause plastic flow. In this work, eight pure single-phase metals of different lattice structures and degrees of elastic anisotropy were examined after pressure exposures of up to 26 kbars in order to establish the conditions under which such deformation occurs and the type of deformation as a function of pressure. The effects of pressure cycling on the mechanical proper-

Jan 1, 1965