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By D. C. Porterfield

-4s a result of world economic conditions brought about by the European war and the consequent loss of European markets for Venezuelan crude, production in Venezuela decreased from 205,433,000 bbl. in 1939 to 186,134,000 bbl. in 1940, or 9.4 per cent. This decline interrupted a steady annual increase in production from the country, which had been continuous since 1932. The Maracaibo Lake Basin in western Venezuela accounted for 135,738,000 bbl. during 1940, or 73 per cent of the total, as against eastern Venezuela's production of 50,396,000 bbl., or 27 per cent of the total for the country. Corresponding figures for 1939 were 167,688,000 bbl., or 82 per cent of the total from the Lake Basin, and 37,744,000 bbl., or 18 per cent from the east. These comparisons indicate the increasing importance of the newer developments in eastern Venezuela, but it is unlikely that production from the east will ever exceed that from the west. In all, 429 new wells were completed in Venezuela during 1940 as compared to 438 in 1939. Of the 1940 completions, 411 were in proved areas and 18 were wildcats; 389 (95 per cent) of the exploitation wells drilled were producers, I was a gas well and 21 were dry holes, while 6 (33 per cent) of the wildcats drilled were producers and 12 (67 per cent) were dry; 263 of the exploitation wells and 4 of the wildcats were in the Maracaibo Lake Basin and 148 exploitation wells and 14 wildcats were in the east, mostly in the State of Anzoategui. Exploration of the light-oil areas in the Tia Juana and Lagunillas fields was con- tinued during the year with favorable results. Along the lake shore, drilling was confined principally to fill-in wells in the La Kosa, Tia Juana, Laguni!las and Bachaquero fields, although some important semi-exploratory drilling was done in the Pueblo Viejo area between Lagunillas and Bachaquero fields. The only other developments of particular interest in the west were the successful completion of a step-out well lying to the south of the former southern limit of Mene Grande field, which adds considerable proved acreage to that field; and the failure of two wildcats that were drilled near the Colombian border to determine whether the productive sands of the Petrolea field in the Barco Concession, Colombia, extend into Venezuelan territory. In the east no new fields were discovered during the year, although five wildcats were completed as producers. Three of these, Guario No. I, Jusepin-56 and LM-I (Leona), discovered important extensions to previously proved areas and the remaining two, Cantos-I and SPN-I (Socororo), were subsequently proved to have tapped productive sands of small extent and operations in these areas have been suspended. The greatest activity in the east was concentrated in the Oficina field, where 105 new wells were completed during the year and production was increased from 22,540 bbl. per day during January 1940 to 57,627 bbl. per day during December 1940. Important development work was also done in the Jusepin. San Joaquin and El Koble fields.

Jan 1, 1941

By Frank Woolley, John F. Elliott

The high-temperatzrre solulion calorimeter has been modified and an extension dynamic analysis of i/s transient behavior has provided an improved basis for interpretation of the experimental data. The partial molar heats of mixing, Him, of aluminum, copper, and silicon in the corresponding binary liquid-iron alloys haz,e been measured directly at 1600°C. The resulls in kcal per g-alom are: THE high-temperature1,2 solution calorimeter was planned and designed to obtain directly the partial molar heat of solution of solute elements in a variety of liquid metals that might serve as solvents. Results obtained in the temperature range of 1000° to 1200°C were described in the two earlier publications. However, a long-term goal in the program has been to obtain measurements at 1600°C of the heats of solution of elements in liquid iron. Hence, subsequent to the work last reported,2 and based on the experience of the first two studies, the calorimeter system has been modified and operated at 1600°C to obtain measurements on the Fe-Si, Fe-A1, and Fe-Cu systems. This paper reports the results of the work. THE MODIFIED SYSTEM It was clear from the earlier results that the most precise measurements can be obtained with the high-temperature solution calorimeter when the solution time is less than approximately 30 sec. This can be achieved only when the bath temperature, Tb, is above the melting (liquidus) temperature of the solute material being dropped. A second point of importance is that the refractories used earlier liberated considerable water so that, although the oxygen potential of the atmosphere in the system was satisfactory for work with nickel and copper, it was too high for work with iron, silicon, and so forth. Third, for work at 1550" to 1650°C, an improved temperature-sensing unit is needed. The modified calorimeter in its latest form is shown schematically in Fig. 1. The refractory brick and the Kanthal-Super heating elements which formed the furnace for heating the solute addition have been replaced by a smaller resistance-heated tube furnace which is insulated with bubble-grain alumina. A more elaborate gas purification system was installed to maintain a satisfactorily low oxygen potential in the atmosphere of the calorinieter chamber and of the solute addition furnace. As was the case in the earlier work, at the beginning of a run the calorimeter chamber is evacuated and then filled with purified argon so that a total pressure slightly in excess of 1 atm is maintained during the run. When either of the furnaces within the enclosure attains a red heat, 6 pct H2 is added to the atmosphere. Throughout the run, gas from the calorimeter is circulated through a de-oxidation system at the rate of approximately 5 standard liters per min; the changeover time for the gas within the calorimeter enclosure is approximately 1 hr. The essential stages of the deoxidation system are: a copper furnace at 600°C, a drying tube of anhydrite, and a titanium furnace at 800°C. The oxygen potential of the gas stream in the purification system can be monitored at several points by a stabilized zirconia oxygen cell. This cell is invaluable in detecting the development of leaks in the container vessel and the gas circulating system. The installation of the solute-addition furnace allowed either of two types of temperature-sensing units to be mounted on a water-cooled probe. The first was a small optical device using a silicon diode that was

Jan 1, 1968

By W. D. Brown

Newell1 has shown that hydrogen is removed from steel in a vacuum at a temperature of 500° to 900° C. within 136 hr. Holm and Thompson2 also state that, especially when the hydrogen is high, the results obtained by extraction at 800°C. agree with those obtained by vacuum fusion. In this laboratory, an apparatus was designed (Oct. 1940) and used wherein the piece of steel was heated at 100°F. in a Torricellian vacuum (Fig. I). This vacuum was produced by raising and lowering a "leveling bottle or bulb" containing mercury. Extensive tests were made with this apparatus. The weakness of the procedure was the possibility of leakage of air into the tube around the rubber stopper at the bottom. The present plan is to raise and lower the mercury by applying vacuum from a mechanical pump, to the top of the tube and then to a bottle of mercury in which the bottom of the tube is immersed. The Apparatus The extraction is made in a Pyrex tube A, I in. i.d. It is drawn down at the top at N and an 8-mm. i.d. tube is welded on it. To this a Pyrex stopcock E, 4-mm. bore, is welded. The tube between N and E is preferably 8 in. long and is graduated in milliliters. The total length of A to the stopcock is 28 inches. Heat is supplied by a split-tube furnace. 0 in. long by 6 in. in diameter, hinged on the side. A short pyrometer couple P is inserted through the walls of the furnace. In use this is standardized by comparison with a standard couple in a I-in. tube, just opposite this short couple. In the I-in. tube A a glass rod R, flattened at one end, supports the sample S and the length of R is so adjusted that the sample is in proper place. (It may float in the mercury.) The tube A is inserted in a large rubber stopper to fit the bottle B and extends nearly to the bottom. It is so placed that the handle of the stopcock faces the operator. In this stopper also is a short tube D bent at right angles for the application of vacuum. Also in this stopper is an 8-mm. tube C reaching nearly to the bottom; this is bent so as to miss the heater and again bent to bring it in line with and as near as possible to the upper and reduced part of tube A. On this tube C is welded a 4-mm.-bore stopcock H. In order that the two tubes may be as near as possible, the handle of stopcock H faces the back, or is the reverse of stopcock E. The total height of II above the bottom is 28 in. As explained later, this is the pressure-measuring tube C. The bottle B is about half full of mercury. The mercury must cover the ends of A and C by about I in. when they are evacuated. By tees T and rubber tubing RT the tubes A, C and D, the latter through the three-way stopcock M, are connected to the vacuum pump. A sheet of asbestos at K prevents the stopcocks E and H from becoming so hot that they leak. A good grade of vacuum lubricant prevents leakage of E and H. In order to measure the pressure of the

Jan 1, 1945

By Howard Martens

HARDNESS behavior of commercial titanium alloys following various heat treating processes has been studied for some time. However, the hardness of such alloys following a definite measured cooling rate from the single phase region has not been reported. Therefore, a series of experiments was conducted on the three alloys: 150A, 130A, and 130B. The manufacturers gave the following compositions for these alloys: Ti 150A: 0.045 pct C, 0.069 pctN, 1.32 pct Fe, and 2.68 pct Cr; RC 130A: 0.13 pct C, and 7.9 pct Mn: RC 130B: 0.11 pct C, 3.5 pct Mn, and 3.2 pct Al. A method similar to that described by Greninger was used, and the cooling rate varied from 6" to 9200°F per sec. The cooling curves were recorded on a Speedomax or a recording oscillograph, and the hardness was measured on a Tukon hardness tester using a Knoop indentor with a 500 g load. The curves of Fig. 1 show the variation of the Knoop hardness with the cooling rate for these three alloys. The alloys 150A and 130B behave in a similar manner, as shown by the curves. In neither of these alloys was it; possible to retain the high temperature ß phases, even by the most rapid cooling rate. The slow cooling rates produced a typical ß structure in these two alloys. As the cooling was increased, these phalses became more finely dispersed and the hardness increased. This increase in hardness with cooling rate continued for the 150A alloy until a rate of approximately 1000°F per sec was reached. At this cooling rate, the structure showed the result of the dif-fusionless transformation of the ß solid solution to the supersaturated a solid solution which is referred as a'. At higher cooling rates, the structure showed no change and the hardness showed no great change, as indicated by the dotted portion of the curve in Fig. 1. The increase in hardness with cooling rate for the 130B alloy continued until a rate of approximately 300°F per sec was reached. At this rate the structure was produced and the maximum hardness was reached. Higher cooling rates caused no marked changes in hardness or structure. The alloy 130A behaved in a different manner, as shown by the curve of Fig. 1. The structure produced by the slow cooling rates was a typical a+ß structure, which became finely dispersed as the cooling rate increased. The hardness increased with cooling rate and reached a maximum at approximately 120°F per sec. At this maximum hardness, the structure showed the result of the decomposition of the ß phase into the ß + phase structure. As the cooling rate was increased, smaller amounts of the at phase were formed and the alloy became softer. For cooling rates of 3000°F per sec and above, the high temperature ß phase was retained and the hardness of the alloy did not change markedly. Acknowledgment This paper presents the results of one phase of research carried out at the Jet Propulsion Laboratory, California Institute of Technology, under Contract No. DA-04-495-0rd 18, sponsored by the Dept. of the Army, Ordnance Corps.

Jan 1, 1957

By D. T. Peterson, F. A. Schmidt, J. D. Verhoeven

The velocity of mig-ration of these solutes in thorium metal due to a high amperage dc current was measured at several temperatures. All three solutes migrated in the same direction as the electron flow. Diffusion coefficients were combined with the migration velocities to predict the steady-state concentration profile. Electrotransport was shown to be an effective method of removing carbon, oxygen, and nitrogen from thorium and resistance ratios several times larger than for crystal bar thorium were achieved. However, the maxirnum attained purification was well below the predicted values, apparently due to some form of contamination. THE migration of solutes in metals when a high current density dc current is passed has been known for many years. This phenomenon has been the subject of a number of reviews published in recent years.l~* Many of the earlier investigators of electromigration of interstitial solutes interpreted the direction and velocity of migration on the basis of a charge on the solute particle. This interpretation is unclear because, in a given experiment, one cannot separate the contributions to the force moving the solute which arise from the charge on the solute in an electric field gradient and the momentum transfer from the scattering of the electrons or holes. There have been only a few quantitative measurements of interstitial migration velocities and some of these have been characterized by poor accuracy or by errors in the analysis of the experiment. The electrotransport mobilities of carbon, nitrogen, and oxygen in thorium were determined by observing the rate of movement of a discontinuity in the concentration profile. This method was used by Carlson el a~.~ to determine the mobilities of interstitial solutes in yttrium and its advantages are discussed in that paper. The specimen is essentially a pair of semi-infinite, bar-type diffusion couples. The electro-transport velocity was measured by observing the displacement of the mean concentration point in a measured length of time. The diffusion coefficients were calculated by the Grube method from the concentration profile. The transport mobilities and diffusion coefficients can be used to predict the concentration profile in an electrotransport experiment of a given time, current density, and specimen length. At long times a steady state will be reached such that the natural logarithm of the concentration of each solute varies linearly along the bar with a slope which is equal to the velocity of the migrating solute divided by its diffusion coefficient. The estimation of the purification potential of electrotransport is discussed by ~erhoeven.' A comparison of the concentration profile and the purification which should be expected from known solute velocities and diffusion coefficients with that which is obtained in an experiment can be very useful in indicating whether the maximum possible purification has been achieved. EXPERIMENTAL METHOD The apparatus used in this study consisted of a sample chamber which was equipped with two electrodes, a sight glass, and a flange for mounting to a high-vacuum system capable of obtaining pressure of 2 x lo-? Torr. The chamber was made of stainless steel and was 4 in. in diam and 6 in. long. The seal between the electrodes and the chamber was made with "0" rings and compression bushings. The thorium specimens were threaded and screwed into tantalum "U"-shaped adapters attached to the ends of the electrodes. This arrangement is shown in Fig. 1

Jan 1, 1967

By Caudle B. H., Cooper H. E. Jr., O. K. Kimbler

The results of studies of areal sweepout behavior in nine-spot injection patterns are presented in this paper. The ratio of the production rates of the two types of producing wells in a nine-spot pattern is considered in its effect on breakthrough. A model study was used to develop a method to predict the production histories of both side and corner wells and to develop the production history of the pattern or the field. INTRODUCTION The importance of areal sweepout behavior in injection operations has been recognized for a number of years. We now know that the fraction of a reservoir which will be reached by an injected fluid is a function of the geometric arrangement of the production and injection wells and of the mobility ratio of the fluids. Of these two, the mobility ratio (which is a measure of the relative ease with which the oil and the injected fluid can move through the reservoir) appears to have the greatest effect on recovery. The effect of mobility ratio on production history has been reported for most of the dispersed injection patterns (i. e., five-spot, staggered and straight-line drives).4-6 ere has been very little published on the nine-spot injection pattern. This may be because the two types of producing wells in the nine-spot pattern can—and sometimes should —be produced at different rates. This adds a second variable (the fist being mobility ratio) to be covered in a study of the nine-spot flooding network. Muskat1 has reported the sweep pattern efficiency at breakthrough for a mobility ratio of one at various ratios of producing rates for the two types of wells. For these conditions he found maximum oil production at breakthrough to occur when the corner well (the producing well farthest from the injection well) is produced 10.7 times as fast as the side well. Other investigatorsa2,3 have studied a few of the many possible operating conditions possible with this injection pattern. This paper presents the results of studies which are strictly applicable to a nine-spot injection pattern surrounded by other nine-spots in a homogeneous reservoir and which has no flowable gas present (no oil bank is formed). Although the full ranges of the variables mentioned above have not been covered, the results reported here should allow reasonable estimates of field operations to be made. EXPERIMENTAL The well spacing geometry of the nine-spot pattern is shown in Fig. 1. Here we see that each injection well is surrounded by eight production wells arranged in a square. Four of these producing wells are located at the corners of the square; the other four are in the midpoints of the four sides. These two types of producing wells are called "comer wells" and "side wells" respectively. It is the ratio of the producing rates of these two types of wells that affects the breakthrough recovery, as calculated by Muskat.' This ratio is defined as the total producing rate of a corner well divided by the total producing rate of a side well. The model used in this study represented one-eighth of a nine-spot pattern, as shown in Fig. 1. It contained the equivalent of one-fourth of a side well, one-eighth of a comer well, and one-eighth of an injection well. Because of these proportions, the actual producing rate ratios in the model were one-half of the ratios of producing rates as

Jan 1, 1965

By W. O. Philbrook, K. Natesan

A mathematical model was derived for the temperature change in a particle undergoing heterogeneous chemical reaction and heal exchange with a stream of fluid. The model has been applied specifically to the exothermic oxidation of spherical particles of zinc sulfide. Very good agreement was obtained between the theoretical prediclion and experimental observation. It was found that at no stage of the reaction was the particle subjected to a state of thermal instability. In studies of gas-solid reactions, it is generally assumed that the gaseous and the solid phases are at the same temperature throughout the reaction period. This is not strictly true since most of the reactions are either exothermic or endothermic in nature. This paper deals only with heterogeneous gas-solid reactions of noncatalytic type. In the case of exothermic reactions (e.g., roasting of sulfides), there may result an increase in temperature of the particle due to the release of the heat of reaction, and this increase in temperature may enhance the reaction rate and thus cause further release of heat. Such a "spiralling" effect may lead to a thermally unstable situation within the particle. On the other hand, in the case of endothermic reactions (e.g., calcination of limestone), there may occur a decrease in temperature of the particle as the reaction proceeds, leading to a virtual "quenching" of the reaction front. A general discussion of autothermic processes with emphasis on properties and reactor design is given by Van Heerden. 1 Levesque and cubicciotti 2 observed self heating of the sample in the oxidation of thorium. Anous et al. 3 observed self cooling of the sample in their dehydration studies on chrome alum. Cannon and Denbigh, 4 and Ong et al.5 reported self heating of the sample in their studies of the kinetics of oxidation of zinc sulfide. A typical plot of heat generation and heat loss rates with reaction temperature is shown in Fig. 1.'14 The Sigmoid curve represents the heat generation rate while the straight lines represent the heat loss rate. The main objectives of this analysis are to estimate the temperature increase of a single particle undergoing an exothermic reaction and to study the possibility of a thermally unstable situation within the particle. Also, a study of the dependence of the tempera- ture increase on different process variables has been attempted. A "she ll-by-shell" method has been proposed to analyze the problem. The shells may be made as small as necessary, and equations may be set up for heat transfer by various mechanisms. Such equations will have to be solved for each shell, with the boundary conditions for ith shell being obtained from the (i - 1)th shell, using a high speed digital computer. The oxidation of spherical particles of zinc sulfide is considered for this analysis, but the treatment is general enough to be applicable to other reactions and particle shapes. The oxidation reaction can be written as AH =« -115 kcal/mole ZnS [1] Such a reaction has been found to proceed in a topo-chemical manner; i.e., as the reaction proceeds, a progressively thicker outer shell of zinc oxide is formed, while the inner core of unreacted sulfide decreases. It has been found experimentally, both in the present work and in the previous investigations,4-6 that the particle retains its original dimensions. Therefore, at any given time, the particle may be

Jan 1, 1970

By S. J. Rothman, J. J. Hines, D. Rokop, R. Bastar

Self-diffusion has been measured along each principal axis of mosaic-structured and relatively perfect a uvaniuln single crystals. The anisotropy reported before,' D[loo] = D[001] » D[010 , has been confirmed. and is shown to represent volume diffusion. Diffusion along the dislocations in the imperfect crystals increases the measured value of D significantly without changingr either the shapes of the penetvation plots or the anisotropy. SINCE a uranium is orthorhombic, the diffusion coefficients along the principal crystallographic axes need not be equal. A previous measurement of self-diffusion in three mosaic-structured a uranium single crystals1 showed that Dil00| =D[001 >>D[010 This result was deduced from penetration plots similar to those in Fig. 1. That is, the plots for diffusion along [loo] and [001] were the straight lines expected from the solution of the diffusion equation for the boundary and initial conditions of the experiment, while the plot for diffusion along [010] was curved so that the part near X = 0 had a much steeper slope than the straight lines did. The observed anisotropy could have been due either to volume diffusion or to diffusion along the very large number of dislocations in such crystals, according to the model proposed by aart.' The authors of Ref. 1 argued that the latter possibility was unlikely, but did not exclude it. To settle the question, a simultaneous measurement of self-diffusion in mosaic-structured and mosaic-free crystals was carried out. This paper reports the results, as well as a measurement of the temperature dependence of self-diffusion in a uranium. EXPERIMENTAL The single crystals of a uranium used in this experiment are listed in Table I, together with other experimental details. '(Mosaic structured" here means crystals grown by solid-state phase change.3 The Laue spots from such crystals spread over about 6 deg of arc, which indicates that the density of dislocations is very large. "Perfect" crystals are ones grown by Fisher's grain-coarsening tehnique; the Laue spots from these are quite sharp,4 which shows that the dislocation density in these crystals is much lower. All of the crystals were made from high-purity uranium; a typical analysis is given in Ref. 1. The actual diffusion directions were within a few degrees of the nominal. The thin layer-sectioning method was used. The experimental techniques were similar to those described in Ref. 1, except the sectioning. Crystals Za, 2b, 2c, 4a, 4b, and 4c were sectioned on a lathe. Crystals 3a, 3b, and 3c were sectioned by grinding on glass plates, with 1 p diamond powder in butyl alcohol as the abrasive. The use of glass plates eliminated the crowning noted in Ref. 1, and the use of butyl alcohol as a lubricant simplified the chemical processing of the sections. Material from the sides of these three crystals was removed by spark machining before sectioning. The sectioning of crystals 4a and 4c was botched. A small cylinder of metal was extruded from the center of crystal 4a while turning off material from the side of the sample. Crystal 4c was misaligned by 2.1 x 10"3 cm because the dial gage jammed during the alignment. Notwithstanding these errors, useful data were still obtained from these crystals. The points

Jan 1, 1967

By Hsun Hu

Direct observations were made on the recrystalli-zntion of a cold-rolled (110) [001] crystal of 3 pct Si-Fe by annealing thin sections of the specimen inside the electron microscope during examination. Subgrain formation, presumably by polygonization, occurs rapidly, followed by a slow increase in the subgrain size before recrystallization can be detected. No migration of the subgrain boundaries was observed, and the formation of recrystallization nuclei by coalescence of subgrains is suggested. Re-c~ystallization characteristics of various texture components were individually examined. AS reported in a previous paper: orientations of the various deformation texture components in a 70 pct cold-rolled (110) [0011 crystal of 3 pct Si-Fe were completely confirmed by transmission electron diffraction patterns taken at various areas of the specimen. These orientations can be summarized as follows: Over most areas of the deformed crystal (111) planes are parallel to the surface. These areas correspond to either $- the two main defor: mation textures, the (111) [112] or the (111) [112].2-5 Within a few areas, (100) planes parallel to the surface are frequently found. These areas represent the minor deformation texture component, the (100) [Oll] type, which is derived from rotated Neumann bands.2'4'5 Only in very rare instances are (110) planes found parallel to the surface. This orientation corresponds to the very weak texture component of the (110) [001] type. In a pole figure, this (110) [0011 orientation is located within the orientation spread between the two main (111) [1121 type textures. Usually the electron diffraction patterns taken at various areas of the specimen are very simple ones, i.e., for any one area the pattern consists only of sharply defined spots of a single orientation. In some areas, however, more or less continuous rings of weaker intensities are superimposed on a simple pattern, indicating the presence of a wide range of additional orientations. In order to understand in detail the recrystallization of this cold-worked crystal, the recrywtallization characteristics of each of the various deformation texture components must be examined. This can best be done by using a high-temperature stage and transmission electron microscopy techniques, whereby small areas corresponding to the various texture components can be selected individually for examination. The present paper presents results obtained by this method. MATERIAL AND METHOD Specimens were prepared from a previously studied1 single crystal of a commercial grade 3 pct Si-Fe alloy with an initial orientation (110) [00l], cold-rolled 70 pct from 0.014 in. to 0.004 in. in thickness. Thin foils were prepared by electrolytic polishing as described in the earlier paper.' Specimens of this thin foil were then mounted in a high temperature stage similar to the one designed by Whelan.6 A direct current was used as the heating source. The temperature of the heating grid was calibrated by means of a Pt vs Pt-10 pct Rh thermocouple in a vacuum chamber (0.01 to 0.02µHg). Maximum variation within the temperature range of 250o to 000°C was about ± 25°C. After the desired areas were selected for examination, the specimen was heated to a low temperature for some time while the selected area was constantly watched. When no further change in structure and orientation could be observed, a higher temperature was applied. Photomicrographs and diffraction patterns were taken from time to time during examination, with the heater temporarily turned off* while the photographs were

Jan 1, 1962

By J. S. Aronofsky, J. P. Heller

This paper presents a mathematical analysis of the fluid mixing which occurs during flow through porous media. The analysis is based on the well-known diffusion equation with mass transfer term. It is pointed out that the use of this equation is justified by a single general assumption which does not specify any particular mechanism for the mixing process. Formulas are given for tracer fluid concentrations for two different boundary and initial conditions. Calculated numerical values compare closely with some published results of experiments in which no viscosity or density difference existed between displaced and displacing fluids. INTRODUCTION A phenomenological theory of the mixing and diffusion process is presented for the flow in a porous medium of two miscible liquids of equal viscosity and density. This is based on the classical diffusion equation which has been used to explain such processes as Brownian motion and heat conduction. More recently, this same differential equation has been derived (in a manner analogous to the work in this paper) for the problem of heat transfer during fluid flow in porous media.' The purpose of this paper is to show that the available experimental evidence justifies the use of this simple diffusion equation. The recent data of Koch and Slobod' and of von Rosenberg' are used to test this diffusion model by comparing concentration curve shapes and calculating "effective diffusion coefficients". The same differential equation for miscible displacement, as well as some experimental confirmation, has been published in Japan by K. Yuhara,4 and recently in this country by Day.5 Yuhara arrives at his con-clusions through analogy with turbulent diffusion, de- scribing the microscopic velocities in terms of "eddy motion" — although unlike the ordinary eddies of hydraulics. Day proceeds directly from Scheidegger's6 concept of the dispersion of velocities. After a very clear description of Scheidegger's theory, he applies it to calculate the motion first of a drop and then of a thin layer of brine into a bed of sand. For the latter case, Day derives and solves the same differential equation as Eq. 1. In our presentation no detailed assumptions are made as to the mechanism of displacement. In fact, several microscopic processes may well be involved—the only general assumption made is that their cumulative microscopic effect may be described by the addition of a second-order term to the differentirl equation of continuity. DERIVATION OF EQUATIONS Let us consider the displacement in one spatial dimension of Liquid A from a sample of porous media by Liquid B. A and B are assumed completely miscible and of the same viscosity and density. Both fluids are also assumed to be incompressible and it is specified that the total flow is the same for any cross section. It is evident that the rate at which Liquid B would be carried across such a section would be given by the product of the fluid velocity, u, and the concentration, C, of B in the fluid mixture assuming the constancy of u and C over the cross section. Since the volume of Fluid B is conserved, this leads to the one-way wave equation - u ?C/?x = ? ?C/?l where x increases in the direction of u—the "Darcy velocity", Q/A (or velocity flux per unit area), and 9 is the fractional porosity. Consider the case where fluid B is displacing Fluid A and the concentration distribution is represented by a step funcLon (with a vertical discontinuity). The solution of Eq. 1 predicts no mixing

Jan 1, 1958

By M. E. V. Plummer, F. E. Beamish, J. C. Hole, J. M. Kavanagh

Prior researches have shown that the iron-copper-nickel content of platinum concentrates may be reduced by carbon to form a collecting alloy for the platinum metals in a manner exactly analogous to the formation of the lead fire assay button. The nature of the iron-copper-nickel alloys with the platinum metals has been investigated. Palladium and probably some of the other platinum metals form solid solutions with this alloy. ONE of the greatest known reserves of the platinum metals is to be found in the nickel, copper, iron sulphide deposits of Northern Ontario. The very large tonnage of copper-nickel ore processed makes the Sudbury district a significant world source of the platinum metals, the latter being recovered as byproducts in the electrolytic refining process for copper and nickel. Chalcopyrite, pentlandite, and pyrrhotite are carriers of the platinum metals;' these minerals contain from 0.03 to 2.0 ppm of the noble metals. The analysis of the above ores for the platinum metals necessitates an initial concentration of the minute quantities present in the ore. The method usually employed involves the extraction of the noble metals by molten lead with or without small proportions of silver. The lead is then separated from the precious metals by cupellation, a process which involves absorption by the cupel and volatilization of lead oxide, leaving the precious metals collected in a silver bead on the surface of the cupel. The collection by lead is probably achieved by a process of dissolution followed by some degree of segregation in the lead alloy since the cold button is seldom uniform in composition. With silver-lead alloys this process of liquation has been used to re- cover pure lead since, during the cooling period, crystals of pure lead separate and silver becomes concentrated toward the center of the solid.' Platinum and gold liquate in a manner similar to silver which may account for the successful recovery of platinum by fire assay even though it is considered insoluble in a solid lead medium. 3 In the case of iridium the liquation process may be a disadvantage in that the iridium may appear in the outer areas of the button and thus be lost mechanically during the handling processes. In any case the deficient alloying characteristics of lead for the platinum metals encouraged attempts to find an improved collecting alloy, and by analogy with aqueous extractions, preferably one in which solid solutions would be readily formed. Since, beyond a certain minimum solute to solvent ratio, the formation of a solid solution is accompanied by a lattice distortion, one would expect some relationship between the atomic diameters of any two metals and their tendency to form a solid solution. In one survey5 the hypothesis was put forward that when the atomic diameter of the solute differs by more than about 15 pct from the solvent, the "size factor" is unfavorable and the formation of a solid solution is restricted. On the other hand, when the atomic diameters are within this limit the size factor is favorable and solid solutions are readily formed, e.g. the transition metals of Group VIII having similar radii form solid solutions over wide ranges of composition.6 Even with a favorable size factor, elements having different crystal structures cannot form a continuous series of solid solutions. Cell dimensions alter as the concentration of the solute increases, causing a

Jan 1, 1962

By Leon Merker

FLAME fusion growth of strontium titanate crystals was undertaken to obtain large transparent crystals on which physical data could be gathered. The fact that strontium titanate is a cubic crystal and probably has a high index of refraction indicated that it might be useful in optics. No large boules* of strontium titanate had been synthesized previously, and none have been found as natural occurrences. The flame fusion process, invented by Verneuil in 1904,1 was first applied to grow large crystals of ruby and has since been used extensively to grow crystals of corundum and spinel on a commercial scale. Only during the past nine years has this technique been applied to the growth of large single crystals of other high melting compounds such as rutile,2 mullite,3 and ferrites4 One of several significant advantages of the flame fusion technique over more conventional methods of growing crystals from a melt is that the crystals are grown without the use of melt containers. This fact alone excludes other methods of growing crystals from a melt whenever there are compounds with a melting point near 2000°C or above. Growth of the crystal, similar to that of a stalagmite, takes place inside a circular refractory furnace that does not contact the crystal, Fig. 1. This eliminates a serious source of impurities as well as the likelihood of multi-crystalline solidification along the container walls. The flame fusion technique is particularly advantageous for research because it produces large crystals quickly and at lower cost. A corundum crystal 25 mm long and 12 mm in diam can be grown in 2 to 3 hr. Operating conditions permit observation of the growing crystal during any stage of the process. The various flame patterns and a wide range of oxidation-reduction conditions can be investigated in a relatively short time. It should be pointed out that formation of the foot of the boule is the most critical step in growing an unfractured boule. When the tip of the sintered cone is fused to initiate crystal growth, the original area of fusion should be kept as small as possible so that one crystallite can assume dominance over all the others and propagate as a single crystal. If the boule start happens to consist of several crystals, usually the entire boule also contains several crystals. In many cases these tightly cohering crystals exhibit different crystallographic orientations. Although the basic technique of the original Verneuil process has been retained, refinements in method and burner design have been developed in recent years. For example, boule growth may be initiated on a seed crystal rather than by fusion of a sintered cone. The danger of multicrystalline or twinned boule starts is thus eliminated, and this procedure may also be used to grow boules of predetermined crystallographic orientation. Prior Art: Of the many attempts to produce artificial gem crystals, most early efforts% ere directed toward synthesis of .rubies and sapphires and diamonds. None of these efforts resulted in a process that was commercially practical. Without exception only tiny crystals of corundum were obtained, and those at excessive cost. Verneuil announced his process in 1904; Ravier6 in 1934 presented a detailed description of a commercial boule production plant. Descriptions of the German synthetic stone industry, the process details, and the equipment used have been published by Barnes7 and by Merker.'

Jan 1, 1956

By David J. Mack, T. V. Philip

The effects on the Cu-A1 eutectoid transformation of the addition of a third element from the second long period and from Group IIB of the periodic table, were studied by isothermal transfomation techniques. The valency of the third element shows the most pronounced effect on the eutectoid decomposition. With some exceptions, the atomic size factor of the third element shows some consistent effects on pearlite nucleation and growth. The atomic concentration or the electronegative valency of the third element shows wo consistent effect on the eutectoid decomposition. Addition of the third element helps spheroidization of pearlite. SINCE the pioneering work of Davenport and Bainl on the eutectoid decomposition in steel by the isothermal transformation technique, an enormous amount of data has been accumulated on the eutectoid transformation characteristics of commercial steels. While the eutectoid reaction in steel is of tremendous industrial importance, it appears to be quite unique among the eutectoid reactions. Though the metal-lographic structures developed in the Fe-C and Cu-A1 eutectoids are similar, it is now evident that the two eutectoid reactions are quite dissimilar. The eutectoid decomposition in the Cu-A1 system is more complicated partly due to the appearance of transitional phases in addition to the equilibrium phases. * The first isothermal studies of aluminum bronzes were made by Smith and Lindlief2 and later by Mack,3 Klier, and Grymko,4 and recently Hayne 5,6 A few other systems, structurally analogous to the above, have been studied by the same technique.7-10 Kasberg and Mack1' studied the effects of iron on the transformation of commercial aluminum bronze. The effects of manganese and nickel on this aluminum bronze and those of iron, manganese, and nickel on the binary Cu-A1 eutectoid have been recently studied.12 The object of the present investigation was to study, by isothermal transformation techniques, the effects (with respect to the fundamental factors of alloy formation) of the addition of a third element on the transformation of a Cu-Al eutectoid alloy, 11.8 pct Al. The third elements chosen for the purpose were silver, cadmium, indium, tin, and antimony from the second long period, and beryllium, zinc, and cadmium from group 11B of the periodic table. They were added in nominal amounts of 0.5, 1.0, and 2.5 pct by weight to a eutectoid Cu-A1 alloy. These third elements were expected to show the effects of electron valency, atomic size factor, and electronegative valency effects on the binary eutectoid transformation. PFLEPARATION OF ALLOYS AND PROCEDURE All alloys were made as 400-g heats by induction melting in a graphite crucible with a graphite lid. The molten alloy with a layer of graphite powder on top was stirred at regular intervals by a graphite rod. When the melt was judged to have acquired sufficient superheat, a pinch of common salt was sprinkled on top and any crust scraped off to one side. It was then chill cast into a 5/8 in. rod. The alloys with silver were prepared by remelt-ing the high-purity Cu-A1 eutectoid alloy used by Mack3 with high-purity silver (99.99 pct pure from Hardy and Harmon, Inc.). In the preparation of all the other alloys, polished electrolytic copper shot (from Belmont Smelting and Refining Co.) and high-purity aluminum (99.99 pct pure from Aluminum Co. of America) were used.

Jan 1, 1962

By H. P. Stuwe

It is investigated whether a model of sphere packing permits the computation of the lattice parameters and their variation with concentration in various a phases; it does, except in phases containing silicon. Several observations indicate that a particular role is played by the "E-atoms" in the lattice. THE phase is formed by many alloys of the transition elements.'-' Its structure has been described by Bergman and Shoemaker,5,8 Dickins, Douglas, and Taylor,7 and Kasper, Decker, and elanger.' There are five kinds of crystallographically equivalent positions for the 30 atoms in the unit cell which have been designated by these authors I, 11, 111, IV, and V or, in the same sequence, A, B, C,D, and E, as shown in Fig, 1. For the purposes of the present paper two elements of the lattice will be discussed separately: The lattice planes perpendicular to the c-axis which are made up by the atoms A, B, C, and D and which determine the a-parameter; and the close-packed "vertical rows" of E-atoms which determine the c-parameter. asper,9 and Frank and kasper10 described the a phase and other complex structures as sphere packings. In these papers particular attention is paid to the coordination polyhedra. Lattice positions with an equal number of nearest neighbors are considered equivalent regardless of their crystallo-graphic description. This seems to work very well for the explanation of ordering phenomena in the a phase as far as information on this is available." If the description of the phase as a sphere packing is valid it should be possible to draw two further conclusions: 1) It should be possible to calculate the lattice parameters from the atomic radii of the components, and 2) It should be possible to predict the variation of the lattice parameters with concentration within one phase field. Both points are discussed in this paper. PROCEDURE Measurements of lattice constants were done with a symmetrical back-reflection X-ray camera. With Cr-radiation, about a dozen a1 lines could be observed with most samples. From the (550/710)- and 720-lines the a-value was computed. With this value the c-parameter was calculated from the (333)- and (413)-lines. Finally, the position of all the other lines was computed from the lattice parameters thus obtained. In this way the validity of the indexing could be checked and, at the same time, this proce- dure gave an estimate of the accuracy of the lattice parameters. A correction was made for film shrinkage. The atomic radii used in the computations are the Goldschmidt radii for coordination number 12. Actually, some of the atoms have a coordination number of 14 or 15, but the error thereby introduced is small. * For the computation of lattice constants from the atomic radii an "average atomic radius" of all components was used, i.e., complete disorder was assumed. Although this seems to be true for some systems like OS-Cr,4 it is not true in others.11,13,14 The ordering model, however, which has been suggested by Kasper on the basis of extensive research lets both kinds of atoms contribute to both the ''vertical rows" and the lattice planes perpendicular to them, so that even with ordering of this kind the assumption of complete disorder cannot introduce a very serious error in calculating lattice parameters. RESULTS Bergman and shoemaker6 gave a table of the interatomic distances in the Fe-Cr phase. If we disre-

Jan 1, 1960

By B. H. Kear

All the possible ways of forming simple two-dislocation boundaries in multiple slip orientations are deduced. It is shown that the principal differences in work hardening behavior of crystals plastically deformed in these orientations can he accounted for in terms of the likelihood of formation of diffuse boundaries, or deformation bands, consisting of the products of dislocation reactions of the Lomer type. Also, it is pointed out that the break-up of boundaries under the applied stress is more likely to yield interstitial jogs than vacancy jogs, whereas glide dislocations cutting throuqh dislocations trapped in particularly stable boundam'es can acquire both types of jogs. USING Frank's' grain boundary formula, Ball and Hirsch2 determined the parameters of all the possible boundaries that can be built from two sets of glide dislocations in the fee structure. This paper deduces the geometrical possibilities for the formation of these two-dislocation boundaries in the different multiple slip orientations, and discusses the role of boundaries in work hardening and in jog formation. Formation of Two-Dislocation Boundaries in Mutiple Slip Orientations. The parameters of the possible two-dislocation boundaries for any crystal orientation can be deduced most conveniently using the boundary formulae due to Amelinckx,3 in conjunction with the stereographic projection shown in Fig. 1. Suppose that slip occurs in a-BC and ß-CA. According to Amelinckx, when the two systems of glide dislocations are parallel, the rotation axis, u, of the boundary is parallel to (a x ß) and the unit normal, u, defining the boundary plane is parallel to (a x ß) x (BC x CA). On the other hand, when the glide dislocations form a crossed grid, u is parallel to (BC x CA) and u is parallel to (a x CA) x (ß x BC). In Fig. 1, (a x ß) is the pole of the great circle passing through a and ß, so that u is parallel to DC, i.e., parallel to [110]. Furthermore, (a x ß) x (BC x CA) is the pole of the great circle passing through DC and 6, so that u is parallel to [110]. Hence, when all the glide dislocations are parallel, the boundary is pure tilt with parameters u = [110], v = [ 110]. Similarly, it may be shown that the corresponding crossed grid of dislocations is a mixed boundary with u = [111], u = [111]. Consider the case of a crystal compressed in the [001] orientation. In this orientation, according to Fig. 1, there are eight available slip systems. In what follows, the different types of two-dislocation boundaries that are possible in this orientation will be examined by taking different combinations of two slip systems out of the available eight, with a-BC always one of the systems. Case (i)—Slip Systems a-BC and d-CB. The formation of a tilt boundary in (101), which consists of a crossed grid of edge dislocations with vectors BC andCB, is illustrated in Fig. 2(a). As the figure shows, slip is required in the system a-BC in one part of the crystal and in d-CB in the other part, with both sets of dislocations entering the boundary from opposite sides. The rotation axis of the boundary lies in (101), depending on the relative densities of the two sets of edge dislocations in the boundary. Case (ii)-Slip Systems a-BC and ß-CA. The two possible boundaries are shown in Figs. 2(b) and (c); boundary (b) is tilt with u = [110], u = [110], and

Jan 1, 1962

By Jean-Jacques Prompsy, J. S. Rinehart

The purpose of the present study was to arrive at a better understanding of the stress fields developed within drill bits under dynamic loading and the influence that these stress fields could have on the failure of the drill bits. Particular emphasis has been given to the influence that bit shape has on the establishment of highly localized, potentially destructive stress inhomogeneities within the bit. This paper covers one phase of a broad study involving three velocity regimes: impacts at very low velocities, 0 to 20 ft/sec, a velocity range in which dynamic effects are just beginning to be found; impacts at velocities ranging from 20 to several hundred ft/sec, arange commonly encountered in practical drilling operations; and impulsive loading through the detonation of explosives, a velocity range in which the dynamic effects are greatly exaggerated and made more identifiable. Specifically, fracturing which develops at the end of a cylindrical bar undergoing a shock produced by impact was experimentally investigated, the impacted area being the face of a truncated cone terminating the cylindrical bar. Fracture characteristics were looked at as a function of impact velocity, area of impacted face and cone angle — parameters of considerable importance in the design of drill bit. EXPERIMENTS The specimens were made of Plexiglas, which enabled the fracture pattern to be easily observed. In Plexiglas, the velocity of a longitudinal wave is about 8,700 ft/sec, and the velocity of the transverse wave about 4,200 ft/sec. The specimens were impacted by a steel cylinder whose lower face was parallel to the terminal face of the experimental bit, as shown in Fig. 1. The impacting mass was shifted up and down on two steel rods, dropping from a maximum height of 11 ft, with the velocity varied between 0 and 6.5 m/sec. The initial mass of the cylinder was 4.1 kg, which could be increased by steps up to 8.8 kg by addition of steel plates. The velocity of the steel cylinder was determined by using an electric circuit; as the cylinder dropped, it contacted two steel wires, the first one starting and the second one stopping an electronic counter. Two similar counters connected in parallel were used to corroborate each other. The velocity was checked at each drop as it had been observed to vary somewhat from drop to drop. The two contact wires were located just above the top of the specimen so as to give a value of the instantaneous velocity of the steel cylinder at impact. The distance between the wires was 40 ± 0.5 mm and the discrepancy between the counters was never more than 0.1 m/sec. Thus the velocity was known to within 3.25 per cent. This precision was sufficient with respect to the precision of the machining of the specimens. A very small lack of parallelism between the face of a specimen and that of the impacting hammer resulted in an oblique impact which required a much lower velocity to break the specimen. The experiments had a twofold purpose — (1) to observe the fracture pattern, and (2) to find relationships between the minimum velocity required to generate a fracture within a specimen and the

By J. H. Henderson, R. J. Wygal, J. Naar

Experimental work is reported which shows that consolidated rocks and unconsolidated porous media exhibit different imbibition flow behavior. At a given saturation the imbibition nonwetting permeabilities for a rock are smaller than the drainage permeabilities. The contrary happens for unconsolidated aggregates — imbibition nonwetting permeabilities are larger than drainage ones. A similar difference is observed for the wetting phase. Imbibition permeabilities are larger than drainage ones for a consolidated rock but smaller than drainage permeabilities for an unconsolidated medium. The results of these differences are examined for two cases. 1. Flooding Efficiency — Craig's1 scheme for the computation of production history of a five-spot water flood is shown to agree extremely well with experimental results obtained when using a system packed with glass spheres if imbibition relative permeability curves are used. 2. Alcohol-Slug Displacement — Published theory 2 on oil displacement by alcohol slugs has been questioned despite the apparent agreement between predicted and observed results. The present work suggests that, if imbibition relative permeability curves characteristic of the unconsolidated media used in the early experiments had been available to make the predictions, the inadequacy of the theory would have been immediately evident. The experimental work shows that poorly consolidated formations tend to behave like unconsolidated media. Finally, it is shown that the difference in imbibition behavior is directly related to pore-size distribution and cementation. PART 1 — THE FLOW BEHAVIOR OF UNCONSOLIDATED AGGREGATES Experiments on scaled models of field reservoirs are useful for studying new displacement processes which are incompletely understood. Even when a mathematical description is possible, the solution might be difficult and complex. An answer obtained from a scaled model is extremely valuable in such cases. A great amount of work$-5 therefore, has been devoted to the derivation of scaling laws. Similarity groups have been defined which assume (1) that the relative Permeability curves of the prototype and the model are the same whether the displacement is an imbibition or a drainage process and (2) that there is a linear relationship between the capillary pressure of the model and the prototype. For practical reasons (simplicity in the preparation of models, duration of the experiments, etc.), the porous media of laboratory models are usually unconsolidated packs of sand or glass particles. Hence, unless the capillary and flow characteristics of unconsolidated and consolidated systems are identical, the model data are applicable only to unconsolidated formations. The usefulness of scaled-model studies may then be seriouslyrestricted since most oil-bearing sands are consolidated. Perkins and Collins6 suggested the use of model and prototype curves normalized with respect to both relative permeability and saturation to improve compliance with scaling criteria. Even this technique does not give a satisfactory model-prototype match. This paper reports an observation of two-phase flow in unconsolidated sands which shows that, for most displacements, "scaling" in the strict sense of the word is not even qualitatively feasible with a sand model. It provides, however, a firm foundation for testing a theory by matching it with observed performance of laboratory-size models. EXPERIMENTAL AS a part of a basic study of packed aggregates, the relative permeability of glass-spheres and sand-grain packs was measured with capillary control. The fluids were oil and air' The spheres were packed in plastic cylinders 1-in. in diameter;

By B. F. Buie

For over 200 years meerschaum has been a significant item of trade between the Near East and countries to the west. Best-known for its use in making smoking pipes and cigar and cigarette holders, it is also used in making a variety of decorative and ornamental items. It is easily carved, and is given an attractive, lustrous finish by rubbing. No published specifications are known, but compact form, purity, uniform whiteness or light color, and uniform, porous texture appear to be requisites. Fig. I a shows the appearance of a piece of crude meerschaum from Turkey; Fig. lb a carved item; and Fig. lc, d, and e show the fibrous and porous texture as revealed by the scanning electron microscope. Neither a major industrial mineral nor a precious stone, meerschaum nevertheless holds a record of long-and continuing-association with art and personal affairs of man. It is also a mineral to which modem methods of exploration and production probably could be profitably applied. An interesting account of the preparation and use of meerschaum in the early part of this century is given by Sterrett (1907). Geology and Mining Mineralogy Meerschaum is recognized now to be the compact variety of sepiolite, the claylike variety of which is one of the minerals included in the chapter on clays. The word meerschaum, from the German meaning "sea foam," was in earlier times used synonymously with sepiolite, a word of Greek origin suggestive of the resemblance of the compact variety-then the only variety known-to cuttlefish bone. At one time meerschaum was used as the commercial name, and sepiolite for the scientific name. However, to think of meerschaum as the name by which sepiolite is known commercially is no longer valid, in view of the recent emergence of the earthy variety as a competitor of bentonite and other clay materials. Present tendency is to restrict the term meerschaum to the compact variety, and to use the term sepiolite as a more general name to include both the compact and earthy varieties. This is the terminology followed in this publication. Meerschaum typically is white to light gray, has a hardness of 2 to 2 ½ , conchoidal to irregular fracture, and specific gravity varying with porosity from about 2 to less than 1. The composition is somewhat variable,- but is near that determined by Schaller (1936) for sepiolite as 2MgO.3SiO2.4H2O. A more modern for mula based on X-ray determination of the crystal structure is given by Caillére and Henin (1961) as (Site) (Mg9)030(OH)6(0H2)46H20. The reason for the firm, compact nature of meerschaum, in contrast with the earthy variety of sepiolite, is not fully understood. It may be due to the presence of a very minor amount of some bonding material such as silica. Scanning electron microscope (SEM) photographs of a specimen from Turkey (Figs. lc-le) indicate that the meerschaum is composed predominantly, if not entirely, of crystalline fibers which merge into sheetlike forms in a complex, intertwined mass. Detailed X-ray and crystal structure data, as well as some data on thermal stability range, are given by Caillère and Hénin (1961) for the sepiolite family of minerals. Mode of Occurrence and Origin Most authors have attributed the origin of meerschaum to alteration of serpentine or other magnesian minerals. The occurrence in alluvium and other surficial deposits does not necessarily indicate that the alteration was supergene. In fact, hydrothermal origin appears more likely for some deposits, though not for all.

Jan 1, 1983

By A. T. Aldred, K. M. Myles

The vapor pressure of chromium over solid V-Cr alloys has been measured by the torsion-effusion method in the temperature range 1450" to 1650°K. The chemical activities as well as the free energies, entropies, and enthalpies of formation of the alloys have hem computed from the vapor-pressure data. The activities of both chromium and vanadium exhibit fairly small negative deviations from Raoult's law over the entire composition range. The values of the excess entropies and the enthalpies of formation found at 1550°K are considered in terms of the changes in the clzaracteristic properties of vanadium and chromium that occur upon alloying. RECENT interest in the theory and properties of transition-metal alloys has shown the need for relevant thermodynamic data. The present investigation of the V-Cr system was undertaken as part of a more general study of thermodynamic properties of transition-metal alloys. The system was chosen because of the reported complete solid solubility of the components at all temperatures1 and because the difference in magnitude of the vapor pressures of vanadium and chromium allowed a simple effusion technique to be used. The torsion-effusion method was employed to minimize compositional changes in the alloys during the experiments. EXPERIMENTAL Equipment. The torsion-effusion apparatus, shown diagrammatic ally in Fig. 1, consists of a torsion system suspended inside a vertical vacuum chamber whose lower end is surrounded by a high-temperature vacuum furnace. An auxiliary optical lever and screen are used to determine the angle of rotation of the effusion cell. The torsion fiber is an annealed, high-purity, tungsten wire usually between 0.003 and 0.007 cm in diameter and about 43 cm long. The wire supports a concave mirror and a rigid tantalum connector tube. The effusion cell is fastened to the lower end of the connector tube by a dovetailed yoke and har- ness arrangement. The cells, which are 1.27 cm in diameter and 4.4 cm long, were made by electron-beam welding tantalum endcaps onto a 0.025-cm-wall tantalum tube. The two orifices, formed by drilling, have similar areas ranging from 0.011 to 0.031 sq cm. The noninductive sheath-type furnace consists of a single-phase tantalum heating element 6.35 cm in diameter, 15 cm high, and 0.013 cm thick; the element is completely surrounded by radiation shields in an evacuated chamber. The power system includes a 17-kva self-saturating saturable core reactor with an integral magnetic amplifier and a water-cooled, 15-kva, step-down power transformer with a multitap secondary. Temperatures are controlled to within ±1/2o at 1600°K by means of a unit that includes a reference voltage source, a dc electronic null detector, a current-adjusting proportional band relay, and the self - saturating reactor. At 1650°K the temperature variation within the cell, vertically and horizontally, is 3/4oK. Unwanted motions of the torsion system are damped by the movement of an aluminum vane, which is mounted on the connector tube, between the poles of an electromagnet. When not in use the magnet is withdrawn and rotated 90 deg to prevent its interaction with the rotation of the system. Measurements. The vapor pressure, p, is re-

Jan 1, 1964

By Mark J. Klein, A. H. Clauer

The Snoek peak induced by solute nitrogen in chromium was studied. A rapid quenching rate is required to maintain nitvogen in solution in sufficient concentrations to be detectable by internal-friction techniques. An expression for the tertninal solubility, c, of nitrogen in weight pevcent is log,,c = -4130/T + 1.05. The diffusion of nitrogen in zhromium determined by internal friction is SNOEK&apos;S theory1 of the stress-induced ordering of interstitial solute elements in the bcc lattice has been used to deduce information on the behavior of interstitials in a number of bcc metals. Under the proper conditions, internal-friction studies of "Snoek peaks" permit a determination of diffu-sivities, solubilities, and precipitation kinetics of interstitial elements. Internal-friction studies of this type have been reviewed by Nowick,&apos; Wert,= Powers and Doyle,4 and others. The behavior of interstitial atoms in Group V-A metals and iron has been rather extensively studied in this way, and results are reasonably self-consistent. However, the study of interstitial elements using internal friction in the Group VI-A metals—tungsten, molybdenum, and chromium—is much less clear. Thus far, internal-friction studies have revealed a peak in tungsten tentatively attributed to carbon5 and a peak in chromium associated with nitrogen. In chromium, Bungardt and Preisendanz6 observed a peak at 180°C for a frequency of -1 cps which they associated with nitrogen. More recently, de Morton7 determined the diffusivity of nitrogen in chromium from studies of a nitrogen peak at 160°c for a frequency of -1 cps and from elastic aftereffect measurements. The objective of this investigation was to study the nitrogen peak in greater detail, to obtain information on the behavior of nitrogen in chromium. This information is of interest because nitrogen is known to have an embrittling effect on chromium. However, efforts to deduce the role played by nitrogen in influencing the mechanical behavior of chromium are handicapped by lack of information on its behavior and distribution in chromium. EXPERIMENTAL PROCEDURES The chromium wires used in these studies were drawn from arc-melted iodide-chromium stock that had been extruded, swaged, and centerless-ground to 0.25 in.-diameter rods. Internal-friction measurements were made in an evacuated torsional pendulum similar to the one described by Kê,8 utilizing wires 0.048 to 0.030 in. in diameter. A series of analyses of the fabricated wires showed that the major impurities were approximately 100 ppm Fe and a combined interstitial content of 50 to 100 ppm C, O, N, and H. A typical interstitial analysis of the wires after processing is 15 ppm C, 2 ppm H, 50 ppm O, and < 5 ppm N. Interstitial concentrations were determined by vacuum-fusion, combustion-conductimetric, and micro-Kjeldahl analyses. In general, vacuum-fusion analysis yielded a lower nitrogen concentration than did micro-Kjeldahl analysis. For consistency, the nitrogen concentrations reported in this investigation were determined by the former method. Heat treatments were carried out by encapsulating the specimens in quartz tubes containing argon. The fabricated wires were annealed at 1150°C for 1 hr to produce uniform recrystallized grain size of about 0.015 mm. Nitrogen and carbon were added to the fabricated wires by encapsulating them in a quartz tube containing ammonia or methane, and annealing them at 1150°C for 48 hr. In some instances, carbon was also introduced by coating the wires with carbon black prior to annealing.

Jan 1, 1965