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By Paul B. Crawford, Melvin B. Burton

A procedure is given which describes how the gelatin model may be-used to study the effect of mobility ratio on areas swept before and after breakthrough in fluid injection programs. Differences in mobility ratio are achieved through the proper selection of electrolytes for the invading and displaced fluids. The progress of the front is followed visually and photographed or traced onto plates (it desirrd positions before and after breakthrough. Example applications to standard flooding pattrrrls, vertically and horizontally fractured reservoirs and the in-situ combustion recovery process ore presentrtl. INTRODUCTION In water flooding, gas cycling and other fluid injection programs, the sweep efficiencies and swept areas after breakthrough are known to depend on the mobility ratio. Previous investigators have described the application of the X-ray technique, stepwise use of the poten-tiometric model and numerical procedures for studying these problems.1,3,4,6 Others have described the application of the electrolytic and gelatin models for studying sweep efficiencies at a mobility ratio near one.2,5,7,8,9 It is a purpose of this paper to describe a procedure for using the gelatin model to study swept areas at mobility ratios other than one. PROCEDURE The technique developed for gelatin model studies is composed of four phases: the construction of the model, the selection and preparation of electrolytes, the running procedure, and the recording of the swept areas. Each phase is discussed independently in order that the over-all procedure will not be confused. Construction of the Model Since some of the advantages of the gelatin model are its economy, and speed, it is desirable to have as small a model as is possihle without creating error through the inability to follow visually the progress of the colored front. For this work, it was found that a distance of 4 to 6 cm between salt bridges is convenient to use. Since the size is relatively small, it follows that the gelatin layer must be thin. For this study, 1/8 in. was used as the gelatin thickness in each case. The model was constructed on a flat 1/2-in. plastic (Lucite) base with boundaries constructed from 1/8-in. plastic strips. The well pattern to be studied was outlined with the 1/8-in. plastic strips on the base with narrow apertures (- - 2 mm) in the boundary to represent the wells. Behind each well opening, a rectangular area of some 8 to 12 sq cm was outlined by additional plastic strips. The plastic strips were fastened to the base by any suitable non-conducting cement so that the boundary system was liquid tight. The electrodes were constructed from sheet copper in a fashion that would allow easy removal. In this work, it was found that by simply forming a large U-shaped clip from a strip of the copper that would slip over the boundary was satisfactory and easily removed. The electrodes should be as wide as one of the sides of the rectangular area behind the well opening. If a fractured pattern is studied, fractures inside the well pattern may be represented by strips of copper of the desired shape mounted in some easily removable fashion. The reason that apertures (or salt bridges) were used for wells instead of metallic electrodes is that the plating of the electroactive ion at the cathode and the introduction of an electroactive ion at the anode has a detrimental effect on the swept area outline if these operations \rere permitted to be carried on within the well pattern itself. The anode may be used within the well pattern

Jan 1, 1957

By P. R. PaeIay, A. Slibar

Using three-dimensional, stress-deformation rate equations for a Bingham plastic, an approximate solution for the laminar flow of drilling mud between the drill pipe and casing is given for the case when the velocity gradients due to axial flow are large compared to velocity gradients due to rotation. INTRODUCTION The fact that drilling mud can be treated, to a first approximation, as a Bingham plastic has been established by several investigators12,3,4,5. A Bingham plastic is a material that is rigid until a certain combination of stresses exceeds a critical value. When the stresses are sufficiently high to initiate flow, the reduced stresses have a constant ratio to the corresponding deformation rates at each point. Prediction of pressure and torque requirements to produce the desired flow of the material have been made by analytical solutions due to Bingham6, Reiner7, and Buckingham8. One part of the present means of calculation that requires clarification is the flow between the drill pipe and casing. Present methods treat the flow resulting from the rotation of the drill pipe separate from the flow produced by the axial pressure gradient. Since the non-linear behavior of a Bingham plastic implies that solutions cannot be superposed, the present analytical investigation was undertaken to obtain a solution to the problem of combined axial and tangential flow. Equations have been proposed for the three-dimensional flow of a Bingham plastic9,10,11 and it has been shown that the proposed equations lead to the usual solutions from the one-dimensional theory". Applying these equations to the flow between two concentric cylinders the problem of axial flow only13 and tangential flow only12 have been solved. In addition to the separate problems of axial and tangential flow between concentric cylinders, the authors have determined the relationship between external torque on the inside cylinder and axial pressure gradient necessary to initiate flow". The solution given in Ref. 14 points out the complexity of the present problem. The problem considered here is the case of simultaneous axial and tangential flow. The axial flow, necessary to keep the drill bit free of removed material, results from the axial pressure gradient. Tangential flow is present since the drill pipe rotates in the hole. The solution given here is valid when the shear stresses in the drilling mud due to rotation are small compared to the shear stresses resulting from the axial flow. This case for laminar flow can be encountered in oilfield drilling. STRESS-DEFORMATION RATE EQUATIONS A detailed derivation of the equations governing the behavior of a Bingham plastic is given in Ref. 10. This section briefly outlines the physical reasoning leading to this three-dimensional formulation of the stress-deforma-tion rate law. Experiments on a certain class of materials have demonstrated the following properties: (1) when a certain combination of the stresses is below a critical value no deformation of the material occurs, (2) to a first approximation this combination of stresses is independent of the hydrostatic pressure (negative average normal stress), (3) if flow occurs the ratio of the reduced stress to the corresponding deformation rate is constant at each point of the flowing material, and (4) incom-pressibility. For the one dimensional case of shear deformation, Bingham6 formulated the governing equation as dxx, = 0 for sxx < to and µdxx = s xx = to for sxx to , . where dxx is the shear deformation rate, sxx the shear stress, to the critical shear stress, and p the slope of the sxx vs dxx curve for dxx > 0. Fig. 1 shows the behavior for pure shear between two parallel plates. Taking into account the physical properties given above and assuming the critical combination of stresses to be measured by the octahedral shear stress, the fol-lowing three-dimensional equations have been proposed.

Jan 1, 1958

By M. Cross

INTRODUCTION The development of refining processes involving gas injection into liquid metals has seen the evolution of a variety of designs [I]. During the last few years or so the top, bottom and side blown processes have, essentially, been combined to produce a series of mixed gas blowing schemes (eg. the LBE [2,3] and LD-OB processes [4,5]). Although, a variety of mixed gas blowing processes have been carefully engineered and recorded operational experience now attests to their effectiveness, optimising their design and operational performance is another matter. Detailed process analysis of such operations is impeded by the complex nature of the physical and chemical processes involved: . fluid flow behaviour of the liquid metal and injected gases, . chemical reactions in the vessel, . mass and heat transfer within the fluid phases in the vessel, . influence of the geometry (eg. location of the tuyeres, etc.), . monitoring the operational performance. By and large, design and operating parameters evolve from a combination of laboratory scale experiments, pilot plant work and, of course, practical engineering experience. Conventionally, the gas-liquid system used to model the industrial process is air-water, for a number of valid reasons including – . visual demonstration plus geometric and (to some extent) dynamic similarity, . the fact that the system demonstrates, at least, qualitatively many of the complex features of gas injection - interaction between the gas and the liquid - the nature of the gas flow behaviour - qualitative indications of mass transfer rates Important experimental contributions to the understanding of gas liquid systems include the work of Brimacombe [6], Wraith [7], T.Robertson [8], D.Robertson [9], Guthrie [10] and Szekely [11] (these references provide starting points to identify the major published work in this area). Although air-water systems have been instrumental in providing a great deal of qualitative insight into the fluid flow behaviour of the processes, reliable quantitative predictions are more problematic. A number of workers (notably Brimacombe and colleagues [6, 12])have demonstrated that different gas-liquid systems can lead to wide variations in behaviour. Other limitations of physical modelling include the assumptions of isothermal flow, scale- up of turbulent flow behaviour and the inability to directly include the effects of chemical reactions. The above limitations in physical modelling, therefore, yield a natural role for mathematical modelling provided that it can be made comprehensive enough whilst remaining cost-effective. The major physical processes involved in gas injection systems may be summarised as . behaviour of the gas envelope structure in moving through the liquid metal . the global movement of the liquid metal . momentum interaction between the liquid metal and the injected gas . heat transfer and mass exchange via the chemical react ions. The behaviour of gas envelopes forming at a tuyere and then moving through the liquid has been investigated thoroughly by a large number of workers. Such models are generally based upon simple physic- al concepts derived from visual observation. An early model by Davidson and Schuler (DS), based upon a force balance between the liquid inertia and a combination of the bouyancy and momentum forces on the envelope, has proved to be a useful tool in predicting bubble volumes at detachment [13]. The DS model has been successfully modified in a number of ways [14-17] to cope with heat transfer and reactive transfer rates, etc. Furthermore, the basic ideas summarised above have also been confirmed on industrial converters by both Guthrie [I8] and Brimacombe [19]. Models have also been developed to predict the behaviour of the "chemical" component of some gas-liquid systems. However, one thing is clear; all gas-liquid reactive systems are governed by the degree to which effective mixing occurs. As such, the processes are dominated by the behaviour of the heat and mass exchange of the two phase fluid flow system. It is the pattern of movement of the turbulent liquid metal plus its interaction with the injected gas which must be predicted reliably, if sufficiently comprehensive comp-

Jan 1, 1984

By J. E. Hilliard

AS is well-known, the length per unit volume, Lv, of any lineal feature (such as grain edges in a polycrys-talline specimen) can be estimated from a count of the number of point intersections with a random plane of polish. Specifically:&apos; Lv = 2NA in which NA is the expected number of intersections per unit area of the plane. As with all other measurements in quantitative metallography, the estimation of Lv is subject to a sampling error, the magnitude of which will depend on the number of observations. The analysis is usually performed by counting the number of intersections falling within a closed test figure of known area applied to a succession of different areas on the plane of polish. The test figure can be drawn directly on the screen of a projection microscope or else superimposed on the image of the structure by means of a reticule in a focusing eyepiece of an ordinary microscope. The use of micrographs is very inefficient unless, as in the present study, these are also needed for some other purpose. Consider an analysis in which n applications are made of a test figure within which the expected number of intersections is N. The expected total number of intersections counted is then given by NT = nN If (N) is the standard deviation in N for one application of the figure and if the latter is applied to non-overlapping regions so that the N&apos;s are independent in the statistical sense, then the standard deviation in Nt is given by: a(NT) = a(N)vn Eliminating n from Eqs. [2] and [3] we obtain: Thus, k will be a constant for a given specimen and test figure. From Eqs. [ I] and [4] we find As is to be expected, therefore, the fractional error in the estimation of Lv is inversely proportional to the square root of the total number of intersections counted. It is very desirable to have at least a rough estimate of k before starting an analysis, since it will then be possible to decide in advance by use of Eq. [5J how many counts should be made in order to achieve the desired accuracy in Lv. Without such information the investigator, unless lucky, will either do needless (and tedious) work or else obtain a result with too large an error. It is very easy to calculate k for the limiting case of N — 0, i.e., when the test figure is vanishingly small. In this case there is a negligible probability of two or more intersections falling within the figure. There are therefore only two possible outcomes: a count of one with a probability of, say, p and a zero count with a probability of (I-p). Thus and Therefore: We might expect k to depend on N and therefore on the size of the test figure. Such a dependence was observed2 in the analogous case involving the estimation of grain-boundary area per unit volume by a count of the number of intersections of a test circle with traces of the boundaries on the plane of polish. In this case the theoretical value of k at the limit N — 0 was 1.41. It was also predicted that k would display an initial sharp decrease with increasing N. Experimentally, k was observed to decrease to a value of 0.62 at N - 6 and thereafter remain approximately constant. Furthermore, the value of k for N > 6 was found to be insensitive to the form of the structure provided no pronounced anisotropy or grain-size variation was present. In order to determinewhether k would exhibit a similar dependence on N in the estimation of grain-

Jan 1, 1967

By Mason K. Banks

A process has been developed which produces spodumene concentrates assaying 6.0 pct Li2O and 0.45 pct Fe2O3, with 70 to 75 pct recovery of spodumene. Two flotation separations are required: the simultaneous removal of mica, feldspar, and quartz as a froth product using a cationic collector, and the subsequent removal of iron-bearing minerals as a froth product while spodumene is depressed as a tailing. SINCE the close of World War II there have been numerous developments resulting in the increased use of lithium by industry. Notable among these has been the development of an industrial lubricant, based on lithium stearate, which is fluid at low temperatures, stable at high temperature, and insoluble in water. A new bleaching agent, lithium hypo-chlorite, has also been developed recently.&apos; Other well-known uses are listed below: 1—Lithium oxide in ceramics reduces firing time, lowers melting temperatures, lowers thermal expansion, raises refractive index, and improves chemical resistivity. 2—Lithium carbonate and lithium strontium nitrate are used in pyrotechnics for their brilliant red color.&apos; 3—Lithium chlorides and fluorides are used as fluxes for welding aluminum and magnesium.&apos; 4—The affinity of lithium for nitrogen is the basis for a process used in purifying helium." 5— Lithium chlorides and bromides are used in air conditioning for their powers to absorb organic amines, ammonia, and smoke.6-—Lithium hydride has been used as a source of hydrogen to inflate life rafts and balloons used by the Air Forces and Navy.&apos; 7— Lithium metal and lithium hydride are useful in the atomic energy program in connection with the production of the isotopes of hydrogen." During the 1942-45 period of World War 11, Solvay Process Co. operated a flotation plant with a daily capacity of 300 tons of ore for the production of spodumene concentrate near Kings Mountain, North Carolina. The operation was shut down upon reduction of wartime demand for lithium chemicals, and until recently there has been no production from the Carolinas. In April 1950 an investigation of the problem of spodumene flotation from North Carolina ores was begun at North Carolina State College Minerals Research Laboratory at Asheville, N. C. This laboratory, operating under the combined auspices of North Carolina State College, North Carolina Department of Conservation and Development, and the Tennessee Valley Authority, undertook the spodumene project at the request of Foote Mineral Co. of Philadelphia. The work, involving a period of 15 months of intensive batch and pilot plant investigation, was completed in June 1951 and resulted in development of a process for concentration of spodumene by flotation which has been hereto unrecorded in the literature. A similar approach has been used before on iron, ilmenite, and other minerals, but the literature shows no record of its having been applied to spodumene. Parts of the process are now being used on a commercial scale at the Foote Mineral Co. plant near Kings Mountain, N. C., which has been in production since July 1951. It is the purpose of this paper to describe some of the results of the investigation. The Orebodies Spodumene-bearing pegmatites occur in a narrow belt 24.5 miles long and 1.8 miles wide extending southwestward from Lincolnton to Grover, N. C. According to statements judged to be conservative, the tin-spodumene belt of the Carolinas contains an estimated 4,300,000 tons of pegmatite ore averaging 15 pct spodumene. This amounts to "a reserve of at least 650,000 tons of spodumene (more than 20,000 tons of metallic lithium) to a depth of 100 ft."&apos; This figure includes only bodies having an average thickness of 35 ft and length of 550 ft. There are hundreds of pegmatites in the belt, many of which are less than 10 ft wide. The largest pegmatite in the area has a maximum thickness of 395 ft and maximum length of 3250 ft.&apos; These larger pegmatites are composed principally

Jan 1, 1954

By J. A. Roberson, A. M. Adair

A plate of an 18 pct Ni (250 grade) maraging steel was homogenized, cold rolled, and re crystallized to produce a fine grained, homogeneous structure. The strength properties were then measured as a function of cold rollinq prior to aging. Material in the solution heat treated condition exhibited substantial strengthening from cold rolling. The strength increase and the rate of strengthening were proportional to the amount of cold work. This benefit, of up to 50 ksi, was retained during aging for 3 hr at 900°F. The structures produced by cold rolling and aging have been systematically investigated by transmission electron microscopy. The dislocation debris produced by cold work influences the morphology and distribution of precipitate particles as well as the rate of precipitation. An essential feature of the high nickel maraging steels is a martensitic transformation from y to a with the resultant production of a supersaturated bcc solid solution. This solution decomposes to form precipitates during a subsequent aging treatment.&apos; The martensitic transformation in low carbon Fe-Ni alloys and the mechanical properties of these alloys have been studied extensively.2"4 The carbon content of maraging steel is kept very low to avoid the embrittling Fe3C phase and to minimize the acicular, or twinned martensite common to high carbon steels. Maraging steels are usually strengthened only by heat treatment, but cold working is also quite effective in this respect as was shown by Decker, Eash, and Goldman,&apos; Kula and Hickey,5 and the present work. There is, however, some disagreement between the present results and those of Kula and Hickey. Because of their unusual combination of strength and toughness, the maraging steels are suitable for applications where reliability at high stress levels is required. The strength level is achieved through a combination of solid solution hardening,&apos; substructure hardening,3 and precipitation hardening,6 of which the last is the most important. The toughness of these steels has been reported to be a result of the absence of brittle grain boundary precipitates,7 and the fact that the matrix is a massive martensite rather than a twinned martensite.&apos; One of the problems in the manufacture and use of maraging steel has been its tendency to develop laminations in sheets parallel to the rolling plane.9,10 The present investigation was intended to determine whether or not an improvement in the microstructure of this steel by thermomechani-cal processing would improve mechanical properties. I) EXPERIMENTAL PROCEDURES A plate ($ in. thick) of an 18 pct Ni-250 grade maraging steel* was obtained and examined metallo- "Analysis, wt pct: 18.1 Ni, 8.6 Co, 5.3 Mo, 0.4 Ti, 0.07 Mn, 0.05 Cr, 0.05 Si, 0.002B. 0.1 Cu. 0.05 Al. 0.002 Zr, 0.016 C, balance Fe. graphically. It was found to contain laminations as described by previous investigators. This plate was then annealed at 2100°F for 24 hr to promote homogeneity. Pilot studies on small coupons had shown that this was the minimum temperature required to accomplish homogenization in 24 hr. Since this treatment caused excessive grain growth, the steel was subsequently cold rolled and recrystallized. Cold rolling was accomplished at about 5 pct per pass to a total reduction of 66 pct. The standard solution annealing treatment (1 hr at 1500°F) was adequate to cause recrystallization of the cold worked steel. The recrystallized steel was then cold rolled and ground by various amounts to produce specimens with uniform thickness and selected amounts of cold rolling. All tensile specimens with less than 90 pct cold reduction were ground to a thickness of 0.050 in. Specimens with 90 pct cold reduction were 0.025 in. thick. All

Jan 1, 1970

By J. Lohrenz, R. J. Francis, G. C. Clark

Compositional material balance methods have been extended to solution gas-gas cap-water drive reservoirs. Partial or complete pressure maintenance as well as reservoir repressuring by gas and/or water injection may be considered. Increrments of reservoir withdrawals and injections equivalent to a given time increment are used as the basic iterative step to allow the extension. An example field study shows how the reservoir volumes. extent of the water drive, and relative permeability ratio function may be estimated. Knowing these, the methods can be used to predict reservoir behavior. INTRODUCTION The compositional material balance is similar to conventional material balances in many aspects. The assumption of a pressure-homogenous reservoir with no transients is common to all conventional material balances1-1 although numerical methods are different. The compositional material balance also contains this assumption. Inclusion of phase equilibria directly in reservoir calculations is the key difference and improvement in the compositional material balance. A continuous account of the compositions of the reservoir fluids and effluents is determined. Calculations of fluid phases and properties can be based on these compositions. Excepting fluid phases and properties determined in this manner, the compositional material balance uses procedures exactly analogous to conventional methods. Brinkman and Weinaug5 and Jacoby and Berry6 have given compositional material balances for simple solution gas drive reservoirs; subsequently Jacoby and Berry&apos; extended their procedure to include pressure maintenance by gas injection at constant reservoir pressure. This paper extends compositional material balance techniques to a more general combination drive reservoir with various types of pressure maintenance operations. The procedure for this extended compositional material balance and an example of its field application are presented. PROCEDURE The numerical calculations of this compositional mate- rial balance are performed on a high-speed digital computer. The most comprehensive case the present compositional material balance program can handle is a solution gas-gas cap-water drive reservoir with (1) water production, (2) water influx, (3) separate or commingled production from the oil and gas zones, (4) any surface processing arrangement including LPG and gasoline production, and (5) gas injection into the gas and/or oil zones." This description of the procedure applies to this most comprehensive case. Appropriate deletions in the procedure can be made for more simple mechanisms. ASSUMPTIONS In common with conventional material balances, the compositional material balance makes the following assumptions: (1) the total reservoir is at a uniform pressure at a given time; (2) the volumetric ratio of gas to liquid withdrawal at reservoir conditions is directly proportional to their mobilities; and (3) instead of assuming a pressure-dependent total gas-oil equilibrium, the cornpbsitional material balance makes a sianificantly different and more rigorous assumption as follows. Thermodynamic equilibrium exists for each component in the liquid-gas mixtures in the reservoir and in surface processing. These three assumptions are fundamental in the compositional material balance procedure. Other assumptions involve arbitrary simplifications of the reservoir mechanisms. This compositional material balance uses the following simplifying assumptions: (1) no bulk mixing occurs between the gas and oil zone; (2) total mixing occurs between any injected gas and the contents of the injection zone; (3) no in situ hydrocarbon liquid is produced from the gas zone; (4) no reservoir gas is retained behind any net bottom-water influx; (5) the residual saturation of hydrocarbon liquid at reservoir conditions behind any net bottom-water influx is constant; and (6) fluids retained behind any net bottom-water influx exert no further influence on the reservoir. These six assumptions are not fundamental. Modifications of the present procedure can be made for different assumptions of the reservoir mechanisms. THE TIME INCREMENT Previous compositional material balances have used a pressure increment for simple solution gas drive reservoirs"" and an increment of injected gas for pressure maintenance.&apos; These increments are suitable for these simple cases. They are, however, inadequate for reservoirs with time dependent drive mechanisms including water drive. A time increment removes this inadequacy and is

By Paul Gordon

The distribution coefficient, k, of interest in zone refining is generally defined as the ratio of the solid to the liquid solubilities of one element in another at the normal melting point of the solvent. In practice, distribution coefficients are frequently calculated from solubilities at temperatures well below the melting point of the solvent, e.g., at a eutectic temperature. This is done because of the general lack of accurate data on the liquidus and solidus curves in most phase diagrams. The practice is equivalent to assuming that the liquidus and solidus curves are straight lines in the temperature range between the melting point of the solvent and the temperature at which data for the solubilities are known. It is recognized that the k values so derived may be in serious error; we have found evidence in recent zone-refining operations that a case in point is silicon in aluminum. The phase diagram found in the literature for silicon in aluminum indicates a eutectic on the aluminum-rich side at 12.5 at. pct Si and at 577°C,1 and a solid solubility of silicon in aluminum at this temperature of 1.59 pct.2 It also shows the liquidus and solidus as nearly straight lines from 577°C to the melting point of aluminum (660°C)1,2 The k value obtained from this data is 0.13. This is almost exactly the same as the k value (about 0.14)&apos; obtained in a similar way for copper in aluminum. Yet we have found by mass-spectrograph analyses* of three bars of zone-refined aluminum that copper is efficiently removed from aluminum by zone refining but silicon is not. Two of the three bars of aluminum were zone-refined in our laboratories, the other in the laboratories of AIAG Metals, Inc. The former two bars, designated bars PG2 and PG3, were 30 in. long by 1 in. in diameter during refining; each was processed in a high-purity graphite boat* and kept under refining stage of the process consisted of thirty unidirectional zone traverses at rates of about 1 to 2 in. per hr, with a zone length of 1 to 2 in. At approximately the middle of the refining period a 6-in. piece of the impurity-laden finishing end of the bar* was cut off and replaced with new start- ing material. Following refining, the purest 20-in. section of the bar (obtained by discarding approximately 2 in. of the starting end and 8 in. of the finishing end of the bar) was leveled and then analyzed. The AIAG bar was 16 in. long and about 1.2 in. in diameter; it was purified by eight passes of a 2-in. zone at a rate of 1.6 in. per hr, the bar being contained in a sintered alumina boat lined with loose alumina powder of the highest purity. Data for the silicon and copper contents of the three refined bars are given in Table I. Derived data for the efficiencies of the zone-refining operations are listed in Table 11. It is clear from these two tables that the copper is removed much more efficiently—by a factor of 10 to 50 times—from aluminum during zone refining than is silicon. For example, in bar PG3 the thirty zone passes reduced the copper content from 10 to 0.02 ppm, i.e., by 500 times, whereas the same treatment lowered the silicon content by only a factor of about 16. It seem virtually certain, therefore, that the distribution co efficient, k, for silicon in aluminum is well above that for copper and is very likely not much below unity. This also implies that the solidus line on the aluminum-rich side of the A1-Si phase diagram is very probably appreciably curved near the aluminum melting point since the phase-diagram data indicate such curvature is unlikely in the liquidus.

Jan 1, 1965

By Gordon H. Geiger, John M. Svoboda

One mechanism of metal penetration into mold aggregates by cast steels is vapor state mass transport. In order to further understand and quantify this mechanism, the steady-state diffusion of metallic vapors was studied at 1250" to 1600°C in controlled oxygen pressure atmospheres by measuring the weight loss from a metal bath whose vapors were allowed to diffuse through a porous aggregate compact held at constant temperature. Mass transport rates substantially greater than those predicted from the assumption that the vaporizing species was the elemental metal vapor were observed. Vaporization rates were found to increase with an increase in the oxygen partial pressure in the gas atmosphere with maximum vaporization occurring at oxygen partial pressures just below those in equilibrium with a liquid oxide layer at the given temperature. When the liquid oxide layer was present, no vaporization was observed. Metal carbonyls were ruled out as the vapor species involved because of their thermodynamic properties. The species with the enhanced volatility were deduced to be gaseous metal-oxygen molecules, probably the polynuclear oxides (FeO)2 and (Mn0)2. Diffusion followed the ordinary diffusion mode, and Knudsen diffusion is not important in normal foundry molds. The interdiffusivities of the metallic vapors and the mold atmosphere, through the porous aggregates, were measured. The effective interdif-fusivity of the vapor species and mold atmosphere through the aggregate may be expressed as follows: Deff,12 - — where E is the void fraction of the aggregate and 7 is the "tortuosity factor". Aggregate material had no effect on diffusion rates, and the only relevant parameter was the void fraction. Values for 7 ranged from 3 to 6 for normal foundry aggregates with 4 being considered a good representative value. An area of interest in high-temperature metallurgical processes which has received little attention in the past is the diffusion of metal vapor species in porous aggregates. Research into the metal penetration phenomenon&apos; in foundry molds has shown that a major penetration mechanism is the vapor transport of metallic species into the porous foundry sand molds, followed by condensation of metallic films on the sand grains. Considerable study of the vaporization of metallic species has been reported in the literature, and the vapor transport properties of many metallic species have been determined.2&apos;3 Many instances have been observed where the vaporization behavior could only be explained by the existence of a metallic species with a vapor pressure greater than the elemental metallic species. These species were identified or, in some cases, presumed to be carbonyls,4,5 hydroxides,6-9 or vaporous oxides.10-16 The diffusion of gases through porous aggregates has received considerable attention in the literature pertaining to catalysis, although primarily at low temperature, and more recently in the literature concerned with ore reduction.17-22 The following experiments were devised and carried out in order to determine the nature of the metallic species involved in the metal penetration phenomenon, and the effective diffusivities involved. EXPERIMENTAL APPROACH Steady-state diffusion was selected for experimental convenience. The experimental systems used, shown in Figs. 1 and 2, consisted of a porous aggregate specimen contained in an impermeable alumina tube held above, and just separate from a metal bath. The entire system was held at constant temperature and pressure in a furnace with a controlled oxygen partial pressure atmosphere. After a brief period of nonsteady state behavior, a steady flux of metal vapor through the porous aggregate will develop due to the concentration gradient of metallic vapor in the gas phase. The flux of metallic vapor from pressure may be neglected because there is no pressure gradient. There is no flux from convection because there is no net flow of gas within the metal-aggregate system, and thermal diffusion is negligible compared to the flux from ordinary diffusion. If the material being studied is allowed to vaporize until steady-state behavior is attained, the experimental results may be analyzed using Fick&apos;s First Law.

Jan 1, 1970

By S. Weinig, E. S. Machlin

IN a previous study of the grain boundary stress relaxation phenomenon,&apos; the authors had arrived at the conclusion that two successive steps were involved in the complete relaxation of stress at a grain boundary. These two steps were grain boundary migration and grain boundary slip. It was further suggested that the slower of these two processes is rate controlling. In an attempt to ascertain the validity of the above, a study of grain growth kinetics on the same alloys used for the previous study was undertaken. Separate wire specimens, 0.030 in. diam, in the same as-deformed state (99 pct reduction in area) as in the previous investigation were heated to the various temperatures for a sequence of times. The compositions used and the temperatures of testing are summarized in Table I. A statistical determination of grain size was then made on each specimen, using the average length of traverse between boundaries as a measure of the grain size. Examples of typical results obtained are shown in Figs. &apos;1 and 2. The grain growth law D = ktn was found to fit most of the data. Using this relation, values of the grain growth exponent n were calculated and the inter -esting dependence of n on solute content shown in Figs. 3 and 4 was then obtained. Also, it was found that the constant K which was obtained from extrapolation is roughly independent of solute content. (K is uncertain to within ±25 pct.) An estimate of the activation energy for grain growth was made using the technique of Burke? This method utilizes the plot of the reciprocal of time necessary for the grain size to increase from a definite starting size Do, to a final size D, for each isothermal heat treatment. The slope of this curve corresponds to the heat of activation for the process. In Fig. 5 the rate of grain growth is plotted against the reciprocal of absolute temperature. The activation energies were obtained from these curves and are summarized in Fig. 6. Discussion It is interesting to note that the same saturation in the variability with concentration has been found for the grain growth exponent n in this research as was found in the previous researches on both grain boundary stress relaxationa and internal friction of dislocation origin." The significant feature is that saturation in the variation of these properties has been found to occur at about the same level of solute concentration, namely, between 0.1 and 0.5 atomic pct. The only characteristic common to these different quantities is that solute adsorption at grain boundaries and dislocation can take place and affect these values in a similar fashion. Thus, it is concluded that the grain growth exponent n depends upon solute adsorption at the grain boundaries. It appears, therefore, that theoretical considerations leading to the contrary conclusions5,6 have not treated the real process by which solute atoms exert their effect on grain growth. Prior to the analysis of the observed activation energies, certain of the aspects of the aforementioned grain boundary stress relaxation investigation will be briefly reviewed. It was found that with the addition of solute atoms the grain boundary stress relaxation peak (internal friction vs temperature) decreased in intensity. Simultaneously a second peak was observed which increased with increasing solute atoms. Both of these peaks were shown to be the result of grain boundary stress relaxation. They have been termed, respectively, poire peak and solute peak. It was suggested that the solute peak was dependent upon grain boundary migration, whereas the pure peak was a manifestation of grain boundry slip. Activation energies for both peaks were obtained as a function of Solute content. The value of the activation energy in the temperature range corresponding to the solute peak was

Jan 1, 1958

By A. Rosen, P. Wynblatt, J. E. Dorn

The temperature, strain rate, and structure dependence of the flow stress in a poly crystalline Fe-2 pct Mn alloy was investigated between 77" and 370°K. It was possible to identify the low-temperature data from 77° to 160°K with Dorn and Rajnak&apos;s theory of the Peierls mechanism of plastic deformation. Above 160°K one or more thermally activated mechanisms must he operative; however, these were not identified in this investigation. The mechanistic basis for the rapid decrease in the flow stress of bcc metals with an increase in temperature over the low-temperature range is yet being debated. Various dislocation mechanisms have been suggested in the past but the agreement between the various proposed theories and the experimental results have not been definitive. In the present paper, the authors will demonstrate that the plastic behavior of polycrystalline iron containing 2 wt pct Mn plus interstitials agrees excellently with predictions based on the Peierls mechanism from 77" to 160°K whereas from 160" to about 370°K other as yet unidentified thermally activated mechanisms take over. Undoubtedly, the overlap of these mechanisms with the Peierls mechanism contributed to the difficulties of rationalizing the low-temperature behavior of iron and its alloys. I) EXPERIMENTAL TECHNIQUES AND RESULTS Although extensive data are now available on the effect of temperature and strain rate on the flow stress of various bcc metals and alloys, it was, nevertheless, deemed necessary for the objectives of the present study to obtain highly accurate data of the type most pertinent for checking the theories on a series of well-documented work-hardened states. Iron containing 2 wt pct Mn was used in this investigation in order to obtain a lower ductile to brittle transition temperature than that for iron.&apos; The additional elements present were 0.004 pct C, 0.05 pct 0, 0.004 pct S, 0.003 pct P, 0.006 pct N, and 0.001 pct Si. The as-received 3/8 by 3/4 in. bars were cold-rolled to 0.100-in. thickness, recrystallized under argon for 30 min at 1073"~, further cold-rolled to 0.063-in. thickness, and finally recrystallized under argon for 30 min at 1000°K. This treatment gave a stable reproducible equiaxed grain structure having an average grain diameter of about 90 p. Prior to the final recrystallization treatment, the sheet was machined into tensile specimens 1/4 in. wide having a 1.625-in.-long gage section. All specimens were tested in controlled-temperature baths on an Instron Testing Machine. For convenience in the theoretical analyses the data will be given in terms of the shear stress for flow, 7, the shear strain rate, f, and the shear strain, y. A series of three standard strain-hardened states were obtained by prestraining specimens at 300°K and at a shear strain rate of 7.68 x lo-&apos; per sec to three stress levels of 6.89 X l08, 10.3 X lo8, and 14.5 x 10&apos; dynes per sq cm, respectively. Immediately following this prestrain the temperature bath was changed and each specimen was pulled in tension at either f = 7.68 X lo-&apos; or f = 3.06 X 10"3 per sec. In terms of the accuracy that was achieved, the values obtained did not deviate more than *0.1 X 10&apos; dynes per sq cm and * 1°K from the reported values of stress and temperature and not more than *5 pct in the reported values of the strain rate. The tensile stresses were converted to shear stresses using a factor of 1/2. The shear stress for the initiation of flow at the test temperature was determined by taking a y = 6.16 x 10&apos;3 offset. from the modulus line at which point good accuracy could be achieved in determining y and the associated flow stress 7, with only a negligible increase in the pre-

Jan 1, 1965

By C. C. Templeton, S. S. Rushing

Methods previously developed for the study of air-liquid displacements in microscopic capillaries (inner diameters of 3 to 40 microns) have been used to investigate oil-water displacements in capillaries initially filled with water. Displacement calculations assuming perfect displacement and no capillary pressure hysteresis yielded oil effective viscosities smaller than the macroscopic viscosities. For a given liquid pair, the oil effective viscosity decreased both with decreasing capillary size and with increasing oil-water viscosity ratio. This behavior can be explained by the existence of an annular water film (20 A to 260 A thick) on the capillary wall. When the capillary was first filled with oil, the ratio of the oil effective viscosity to the normal oil vircosity was highest for the first water displacement and decreased with subsequent displacements. Sometimes the oil effective viscosity ratio during the initial water displacement was greater than unity. INTRODUCTION In a previous paper1 a technique was described for studying air-liquid and liquid-liquid displacements in very small capillaries of uniform diameter, in the hope that such microscopic data would further the understanding of the nature of multiphase fluid flow through porous media. That paper contained comprehensive data for air-liquid displacements in Pyrex capillaries, with a few data for oil-water displacements in capillaries initially filled with water. The purpose of this paper is to present more complete results for oil-water displacements in capillaries initially filled with water, and to describe for the first time such observations in capillaries initially filled with oil. In this way the effect of the wetting history of the system upon the displacement process may be studied. METHODS The basic techniques employed were described in the previous paper.1 The measuring procedure and the working equations will be briefly summarized here and a few modilications will be pointed out. For the present study, temperature control within 0.l °C was obtained by placing the microscope and its immediate accessories in a thermostated air bath made of a steel frame covered with plexiglas. The water and oils used in this work were the same as in the previous paper,1 with the addition of "Medium Mineral Oil, U.S.P. supplied by the Harshaw Scientific Co. When the first liquid was introduced into the horizontal capillary, the air-liquid static capillary pressure Poc was measured along the observable length of the capillary. From the relation Poc = 4y cos 0/d, the capillary diameter d could be calculated from Poc if the "microscopic" air-liquid boundary tension, To = y cos On, was taken as equal to its known "macroscopic" value. This calculation involves assuming that cos 6 = 1 and that capillary size or the resulting high interfa-cial curvature has no effect on surface tension, or y cos 6, as was verified by our work, and also by that of Cohan and Meyers on air-liquid-solid systems. For several capillaries these calculated diameter values were compared with values measured visually with a filar micrometer eyepiece. For a given capillary the average values for the two methods agreed within 1 or 2 per cent, but the visual values were less reproducible than the calculated values. Further the calculated values were much less subject to personal bias and took into account any slight ellipticity that might be present. Hence diameters calculated from air-liquid Poc data were used throughout this paper. The second liquid was introduced in such a way that a single spherical interface separated the two liquids. Displacements were always made with a constant pressure, Pt, between the ends of the capillary. The times (t) at which the interface passed certain

Jan 1, 1957

By A. E. Powers, J. F. Libsch

THE simultaneous influence of time and temperature upon the tempering process in steel is well known. Hollomon and Jaffe¹ have expressed the effect of time and temperature upon the progress of tempering by a single parameter, T (c + log t); where T is the absolute temperature; t, the tempering time; and c, a constant characteristic of the steel and treatment. It was found that when hardness data for plain carbon and low alloy steels were plotted vs. this parameter, a master tempering curve was obtained which satisfied all the data for a given steel and a given hardening treatment. Roberts, Grobe, and Moersch² successfully applied the tempering parameter to various tool steels, including high speed steel for tempering times ranging from 6 min to 100 hr. The influence of tempering time and temperature upon the physical properties of high speed steel has also been extensively investigated.&apos;" Recently, the use of induction heating for very rapid tempering treatments has received attention. The authors% have shown that tempering treatments of only a few seconds duration may be successfully applied to plain carbon and low alloy steel. However, the influence of such short tempering treatments in high speed steel has not been explored. Since high speed steel is a highly alloyed steel exhibiting physical reactions during tempering that are often relatively slow and large in magnitude, it is an interesting steel for the study of tempering reactions. It was the purpose of the present investigation, therefore, to study the early progress of reactions leading to secondary hardness in high speed steel, type 6-5-4-2, and to determine if induction-tempering treatments of a few seconds duration are able to produce results equal to those of more conventional long-time tempering cycles. Experimental Details The chemical compositions for two different heats of W-Mo high speed steel used in this study are given in Table I. Table II summarizes the heat treatment given various groups of specimens cut 21/2 in. long and centerless ground to ¼ in. diam. Hardening treatments consisted of preheating to 1600°F (870°C) for V2 hr and then oil quenching from 2250°F (1230°C). Heat I was subjected to the high temperature in a salt bath for 3 min, while specimens of heat II were austenitized in a carbon-block muffle furnace for 10 min. Specimens of series E were immediately refrigerated to — 318°F (— 195°C) in liquid nitrogen following the hardening treatment to transform the major portion of the retained aus-tenite. Specimens of series F and G were subzero cooled to — 100°F (-75°C) in a dry ice-alcohol mixture. Induction tempering was performed with a Lepel high frequency converter operating at a frequency of approximately 300,000 cycles per sec. Heating cycles were controlled and recorded for each specimen using a chromel-alumel thermocouple, percussion welded to the surface of the specimen, and a modified high speed instrument with a timer previously used and described.&apos; Temperature oscillation of the specimens, while being held at the tempering temperature, was ± 15°F (8°C). Tempering Treatments: After hardening, specimens were tempered between 900" and 1500°F (480" and 815°C) to study the progress of retained austenite decomposition and secondary hardening. The specimens tempered in salt for 1 hr (series C, G, and I) were considered to represent conventional treatment and are used for comparison with rapid tempering treatments. Several induction-tempering cycles were employed: 1—flash tempering, i.e., induction heating to the tempering temperature and immediately cooling (series A, E, and F); 2—induction heating for 60 sec at temperature (series B); and 3—double-induction tempering by two successive heating and cooling cycles (series D). All specimens tempered by induction were heated at an initial rate of 450°F (250°C) per sec except those of series E, which were heated at an initial rate of 1300 °F

Jan 1, 1952

By M. Ahlers, Horace Pops

Slip features and formation of stress-induced mm-tensite have been studied in single crystals of fl phase Cu-Zn alloys in the composition range between 43 and 48 at. pct Zn. A phenomenological analysis is presented in which the glide planes and directions that me associated with the inhomogeneous shear in the martensite phase are predicted from the experimentally determined habit planes and macroscopic shear directions. It is shozun that the resolved shew component of the applied stress on the secondary shew system is a very important factor governing the formation of stress-induced martensite. A comparison is made between martensitic phases that form spontaneously upon cooling and those phases that arise from deformation. THE ß phase in Cu-Zn alloys is unstable and may transform spontaneously by a martensitic transformation1-3 during cooling. Thermoelastic and burst-type spontaneous martensitic phases have been observed3 in the composition range between 38 and 42 at. pct Zn. Martensite forms in alloys containing more than 42 at. pct Zn only by cold work below a temperature denoted as Md.4 Thus, alloys near the equiatomic composition can undergo a stress-induced transformation upon cold work at room temperature. Although data have been obtained concerning the effect of temperature5 and composition4 on the deformation of polycrystalline Cu-Zn alloys, it is rather surprising that few detailed studies6-8 on 6 brass have been conducted for single crystals where the stresses acting on crystallographic planes can be determined. A thorough understanding of transformation mechanisms requires knowledge of the role of stress and plastic strain. With this objective in mind we have made an investigation of the stress-induced rnartensite transformation in single crystals of 0 brass in the composition range between 43 and 48 at. pct Zn. Particular emphasis is given to the morphology and crystallographic features and a comparison is made between stress-induced and spontaneous martensite. I) EXPERIMENTAL PROCEDURE Master alloys having nominal zinc contents of 43, 45, and 48 at. pct were prepared from spectroscop-ically pure metals (>99.999 pct) in sealed quartz tubes under a partial pressure of helium. These alloys were cut into several pieces and converted into small rods by melting in quartz tubes. Subsequently, single crystals were grown from the rods by the Bridgman technique and homogenized for 3 hr at 820°C. Lattice parameter measurements were made from stress- relieved filings. Compositions, as determined from the reported variation of lattice parameter with atomic percent zinc,&apos; indicated a zinc loss of less than approximately 0.25 at. pct. The single crystals were electroturned using a phosphoric acid electrolyte into tensile specimens each with a center section of 1 mm diam and 1 cm gage length. Cylindrical specimens approximately 3 mm in diam and 10 mm long and spark-machined rectangular blocks approximately 3 by 3 by 10 mm were prepared for compression studies. These blocks were stress-relieved and electropolished to eliminate any stresses introduced by spark machining. Crystal orientations were determined to an accuracy of ±2 deg primarily from back-reflection Laue photographs, and sometimes by the rotating single-crystal method. Figs. 1 to 3 show the location of the tension and compression axes within unit stereographic triangles. All specimens were deformed plastically at room temperature using an Instron machine, the crosshead speed being 0.02 or 0.05 cm per min. A hard steel hemisphere was placed at the upper end of the compression samples to achieve uniaxial loading. Trace normals of slip markings and habit planes of strain-induced martensite were determined to an accuracy of ±4 deg from the rectangular blocks by a two-surface analysis and from cylindrical specimens by measuring the angle between the tensile (or compression) axis and the tangent to the surface trace for different rotations about the deformation axis. Slip

Jan 1, 1969

By O. J. C. Runnals

The intermetallic compounds NpAl2, NpAl3, and NpAl have been prepared, and examined by X-ray diffraction methods. The compounds are isostructural with the corresponding U-Al compounds. NpAl3 is face-centered cubic with a = 7.785A and has the MgCu2 structure. NpAl is simple cubic with a = 4.2624 and has the AuCU3 structure. NpAl4 is body-centered orthorhombic with a = 4.428, b = 6.268, and c = 13.71A. The published description for the atomic positions in UAl, applies also to NpAl4 THE U-Al system has been examined by Gordon and Kaufmann.1 They identified the following intermetallic compounds: UAI2, UAl3, and a third, tentatively identified as UAl5. A structure analysis of the latter compound by Borie2 proved that the formula UAl4 was more consistent with the X-ray data. The crystal structures of UAl2 and UA have been reported by Rundle and Wilson.3 UAl2 was found to be isostructural with MgCu. The atomic positions in the simple cubic UA1, cell were found similar to those in the AuCu3 structure. Three analogous compounds exist in the binary system Np-Al. X-ray data for the compounds NpAl2, NpAl3, and NpAl4 are reported below. The neptunium used in the investigation was supplied through the courtesy of the Argonne National Laboratory as neptunium (V) nitrate. It was stated to have a radioactive content of 99.25 pct Np237. The impurity limits had been established by a spectrographic analysis given in Table I. Any titanium, boron, or bismuth impurity should be effectively removed during the fluorination of NpO2 to NpF4, and sodium, lithium, magnesium or strontium during the alloying procedure. The oxalate ignition technique, described by Westrum for the preparation of finely-divided reactive PuO2, was used.The neptunium (V) nitrate solution, containing 100 mg Np, was evaporated to dryness on oxalic acid crystals. The resulting oxalate was calcined to NpO, by heating in platinum at 650°C for 2 hr. An equimolar mixture of anhydrous hydrogen fluoride and purified hydrogen was passed over the NpO2 for 2 hr at 500°C. According to Fried and Davidson" the above fluorination conditions should produce the purple-black fluoride, NpF3 However, the light-green product showed only NpF4 lines on an X-ray diffraction pattern and contained only an estimated 5 pct of finely dispersed black NpF3 particles on microscopic examination. NpAl3: An intimate mixture of 50 mg of neptunium fluoride and aluminum powder (99.99 pct pure, —28 to $60 mesh) was heated to 900°C in a BeO crucible in vacuum by a tungsten coil for 2 hr, to allow the following reaction to proceed: NPF4 + 7Al + NpAl, + 4AlF. [1] The formation and sublimation of aluminum mono-fluoride in this system was verified from an X-ray diffraction analysis of the sublimate. The diffraction pattern was similar to one obtained in the aluminum reduction of UF4 to UF3,6 showing both Al and AlF3 lines. The resulting alloy was sintered at 1100°C for 1 hr. A powdered specimen of the crystalline product was sealed in a thin-walled quartz capillary for X-ray examination. The photograph was taken on a 14.32 cm General Electric powder diffraction camera, using filtered copper radiation. The films were calibrated in terms of a NaCl standard pattern. The wavelengths used in the calculations were 1.5418A for Cu K and 1.5405A for Cu K The powder diffraction pattern showed that NpAl is simple cubic with a = 4.262 % 0.005A, and is isostructural with UA13 where a = 4.27A1 or a =

Jan 1, 1954

By Donald L. Katz

The compressing of gas into a partially depleted gas drive oil reservoir to bring the contents to a single phase miscible with gas is proposed as a process worthy of serious study. The compressed gas and vaporized oil would be recovered by cycling. This paper explores the Possibility of recovering oil from reservoirs of 6,000 ft or more by this method. INTRODUCTION Many new methods for increasing the recovery of oil from reservoirs are under consideration in the various research laboratories. One method which should be considered for recovering oil From partially depleted gas drive fields is to compress gas into the reservoir until the contents become a single phase miscible with more gas. Natural gas then may be cycled through the reservoir to recover the oil. This procedure ideally will eliminate the capillary effects of the liquid ill the porous media and should permit recoveries similar to those obtained in cycling operations, probably of the order of 80 per cent. This paper explores the information now available for the prediction of the performance of such a project and discusses the problems likely to be encountered in recovering oil by cycling. For many years dry natural gas has been passed through oil reservoirs to recover natural gasoline which would vaporize from the oil. Recently a process of recovering more of the high boiling hydrocarbons has been described by employing high pressures at which the heavier hydrocarbons have a higher volatility. Nevertheless. this reported method does not eliminate the differential nature of the vaporization process and capillary forces are still present in the reservoir. The process under discussion in this paper would minimize differential vaporization and would remove capillary forces which are responsible for low oil recoveries. Higher gas pressures are required to accomplish this removal of the miniscus, but it does not follow that more gas is required for oil recovery. The process is limited to reservoirs deep enough to permit the use of the required pressure — Probabably 6,000 it and deeper. DESCRIPTION OF THE PROCESS Briefly, the process is to compress natural gas into an oil reservoir which is partially depleted so that it has an oil saturation in the range of 40 to 60 per cent of the hydrocarbon pore space. The injected gas will pass through the channels formerly carrying gas to the wells to build up the pressure and dissolve in the oil phase. This injection will continue without production from the field until the reservoir reaches a single phase. If one were to select unique conditions such that the gas composition, the reservoir oil composition and the Per cent saturation would bring the mixture in the reservoir to the critical temperature and pressure simultaneously the phase behavior accompanying injection is shown by Fig. 1. The oil-gas mixture in the depleted reservoir has a critical temperature at C,. Addition of natural gas to the mixture will give a new critical temperature for each addition. Mixture No. 3 is shown to have its critical temperature at the reservoir temperature and reservoir pressure at P: the critical pressure. Under ideal behavior, the entire reservoir would be at iti critical temperature and pressure and the entire hydrocarbon content would be at a single phase. If the reservoir content were a binary system of methane and a heavy constituent. it would be miscible with more methane. However, for complex mixtures the critical locus does not confine the two phase region, and the dew point curves of Fig. 1 occur above the critical locus. If fluid which was injected into the rezervoir at the critical point were of the

Jan 1, 1952

By M. T. Simnad, C. E. Birchenall, M. H. Davies

THE marker movements observed by Davies, Simnad, and Birchenall during the growth of magnetite on wustite have been misinterpreted. It is the purpose of this note to correct the original interpretation. In that investigation wustite specimens were prepared by oxidizing pure iron disks in water vapor until the iron was consumed. Argon was used as a carrier gas for the water vapor and on the complete conversion of the metal to its oxides (&apos;FeO&apos; plus a small amount (10 pct) of surface Fe3O) the water supply was bypassed and the specimens were annealed in argon alone in order to equilibrate the wustite with the overlying magnetite. After cooling, the magnetite was ground off, a crude magnetic test being employed to determine its complete removal. Radioactive silver streaks were then applied to the surface and the specimen was oxidized completely in water vapor to prevent hematite formation. After complete oxidation the markers were found to enclose about 80 pct of the magnetite with the other 20 pct lying outside the markers. This observation was interpreted to mean that oxygen transport was four times as great as iron transport during the growth of the magnetite. The conclusion would have been correct if, as in the case of wustite growing on iron, all the iron had started from one side of the marker and all the oxygen from the other side initially. Actually a material balance is required. For simplicity, take wustite to be exactly FeO, and magnetite to be Fe3O,. Then at the start of the reaction 3 FeO + O -> Fe,O., All of the iron and three-fourths of the oxygen of the final product will be within the marker inter- faces. The other fourth of the oxygen is in the gas phase outside the marker. From the experimentally determined final position of the markers simple calculation shows that one-fifth of the total iron has crossed the marker interface, while one-twentieth of the total oxygen has crossed in the reverse direction. The simplified result favors iron transport in the ratio four to one. However wiistite is not stoichiometric FeO. At 900°C, the temperature of the measurement, the composition is about FeO When the material balance for the same experimental result is repeated, assuming the product will be 1 mol of Fe3O4, it is found that 0.6 gram atoms of iron and 0.19 gram atoms of oxygen must have passed the marker in the same direction. (This is readily apparent from the reaction to form 1 mol of Fe3O4 SFeO The total oxygen supplied from the atmosphere is 0.61 gram atoms, but the iron which passes the marker interface, i.e., 0.6 gram atoms, requires 0.8 gram atoms of oxygen to convert it to Fe,O,.) If all the transport during magnetite growth is due to iron ion diffusion, then the marker would enclose about 85 pct of the final magnetite and the original measurement is in error by about 5 pct. This error is probably less than the uncertainty in locating the markers by autoradiography. The possibility also exists that the starting wustite was not quite saturated with oxygen and underwent further enrichment with consequent envelopment of the marker before magnetite caught up with the marker interface. In any case it appears that these measurements actually favor iron ion diffusion in magnetite overwhelmingly as the factor controlling its growth, although further experimental measurements are needed to eliminate the error in material balance described. Such studies are planned.

Jan 1, 1954

By C. H. Li, S. H. Bendel, J. A. Sartell

Small additions of Mg to a 62 Cu-38 Ni alloy reduce the oxidation yates drastically between 500° and 850°C. Metallographic examination of the oxide scale disclosed a continuous network, identified by X-ray diffraction as (Ni, Mg)O solid solutions, located along the former alloy grain boundaries. This network acts as a barrier to diffusion. It is thought to result from the segregation of mg to the alloy grain boundaries. Oxidation occurs preferentially in these regions because of the presence of this more active element. To substantiate these ideas, homogenized alloys ulere oxidized. No oxide network was produced and oxidation rates were the same as for the parent 62 Cu-38 Ni alloy. An interesting aspect of the oxidation of the 62 Cu-38 Ni alloy previously described by the present authors, is the relationship between oxidation kinetics and changes in the microstructure of the oxide.&apos; The oxide scale on these alloys consists of two layers: a thin layer of CuO outside a much thicker layer of Cu20 in which particles of NiO are dispersed. In addition, immediately beneath the oxide-metal interface there is an internally oxidized zone in which NiO particles are dispersed in a copper-rich alloy matrix. Experimental observations on the growth of the individual layers under isothermal conditions have led to the conclusion that the rate-determining step in the oxidation of this alloy is the diffusion of anions through the Cu20 layer. It has been found that this diffusion rate is increased by the recrystallization of the outer CuO layer. Willardson and Gorton have reported that the addition of magnesium to the 62 Cu-38 Ni alloy causes a decrease in the oxidation rate.2 The present investigation was undertaken in order to determine the mechanism by which magnesium additions improve the oxidation resistance of the 62 Cu-38 Ni alloy. EXPERIMENTAL PROCEDURE Fifty-gram ingots of Cu-Ni alloys were prepared by vacuum melting materials of 99.999+ pct purity. These ingots were then remelted in a He atmos- phere and magnesium added in the form of a previously prepared and analyzed Cu-Mg master alloy. This procedure was necessitated by the high reactivity of pure magnesium at high temperatures. The proportion of the alloying elements was adjusted so as to obtain alloys with the desired magnesium content while maintaining a Cu-Ni ratio of 62 :38 by weight. (All alloy compositions will be given in weight percent.) Since it was considered desirable to study the oxidation of only single-phase, solid-solution alloys, it was necessary first to determine the approximate limit of solid solubility of magnesium in the 62 Cu-38 Ni alloy. A series of melts with various magnesium contents was made, homogenized by annealing for 48 hr at 975OC, then cooled in the furnace. Metallographic samples were prepared, etched in a potassium dichromate solution and examined at XI500 for the presence of a second phase. The alloys were then chemically analyzed and it was found that the approximate limit of solid solubility of magnesium in the 62 Cu-38 Ni alloy was 1 pct Mg. The ingots containing less than 1 pct Mg (0.19, 0.29, 0.58, and 0.87 pct Mg) were cold-rolled down to thicknesses suitable for oxidation experiments and specimens for gravimetric studies were cut to dimensions of 3.0 by 1.5 by 0.05 cm, specimens for inert marker studies, 0.6 by 0.6 by 0.17 cm. These were heat-treated in a protective atmosphere at a temperature (975" C) well above the oxidation temperatures and furnace-cooled to insure structural stability of the alloy during oxidation. Immediately prior to oxidation, the specimens were chemically cleaned by immersion for 1 min in a solution of 2 parts nitric acid, 1 part acetic

Jan 1, 1961

By W. R. Hibbard

It has been proposed1 on the basis of slip and flow that the ideal deformation texture of drawn wire for face-centered cubic metals is a (111) direction parallel to the wire axis. Under these considerations, an intermediate [loo] texture is permissible but should eventually change to the stable [Ill] orientation. This theory does not explain why copper and its alloys containing up to about 30 pct nickel, 5 pct zinc and 1 pct aluminum were found to have a double [Ill], [100] texture, while alloys with at least 8 pct zinc and 2 pct aluminum developed essentially a single [Ill] texture.1,2 It does suggest, however, that the [loo] component of the former alloys should disappear with additional deformation.3 Some evidence exists to support this suggestion. Howald4 reported that annealed copper wire previously drawn 93.8 pct and 98.5 pct developed a single [Ill] texture. Since, in wire, the drawn and annealed textures are similar, these wires probably had a single [Ill] deformation texture before annealing. In addition, although cold rolled polycrystalline copper and 80-20 brass were found to form different types of pole figures,5,6 single crystals of these metals when rolled from the stable (110) [112] orientation were found to produce essentially the same polo figure.7.8 It may be, therefore, that while the addition of solute elements changes the path by which the final texture is obtained, the final texture is not changed. Experimental Results The wires remaining from the previous investigation1 already drawn about 50 pct reduction in diameter were further drawn to 86 pct and then to 96.4 pct total reduction in diameter. The textures were determined after these stages by transmission X ray photograms taken with a copper K-alpha beam rotated 0 degrees toward the wire axis. It was found that the ratio of the [111].] and [100] intensities of the double texture in copper and its alloys containing 0.94 pct aluminum, 30.70 pct nickel, 2.61 pct zinc and 4.99 pct zinc increased from about 3:l for the 50 pct reduction to about 9:l for the 86 pct reduction. After 96.4 pct reduction, these wires exhibited essentially a single [lll] texture. Summary Thus, the drawn wire deformation texture of copper and its alpha solid solution alloys containing aluminum, nickel or zinc is essentially a single [Ill] texture as predicted by theoretical considerations of slip and flow.&apos; For copper and its alloys containing up to about 1 pct aluminum, 5 pct zinc and at least 30 pct nickel, an intermediate [loo] component is found after about 50 pct reduction in diameter, but disappears after about 95 pct reduction in diameter. For alpha solid solution alloys containing at least about 2 pct aluminum and 8 pct zinc the single [Ill] texture is developed considerably more easily and is found after about 50 pct reduction in diameter. No explanation of this compositional effect is apparent. Similar work is contemplated for cold rolled and cold compressed alloys. References 1. w. R. Hibbard, Jr. and M. K. Yen: Trans. AIME 175, 126. Metals Tech., Feb. 1948, TP 2334. 2. W. R. Hibbard, Jr. and D. E. Trout, 11: Trans. AIME 185, 620. Jnl. of Melals, Sept. 1949, TN 18. 3. W. R. Hibbard. Jr.: Disc. of R. M. Brick&apos;s paper on Deformation Textures, ASM Symp. on Cold Working of Metals, Oct. 1948. 4. T. S. Howald: Disc. of Ref. 1. Trans. AIME, 175, 126. Melals Tech. Aug. 1948, TP 2449. 5. R. M. Brick, D. L. Martin and R. P. Angier: Trans. ASM (1943) 31, 675. 6. 0. Dahl and F. Pawlek: Zlsch. Melall. (1936) 28, 266. 7. C. S. Barrett and F. W. Steadman: Trans. AIME (1942) 147, 57. 8. J. E. Burke and C. S. Barrett: Trans. AIME 175, 106. Melals Tech., Feb. 1948. TP 2327.

Jan 1, 1950

By W. J. Rogers, L. J. Dwiggins, R. O. Williams

THERE has been comparatively little work on the importance of strain rate and temperature as variables in the formation of deformation textures. For this reason the present work was started, the choice of compression being in part that of convenience and in part because of the absence of quantitative data in this area. These do not turn out to be important variables within the range studied which permits the results to be condensed into this note. Samples 1/4 in. in diam and 1/2 in. long were machined from 99.9 pct pure commercial aluminum and recrystallized to a fine grain size. A set of five samples were deformed between a spring-driven weight and an anvil resulting in roughly 20 pct strain each time. The strain rate varied because of decle-ration and shortening of the sample but would average about 10&apos; per min. Deformation at temperatures over the range of -76" to 150°C were obtained by heating or cooling the samples. Complete lubrication could be obtained by oil since the short time of contact did not permit its escape. The samples were compressed to a total true strain of 1.4 to 1.6. A larger sample was prepared similarly with spiral grooves in the ends to hold paraffin which provided essentially complete lubrication. This sam ple was compressed slowly at room temperature at a rate of about 2 per min to a total strain of 1.0. Cylindrical samples were carefully machined from these discs with a radial direction as an axis. They came from the center half of the samples and were roughly 1/16 in. in diam. The samples were etched down 15 pct to remove any surface effects. Diffraction data were then obtained for six different kinds of planes by the use of the Norton technique as applied to a fiber texture. Direct comparison between the raw data failed to reveal any significant difference between any of the samples and consequently the data from only one sample was completely reduced. The particular sample had been deformed at O°C The background was subtracted from the intensity data which were averaged over the four quadrants and then normalized to random. These data were then used to obtain fiber-axis distribution charts as previously described.2 For this sample the (422) plane was also run to assist in this inversion. The final result, Fig. 1, provided a comparable fit to that previously obtained from fiber data. About 85 pct of the material forms an elliptical texture around the <110> direction. The radius (at half height) is 5.5 deg along the (111) boundary and 8 deg along the (100) boundary. The remaining 15 pct forms a ridge running from this texture to near the <311> direction. The regions around the (100) and (111) planes are devoid of material. The earlier qualitative results of ~arrett~ are in complete agreement with the present results. The possibility has been dis-cused that the material around the <311> results from friction between the sample and compression plates.4 The present results are considered proof that such is not the case because of the uniformity of sample deformation and the requirement of axial symmetry before fiber-axis distribution charts can be accurately obtained. Recently results have been obtained to show that strain rate and temperature are important variables in the extrusion of aluminum.5 Perhaps the most important observation was that the partial recrystal-lization occurred even at room temperature. Strain rate was considered to be a factor only in that it produced a temperature change. There are many differences between the two experiments, the most important possibly being that their material was roughly ten times the purity of the present material. The present experiment, however, may not be as definitive as desired in this respect since it has been shown that compressed aluminum recrystallizes without a texture change6 although supplementary

Jan 1, 1961