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By A. S. Yue

Small-angle tilt boundaries have been observed in a lamellar eutectic structure. These boundaries are morphologically similar to a theoretical till boundary originally proposed by Burgers. The production of a small-angle tilt boundary in the lamellar eutectic structure is attributed to the formation of an indentation on a planar solid-liquid interface during growth. Constitutional supercooling due to the accumulation of impurity atoms ahead of the interface is primarily responsible for the generation of an indentation. The distance between two faults of like sign, designated as lamellar fault distance, is calculated by the Burgers equation for small-angle boundaries. It was found that the lamellar fault distance decreases exponentially and the tilt angle increases exponentially with freezing rate. SMALL-ANGLE tilt boundaries have been observed in many crystals in the as-grown states by etching techniques. However, a detailed substantial proof for revealing the presence of dislocations in the form of etch-pits was recently accomplished by Vogel. These boundaries are described as an array of dislocations of like sign. The angle of tilt associated with a boundary is treated theoretically by Burgers3 and is given by 0 = b/D for small tilt-angle where b is the Burgers vector of the dislocations and D is their average spacing. During the course of experimentation on the solidi fication of the lamellar Mg-32 wt pct Al eutectic a microstructure which was strikingly similar to a dislocation tilt boundary was observed. An example of the type of microstructure is given in Fig. 1, which shows a combination of transverse and top sections of the lamellar eutectic structure. The lamellae of the top section on one side of the apparent tilt angle are not parallel with those on the other side. Consequently, some lamellae must be terminated somewhere. Fig. 1 shows that the termination of lamellae is analogous to a dislocation tilt boundary in a crystal lattice originally proposed by Burgers. This work represents a closer look at such structures. The factors governing the production of the small-angle tilt boundary are discussed in terms of the freezing condition and impurity distribution in the melt. The effect of freezing rate on the distance between two faults, designated as lamellar fault distance, and the tilt angle is also described. In addition, the lamellar fault distance is determined by means of the Burgers equation for small-angle tilt boundary. EXPERIMENTAL PROCEDURE The slightly impure Mg-32 wt pct Al eutectic rods, 5/8 in. in diam and 12 in. long, were made from 99.99 wt pct Al and 99.98 wt pct Mg and were cast in a high-purity graphite mold. Some of these rods were purified by zone-refining the eutectic using direct-radiation helical-coil furnaces5 A molten zone about 2 in. long was passed through the cast rods at a speed of 5.3 x 104 cm per sec for ten passes of the molten zone. Before the freezing experiment was begun, 2-1/2 in. of the material from the head and tail ends of the zone-refined eutectic were removed from the rod to eliminate any impurity. The freezing experiments were carried out under controlled conditions,8 using a protective atmosphere of purified argon. At the end of each run, observation of the tilt boundary was made throughout the entire length of the rod with a light microscope on the top, longitudinal, and transverse sections (three mutually perpendicular sections) of the lamellar eutectic solid. The experimental data used for the determination of interlamellar spacing and tilt angle were taken from the middle portion of

Jan 1, 1964

By J. Szekely, G. P. Martins

A formulation is given and a general analytical solution is presented for the problem of spherical phase growth in multicomponent systems. The results are valid for describing the rate of bubble growth in the absence of translational motion and, therefore, should be relevant to heterogeneous nucleation processes. Typical worked examples are given, which include bubble growth due to the simultaneous diffusion of hydrogen and nitrogen in molten steel, and the growth of CO bubbles in steel, due to the reaction: C + 0 = CO. The calculations predict that the presence of additional components would enhance the growth rate and also increase the absolute rate at which the individual components are transferred into the bubble. These findings should be relevant to degassing processes. THE problem of spherical bubble growth in the absence of translational motion is appropriate to the description of the heterogeneous nucleation and subsequent growth of gas bubbles while attached to a solid surface. Such problems are of considerable practical relevance to various vacuum degassing and steelmaking processes where bubble nucleation and subsequent growth play an important part in the overall reaction mechanism. A definitive formulation of spherical phase growth has been given in the now classical paper by Scriven1 and the analogous problem for solid phase growth has been discussed by Horvay,2 furthermore a thorough review of the relevant literature has also become available.3 All this previous work was concerned with one component or binary systems. The purpose of this paper is to show that the technique developed by Scriven may be generalized to alow the consideration of multicomponent systems. The results of this generalization would be particularly relevant, as gas bubble nucleation in the metallurgical systems previously mentioned frequently involves the simultaneous diffusion of several components. In the following a general formulation is given and typical worked examples are presented. FORMULATION Consider the growth of a spherical gas bubble due to diffusion in a quiescent liquid of infinite extent, as sketched in Fig. 1. The formulation of such problems requires the statement of the following conservation equations relating to the liquid phase: Equation of continuity Equation of motion Conservation of the diffusing species—Diffusion Equation. The following assumptions are made: a) The fluid is incompressible, quiescent, and of infinite extent b) The bubble is stationary, i.e., it undergoes no translational motion c) Spherical symmetry is observed throughout d) The mass diffusivity is constant e) Chemical equilibrium exists at the surface of the bubble, and any reactions occurring there are fast compared to the diffusion process f) Inertia and surface tension effects are neglected. Equations derived on the basis of these assumptions by previous investigators were found to provide a reasonable representation of bubble growth while attached to a surface in organic and aqueous systems4,5 For gas bubble growth in molten metal systems assumptions a) to d) chould hold reasonably well. The exclusion of chemical kinetics as a possible rate limiting step may not be universally valid, but should hold for many systems operating at high temperatures. The effect of hydrodynamic inertia may be neglected for moderate growth rates; finally, the assumption neglecting the effect of interfacial tension on the bubble growth was made for the sake of mathematical convenience; its inclusion would have required the use of quite cumbersome numerical techniques. While surface tension effects could be appreciable on the initial values of the overall growth rate, especially at low absolute pressures and in case of small bubbles, they are not thought to modify the relative effects present in multicomponent systems, the study of which is the objective of this paper. Within the framework of the above assumptions

Jan 1, 1970

By H. E. Cline

A criterion for stability of a lamellar structure in two-phase alloys ulas derived froflr perturbation analysis of the steady-state solutions for eutectic growth. Both the mass and thermal diffusion were considered with boundary conditions for coupled lamellar growth. The concept of an effective liquidus slope was introduced. Excellent agreement was found with the experiments of Mollard and Flemings and the theory predicted a range of compositions on one side of the eulectic composition in which stable growth may occur with zero temperature gradient. The lamellar structure was found to be more stable than predicted by simple constitutional supercooling. The effect of impurity additions was also considered. DIRECTIONALLY solidified eutectic alloys may have a structure consisting of parallel plates or rods. It is important to know the range of composition and growth conditions that produce this ideal composite structure. Breakdown of the planar solid-liquid interface during growth may lead to a dendritic or cellular structure as in single-phase alloys. In the case of single-phase alloys, the concept of constitutional supercooling1 may explain breakdown, but in two-phase alloys the concept is no longer valid because the melting point is not given by the liquidus of one phase. Both phases must be considered in the mass and thermal diffusion problem. In the following theory a boundary condition at the solid-liquid interface was determined by using the steady-state equations of Jackson and ~unt' to relate the average temperature and composition of the liquid at the interface. Then a sinusoidal perturbation in the shape of the interface was introduced to test for stability as in the analysis of Mullins and ~ekerka.~ The wavelength of the important perturbation was assumed to be much larger than the lamellar spacing, h, and the perturbation was assumed not to change the lamellar spacing, A. Both the mass and thermal diffusion equation were used with averaged quantities at the boundary to determine whether the perturbation grows or shrinks in time. A stability criterion was determined by considering the case in which the perturbation remains stationary. An example of Pb-Sn alloys was used to illustrate the validity of this theory. I) THEORY A) Temperature at the Interface. The temperature of the interface in coupled two-phase growth is close to the eutectic temperature TE even for alloys off eutectic composition. It is assumed that the interface is in local equilibrium and one expects the temperature at the intersection of the liquid phase and the solid phases a and 0 to be near the eutectic temperature. The interface is undercooled because of curvature and the deviation of the composition in the liquid from the eutectic composition. The average interface temperature Ti for steady-state growth of a lamellar eutectic with a plane interface was determined by Jackson and Hunt' as given in their Eq. [17c]. Instead of using the ratio of the lamellar widths hp/h, as a parameter, the volume fraction of the 0 phase, +, will be used as it simplifies the algebra. The average interface temperature using this notation becomes: The average temperature of the interface depends on the volume fraction of the 0 phase, +, which may be related to the composition of the solid or the bulk composition of the liquid Co by: where C, and Cp are the average compositions of the solid phases. In steady-state growth, the composition of the solid is equal to the composition in the liquid far from the interface. The variation of the average interface temperature with composition is illustrated in Fig. 1 for the case of Pb-Sn alloys with a 10"4 cm lamellar spacing as calculated from Eqs. [I] and [2]. The interface temperature goes through a maximum at a composition near but not exactly at the eutectic composition. There is no reason why the composition corresponding to the maximum temperature should be equal to the eutectic

Jan 1, 1969

By F. G. Tahmoush, N. J. Grant, E. P. Abrahamson

The effect of grain size on the brittle -ductile transition temperature of pure iron and three dilute Fe-W alloys has been investigated by slow bend tests. The brittle-ductile transition temperature for pure iron increased with increasing grain sizes from 8000 to 500 grains per sq mm where the transition temperature then decreased with further increase in grain size. The dilute additions of tungsten lowered the transition temperature for fine grain sizes, suppressed a reversal of the curve in the coarse grained portion, and raised the transition temperature for this region. The frequency of twin occurrence increased with increasing grain size and decreasing temperature. The object of the present investigation was to re-examine the effect of grain size on the brittle -ductile transition of pure iron and several dilute Fe-W alloys. Current literature normally presents a simple relationship between grain size and the brittle-ductile transition temperature, namely, that the transition temperature increases with increasing grain size. Heslop and etch' have studied the transition temperature of mild steel by V-notch Charpy impact tests for the grain size range between 200 and 6000 grains per sq mm. The result was a linear plot of In 1-1/2 vs brittle-ductile transition temperature, where 1 is the grain diameter. Recent experiments by sipos2 revealed the possibility of a reversal in the slope of the grain size vs transition temperature plot for pure iron; however, his investigation was limited in scope, and the data obtained were insufficient to permit a definitive curve. The present investigation was undertaken in an effort to confirm or extend the grain size effect observed in iron by Sipos. PROCEDURE All alloys were vacuum induction melted, using 99.95 pct Fe containing 0.03 pct 0 and 0.004 pct C. Four ingots were made, ranging from pure iron to iron with 0.035 pct W. The chemical composition of each alloy, after processing, is given in Table I. The four alloys were prepared as 45 lb heats and cast into 4 in. diam by 16 in. molds. They were then extruded into 3/4 in. rods at 1800oF. Specimens for chemical analysis, taken from various positions along the rods, revealed no variation in tungsten content along the length. The extruded rods were reduced to 0.156 in. diam stock by a series of cold swaging operations involving intermediate vacuum annealing treatments. This procedure yielded a final cold reduction in area ranging from 10 to 90 pct. The swaged rods were chemically cleaned, cold pressed to a cross section of 0.10 by 0.18 by 1.0 in. long (2.54 by 4.73 by 25 mm). The specimens were wrapped in iron foil and vacuum sealed (10-4 mm Hg) in quartz. Various heat treatments at temperatures from 700" to 1300oC were utilized to produce grain sizes from 8000 to 1 grain per sq mm. Grain sizes were determined by measuring the average area of grains in a polished section. In specimens with grain size greater than 20 grains per sq mm, the total number of grains in the cross section was counted and then divided by the area to yield grains per sq mm. Grain size was measured both for longitudinal and transverse cross sections. The average of the two measurements was plotted vs transition temperature. An Instron Table Model testing machine was used to determine the transition temperature of specimens tested in three point horizontal bending. The specimen was supported by two 1/4 in. dowels one in. apart. Bending was accomplished with a push rod, the end of which had an 1/8 in. radius. The head speed was 50 in. per min. Temperatures were established through the use

Jan 1, 1963

By P. M. Blair

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

Jan 1, 1965

By R. D. Pehlke, W. M. Small

The effect of added alloying elements on the solubility of nitrogen in a liquid alloy of 74 wt pct Fe. 18 wt pct Cr, and 8 wt pct Ni has been studied. At 1600°C and 1 atm nitrogen pressure, aluminum, chromium, columbium, manganese, molybdenum, tantalum, and vanadium increase the solubility of nitrogen in the liquid 18 Cr-8 Ni alloy, whereas additions of cobalt, copper, nickel, silicon, tin, and tungsten decrease the solubility. Interaction parameters based on a constant ratio of iron, chromium , and nickel have been used to quantitatively define the changes in nitrogen solubility. INVESTIGATIONS on the solubility of nitrogen in liquid iron have usually treated the pure metal, and often the influence of a single added element. Pehlke and Elliott1 carried out a comprehensive study on the influence of alloying elements on nitrogen solubility in liquid iron, and compared their results with previous studies. Turnock and pehlke2 have recently reported a study on the solubility of nitrogen in multicomponent liquid iron alloys. The original interaction parameter proposed by wagner3 was successfully employed in these studies. The need for obtaining information on the gas solubility in more complex alloys is indicated by the increasingly significant role played by dissolved gases in the processing of iron-based materials. The commercial stainless steels, for example, represent an area where data on nitrogen solubilities are required as a function of composition, temperature, and gas pressure. The nitrogen solubility in liquid Fe-Cr-Ni alloys has been studied by Humbert and Elliott.4 Their investigation covered the entire ternary system, and attempted to use data from binary alloys to predict nitrogen contents in liquid ternary alloys. Langenberg,5 using a method developed by Morris and buehl,6 calculated the nitrogen solubilities in multicomponent iron alloys from data available on binary iron alloys at 1600°C. Nelson,7 using a regular-solution model, extended this prediction to include the effect of temperature. Chipman and corrigan8 and Turnock and pehlke2 also employed empirical methods to calculate nitrogen solubilities in liquid multicomponent iron alloys. In order to provide a basis for predicting nitrogen solubility in liquid stainless steels, and to extend the data available for thermochemically modeling systems of this type, the following investigation was carried out. The study involved measupements of nitrogen solubility in a nondilute alloy and examination of the influence of a small addition of an alloying element on the nitrogen solubility. The study was focused on a very narrow range Of composition in the vicinity of 18 wt pct Cr-8 wt pct Ni (herein referred to as 18-8) with special emphasis on the change of.nitrogen solubility with the addition of a fourth metallic element. EXPERIMENTAL PROCEDURE A Sieverts' apparatus, described previously,1,2 was used in this investigation. The charge materials were contained in a recrystallized alumina crucible and inductively melted. Additions were dropped from a glass side-arm. All materials were hydrogen-refined and vacuum-treated to eliminate or minimize the oxygen content. The source and purity of materials is given in Table I. The initial metallic charge was heated to approximately 1000°C and held under vacuum for 30 to 45 min as a means for removing hydrogen and adsorbed gases. Argon was used as the calibrating gas to determine the "hot volume". The average hot volume of the system was approximately 50 cu cm STP and a few experimental runs were made using a reaction bulb with a hot volume of 80 cu cm. The temperature range studied was from 1550° to 1700°C and the temperature was measured optically. THEORETICAL PRINCIPLES General Discussion. Nitrogen dissolves in liquid iron alloys according to the reaction: 1/2N2(g) = N (wt pct in solution) [l] The equilibrium constant, K1, for this reaction is written as:

Jan 1, 1969

By H. M. Otte, A. G. Crocker

SLIP, twinning, stacking faults, and precipitates on well-defined planes in a crystal produce traces that are visible on either a polished or an etched surface. The purpose of this note is to establish the usefulness and limitations of the information obtained from such traces when determining the orientation of single crystals or grains in a polycrystal. Attention will be confined specifically to traces produced by octahedral planes in cubic crystals, though most of the remarks can be taken as more generally valid. The problem of determining the crystal orientation is completely solved in principle' and in particular for the special case under consideration.1'2 Tables are available3 in which the orientations can simply be looked up when given the two or three angles between the traces. Thus a quick and simple method is available for obtaining the orientation of a crystal from surface traces. In this note two examples will be given from recently published papers and will show the use of the method with the aid of the Tables, as well as some of the precautions that should be taken in the procedure. As part of an analysis of X-ray Kossel lines, Peters and Ogilvie 4 determined from them the crystallographic orientation of the specimen. They chose a Cu-5.10 wt pct Si alloy that had been partly transformed from the a (fee) phase to the k (hcp) phase. The latter formed as very straight platelets along the octahedral planes of the a phase, and a photograph (Fig. 7 of Ref. 4) was reproduced in the paper4 of the region for which the orientation had been determined from the Kossel lines. Apart from a very small area (about 0.15 mm diam) near the top center of the photograph, the region shown was clearly all of one orientation. Traces in four different directions were readily identified. Starting with the smallest angle and proceeding counterclockwise, the angles between the traces were measured to be 27° 43', 55° 16', 64 3 43 1, and32"26'. These add up to 1802 08 1, indicating a cumulative error of 8'. The correct values should, of course, be taken from the original glass plates, or at least from prints on distortion-free paper, so the precision of the quoted values is actually greater than their accuracy. The solutions can be looked up in the Tables3 in four ways, depending on how the angles are paired, tion for the minutes of arc is required and the solution obtained will be identical for all the four pairs if the values of the angles are not in error. In a practical situation, such as this one, there are several sources of error, so that the four solutions obtained will all be slightly different. An arithmetic average of the four solutions is mathematically incorrect, but may, in special cases, give an answer close to the correct one. A rigorous procedure for refining the orientation has been described by Mackenzie;5 he gives a least-squares method that can be applied, if necessary, to this type of problem. The orientation thus determined can be compared with the one from the Kosel lines and discrepancies larger than the experimental error assigned to nonor thogonality of the X-ray or electron beam to the specimen. Here the example will only be used to illustrate the general procedure for making a very rapid rough check. From measurements on the stereographic projection, Fig. 9 of the paper,4 we can determine that the surface normal is approximately (0.434, 0.704, 0.562). From the Tables3 we find, using 28°, 55" as a good approximation to the first pair (27°43', 55°16'), that the surface normal is (0.453, 0.678, 0.579).* These two values are within 2 deg of each

Jan 1, 1967

By R. E. Collins, Paul B. Crawford

Potentiometric model studies have been made of ertically fractured reservoirs. All fractures originated at he well and extended into the reservoir for various distances. Studies were made to determine the areal sweep efficiencies for line-drive patterns. The areal sweep efficiency for the unfractured system studied war 70.6 per cent. It was found that if the vertical fracture war paralIel with the direction of flood, the areal sweep efficiency may be considerably reduced, approaching zero in some instances. If the fracture was at right angles to the direction of flood, the areal sweep efficiency may be greater than that for the unfractured system. The sweep efficiency depends on the length and orientation of the fracture and direction of the flood. It is concluded that considerable effort should be made to determine the nature and direction of fractures before initiating a flooding program. Such information is required if the maximum sweep efficiency is to be obtained and channeling is to be reduced to a minimum. INTRODUCTION The recent development and application of fracturing techniques to petroleum reservoirs has served to focus considerable attention on the effect such fractures may have on secondary oil recoveries. The opinion has been expressed that fracturing may not alter appreciably the water flooding possibilities. Others have thought that channeling may occur and that the creation of fractures may serve to substantially reduce the possibilities of conducting a successful water flood in the reservoir The purpose of this study was to obtain quantitative estimates of the effect of vertical fractures on the areal sweep efficiencies of line-drive flooding patterns. In presenting this work on vertical fractures, it is not to be inferred that the present fracturing techniques result in producing a vertical fracture. The types of fracture which are created are not known with certainty in every case. Some are believed to be horizontal, others inclined along bedding planes, and some data indicate vertical fracturing. Each fracture will probably exert a different influence on the sweep efficiency of the secondary recovery program; consequently each fractured well should perhaps be considered as a separate problem, requiring a careful analysis. Preliminary work on the effect of fractures on producing capacity of a well indicated that for practical purposes many fractures may be treated as having a nearly infinite permeability This assumption appears to be valid in some instances where the permeability of the fracture is only 10 or 100 times greater than the permeability of the matrix. It is known to depend on the type, orientation and length of the fracture. In realization that fractures of only 0.01 in. width have calculated permeabilities near 5,000 darcys, it is believed that many fractures may be considered as of near infinite permeability without incurring serious error. Primarily for this reason this study was made with fracture analogs of very great permeability compared to the matrix. FRACTURED SYSTEM ANALOG The pattern sweep efficiencies reported here were obtained with the aid of a potentiometric model. The application of this type model for solving reservoir problems has been described in the literature.1,2,3 The electrolyte reservoir was 20 in. wide by 30 in. long by approximately % in. deep. The liquid depth was constant, which corresponded to a constant permeability-feet product from the injection to the production well. The electrolyte reservoir was built to correspond to a unit in the line-drive pattern shown in Fig. 1. This figure shows part of an infinite array of wells having alternate rows of fractured and unfractured wells. The area within the dashed rectangle shows the area generally studied in the model. In using the model to represent a fractured system, a thin strip of copper was soldered to the well or wells and the length and orientation of the strip was made to conform to the vertical fracture under consideration.

Jan 1, 1955

By H. Kurushima, S. Tsunoda

An investigation of the hydrometallurgy of Cu-Zn concentrates by The Dowa Mining Co. of Japan is described and a detailed flowsheet of the Kosaka hydro-metallurgical plant is given. Sufficient operating experience has shown that the Dorrco FluoSolids system has proven successful for recovery of both zinc and copper by roasting in a fluidized bed. At present, 2400 metric tons of concentrates per month are being treated and recoveries of 93 pct Cu and 65 pct Zn are being effected. A SURVEY of the Japanese chemical industry in 1949 revealed, among other things, that sulphur-bearing blast furnace gas from the Kosaka smelter of The Dowa Mining Co. Ltd., Kosaka, Japan, was being exhausted to the atmosphere as waste. In addition, it was costing that company an average of ten million yen ($28,000) annually for damages to the surrounding property caused by the sulphur dioxide in the gas. It was recommended by this survey that the gas be utilized for the production of sulphuric acid by the contact process. Mr. Kurushima visited the United States in May 1950 to investigate the engineering, construction, and operation of a contact acid plant. Any acid produced at Kosaka would necessarily have to be 98 pct H,SO, to prevent corrosion in railroad tank cars transporting it to the point of use. However, the production of 98 pct acid requires constant volume and SO, content of the gas. As the converter gas at Kosaka fluctuates in both volume and quality, it was decided to evaluate the smelting process with relation to other methods of recovery. This evaluation revealed a relatively new type of roaster, the Dorrco FluoSolids reactor, which was reported to be capable of roasting the copper sulphide concentrate in an excess of air with the production of an SO1 gas strong enough for acid production and a water-soluble copper sulphate calcine. The copper sulphate then could be leached from the roasted ore and electrolytic copper recovered from the leach solution in much the same manner as practiced by the Chile Exploration CO. for treating naturally occuring copper sulphate ores. Use of a reactor of this type would eliminate the need for a blast furnace and converter and, in addition, would effect savings in labor and fuel. The Kosaka ore is a microscopically fine mixture of copper, zinc, and lead sulphides. The lead and a small amount of the barite present can be separated by flotation, but any separation of copper and zinc proved extremely difficult. In blast furnace operation at Kosaka, zinc had been slagged off until a stockpile of about 12,000,000 tons containing 5 pct Zn had been accumulated. Mr. Kurushima felt that if any of the zinc in the concentrate was sulphat-ized and leached with the copper, it also could be recovered electrolytically. The Dorr Co., sponsors of fluidization technique in the metallurgical field, suggested that sulphate roasting be tested out on a laboratory scale and that, if successful, a commercial Dorrco FluoSolids system be constructed. They indicated that the results from such laboratory tests could be transposed to a commercial scale without pilot planting the roasting operation. It was therefore decided to send samples of flotation concentrate to the West-port Mill, laboratories of The Dorr Co., for testing. The objective of Mr. Kurushima's visit to the United States thus had changed from investigation of sulphuric acid plants to a fundamental examination of the hydrometallurgy of copper and zinc. Testing Program In February 1951, Mr. Tsunoda went to the United States to witness the application of the FluoSolids system to sulphate roasting at The Dorr

Jan 1, 1956

By P. S. Kotval, H. Hahell

PRECIPITATION of M23C6 carbides at grain boundaries in austenitic steels and nickel base superalloys has received considerable attention. Hatwell and Berghezan,1 in a study using carbon extraction replicas, observed that dendritic growth of the M23C6 carbide phase occurred at grain boundaries in austenitic steels. Merrick and Nicholson2 have observed that in thin foils of "Nimonic 75" alloy the precipitation of M23C6 carbide was discontinuous at grain boundaries. Lewis and Hattersley3 have suggested that the type of M23C6 precipitation observed in Ni-Cr austenites represented the early stages of discontinuous precipitation. In a more recent study, Singhal and Martin4 have concluded that grain boundary M23C6 precipitation involves the migration of an austenite grain boundary by a mechanism very similar to that proposed by Tu and Turnbull.5 Tu and Turnbull's suggestion essentially involves grain boundary migration under the driving force arising from the difference between the surface free energies of the opposite sides of the plate. Thus, Singhal and Martin's observation4 implies that the first stage of discontinuous precipitation has taken place and that the grain boundary has provided a diffusion short circuit for the sub-stitutional solute atoms (principally chromium) which assist the preferential growth of the M23C6 plate in contact with the grain boundary. However, no evidence of the characteristic cellular morphology typical of discontinuous precipitation was found in Singhal and Martin's studies. Presumably, this is due to the low volume fraction (i.e., low supersaturation of solute) of the precipitating carbide reported by Singhal and Martin.4 Recently, Raymond6 has studied precipitation of the M23C6 carbide in Incoloy alloy 825 and shown that the growth of this phase is perpendicular to grain boundaries at 1400°F and in the plane of the grain boundary at 1700°F. The former aging condition resulting in discontinuous or "cellular" precipitation. Raymond has suggested an explanation for this behavior based on the relative rates of chromium diffusion by volume and grain boundary diffusion at the two different aging temperatures. In the course of a recent investigation on precipitation behavior in several ' strengthened superalloys, precipitation of M23C6 carbide at grain boundaries was studied. Under conditions of a high degree of supersaturation, the M23C6 carbide precipitates in the discontinuous or cellular morphology in Waspalloy (Ni-18.51 Cr-2.0 Fe-3.53 Mo-1.16 Al-2.78 Ti-0.09 C). This note presents electron microscopy observations of this mode of precipitation and its effect on mechanical properties. Specimens of Waspalloy after "solution" heat treatment at 1950°F show essentially single phase matrix with a few residual particles of "primary" MC carbide. In normal commercial heat treatments, precipitation of Ni3(Al,Ti) ordered fcc r' phase occurs throughout the matrix and M23C6 precipitates as globular particles at the grain boundaries. In these commercial heat treatments the degree of supersaturation during precipitation is low. However, if specimens of Waspalloy are quenched from the solution heat treatment temperature of 1950°F and then aged at an aging temperature where the degree of supersaturation is high, discontinuous precipitation of M23C6 occurs. Figs. 1 and 2 show thin foil micrographs of Waspalloy, solution heat treated at 1950°F, water quenched, and aged at 1292° and 1598°F, respectively, for 500 Fig. 1—Thin foil of Waspalloy, solution heat treated at 1950°F for 1 hr, water quenched, and aged for 500 hr at 1292°F. Magnification 15,500 times. Fig. 2—Thin foil of Waspalloy, solution heat treated at 1950°F for 1 hr, water quenched, and aged for 114 hr at 1598° F. Magnification 13,000 times.

Jan 1, 1970

By H. E. Wilborg, G. V. Henderson

This chapter deals with those sands employed by foundries for the manufacture of cores and molds used in the casting of such common metals as steel, gray iron, ductile iron, aluminum-based alloys, and copper-based alloys. Most metal is cast in "green sand" molds in which the sand mixture, consisting primarily of silica, clay, and water, forms a plastic material. This plastic material is rammed around a pattern in a split box called a flask, then the box is separated, the pattern is withdrawn, and the metal is poured. These green sand molds are broken apart when the metal has solidified. The sand is then remixed and a new mold made for the next casting. To a lesser extent silica sand is mixed with resins or oils and, after forming, the mixture is hardened by either heat or chemical reaction. In any case, the foundry mold or core is used just once before being broken up and remixed for manufacturing another core or mold. By far the greatest consumption in the foundry is of those types of sand high in silica, low in clay (under 0.5%), and occurring in numerous locations thoughout the United States. Consumption End Uses Foundry sands fall into four distinct categories: silica sand, lake sand, bank sand, and natural molding sand. Silica sand is a general term used to describe washed, graded, dried, and cooled clay-free sands which are found in many parts of the United States. They generally contain 90 to 100% silica but vary greatly in chemical analysis and grain structure. Lake sand is the sand dredged from Lake Michigan or dug from the dunes bordering that lake. Bank sand is composed primarily of silica grain but containing up to 5% clay. It is shipped in both the damp and dry state but the normal processing of bank sand consists only of drying and scalping. Occasionally two different grades of bank sand are blended prior to or during loading. Natural molding sand is a term used to describe a sand containing sufficient clay as mined to be suitable for molding purposes. Clay content may vary from 10 to 24% with the balance of the material consisting primarily of silica sand. These sands range in fineness from one peaked on the 20 mesh to one which contains more than 50% of the particles passing the 270 mesh. Lake sands are relatively coarse and are used in great quantities for the production of automotive castings. Bank sands and natural molding sands are gradually phasing out as the production of castings becomes more a science and less an art. The vast majority of foundry sands used today are those known as the silica sands and lake sands. Foundry sands are used in two basic ways. The majority are used in the green or uncured state with the balance used in the cured or set state. Green sand is grain bonded with clay and contains varying amounts of water to plasticize the clay in order to bond the grains together. Green sand is either naturally bonded or "synthetic" sand. The naturally bonded molding sands are generally found in the Hudson River Valley of New York State, Cumberland County in New Jersey, and in Tennessee and Mississippi. Synthetic sand is molding sand compounded in the individual foundry. It is made up primarily of graded silica sand, western or southern bentonite, water, and carbonaceous and/or cellulosic materials. Green sand is used primarily in making molds. The use of naturally bonded green sand has been declining for many years, primarily because naturally bonded sands are more difficult for the foundryman to find and control. Synthetic green sand, on the other hand, can be

Jan 1, 1983

By R. I. Jaff, H. R. Ogden, D. J. Maykuth

A phase diagram for alloys containing from 0 to 66.9 pct Mn was determined. Two compounds, tentatively labeled 6 and y, were found in this range. The 6 compound is located at about 66.9 pct Mn and melts congruently at 1330°C. The ? compound has a composition near 53 pct Mn and originates in a peritectic reaction at 1200°C between liquid and 6. ß and Y enter into a eutectic reaction at 1175°C. The ß phase field terminates in a very sluggish eutectoid reaction between a and ? at 550°C. MANGANESE is one of the important alloying addition elements to titanium. As a ß-stabiliz-ing addition for a-ß and ß alloys, it provides high strength with adequate ductility.' Two of the earliest commercial alloys contain manganese as a major alloying component.' The earliest published data on the constitution of Ti-Mn alloys was that of Wallbaum,~ who reported the existence of the compound TiMn2. The transformation range of commercial-purity Ti-Mn alloys was investigated by Craighead, Simmons, and Eastwood' who found limited a solubility and a lowering of the ß transus to about 790°C at 7 pct Mn. Preliminary work by the authors of this paper, using high-purity titanium, confirmed the ß-stabilizing action of manganese and showed the existence of several intermediate phases in the Ti-Mn system." More recently, McQuillan3 determined the ß transus temperatures for high-purity Ti-base alloys containing up through 4.4 pct Mn. These data are given in Fig. 1. The present investigation covered alloys made with a high-purity base. The constitution of titanium-rich alloys made with commercial titanium has been reported4," by Holladay, Kura, and Jackson. The alloy range covered in the present investigation was from 0 to 69.5 pct Mn. This alloy range extends slightly beyond the composition of the first congruent-melting compound encountered on proceeding from the titanium end of the system. A description of the materials and methods used in the determination of the diagram is given in the Appendix. The partial phase diagram determined for the Ti-Mn system is shown in Fig. 2. Details on equilibria in the titanium-rich alloys are given in Fig. 3. Liquid-Solid Relationships Observations on the cast structures of the arc-melted ingots showed that the terminal solid solution of manganese in ß titanium enters into a eutectic reaction with an intermetallic compound labeled ?. As will be shown later, this compound has a composition around 53 pct Mn. The eutectic reaction temperature was determined as 1175°C by thermal analysis which also showed that the ? phase originates in a peritectic reaction at 1200°C. Fig. 4 shows typical eutectic structures for alloy compositions on both sides of the eutectic composition which is located at 42.5 pct Mn. In alloys over the range of 45.4 through 69.5 pct Mn, a second intermetallic compound, labeled 6, was identifield as the primary phase. This phase consistently increased in amount as the manganese content approached 66.9 pct, and, as shown in Fig. 5, an alloy structure containing 66.9 pct Mn consisted essentially of the 6 compound. Thermal-analysis data indicate a melting point of about 1330°C for the 6 intermediate phase. At 67.0 pct Mn, small quantities of a new, higher manganese-content phase were observed in the form of Chinese-script eutectic particles. Comparison of Fig. 6a and b shows an increasing quantity of this phase as the composition was increased to 69.5 pct Mn. The formation of peritectic ? in arc-melted alloys containing 45.4 pct Mn and above appears to be almost entirely suppressed, as shown in Fig. 7a and b. This is because of the narrow freezing range available for its formation by peritectic reaction and because of the very high cooling rates obtained in

Jan 1, 1954

By J. L. Wyatt

THE smelting of ilmenite to produce a slag rich in titanium, with pig-iron as a byproduct, introduces new concepts in electric smelting metallurgy. Titanium slags are characterized by low electrical resistance, thereby requiring a different type of electric furnace operation, and power supply sources substantially different from those encountered in normal steel melting and siliceous ore smelting furnaces. This paper describes special apparatus for determining specific electrical resistivity of titanium slags, and reports the values obtained on slags encountered in the process of smelting ilmenite ores. In an induction furnace under conditions approximating normal furnacing operation, 25 general heats and 6 special standardization and comparison heats were run. Results indicate relatively low resistivity. Resistivity depends primarily on slag composition, and, to a lesser degree, on temperature. During the summer of 1947 a cooperative project between the National Lead Co. and the Bureau of Mines was conducted at Boulder City, Nevada, in the Electrometallurgical Laboratories of the Bureau. The project involved smelting of ilmenite ores in electric furnaces, and has been reported in detail.' The author participated as a member of the staff of the National Lead Co. during the entire campaign. One outstanding characteristic of the titaniferous slags, their apparently excellent and variable electrical conductivity, was considered sufficiently important to warrant further study. Although little electric smelting of iron ores has been done in the United States, a certain amount of information is available for comparative purposes. Personal experience in smelting siliceous iron ores in electric furnaces furnished further information for comparison. In general, silicate slags are relatively poor conductors of electricity. Current flow during smelting operations may be maintained through a floating layer of coke on top of the slag bath, or through a bath of molten iron, beneath the slag. Silicate slags are not sufficiently conductive to furnish a path for current passage when normal furnace voltages are used. Experimental smelting of a silicate iron ore was undertaken for comparative purposes. A fluid melt was not obtained until the FeO content was reduced to 15 pct. The resistivity increased as the FeO content decreased, FeO compounds being the most conductive of the various oxides in such a normal slag system. The converse appears to be true in slags containing large percentages of TiO2. It has been found that titanium slags become molten with FeO percentages in excess of 30 pct, and that the conductivity becomes generally better as the FeO con-tent of the slag diminishes. This phenomenon was demonstrated in the furnacing operations at Boulder City. During the initial smelting period furnace electrodes could be held immersed several inches in the slag, with tap selectors set for a no-load voltage of 70 v. As the FeO content decreased, the furnace electrodes would "walk up," until at one point in the process the tips would break the surface and arc on top of the slag. This was undesirable since direct arcing caused losses of important flux additions. The slag compositions under study at that time were such that losses due to surface arcing sometimes caused viscous slags to be formed. Forcing the electrodes into the slag bath by manual control to prevent surface arcing resulted in excessive power input and lack of control. Suitable low voltages were not available. Measurement of the electrical resistivity of titanium slags at various operating temperatures and slag compositions was undertaken. It was believed that this information would be of use in the design of electric furnaces suitable for smelting ilmenite, with voltages, electrode sizes, and electrode spacing adapted specifically to titanium slags.

Jan 1, 1951

By D. Harker, W. E. Seymour

CERTAIN alloys of iron and nickel, when rolled and annealed, possess a preferred crystal orientation: (001) in the rolling plane and [loo] in the rolling direction, when recrystallized at 850" to 1050°C after approximately 98 pct cold reduction. Since the preferred orientation makes a direction of easiest magnetization coincide with the rolling direction, these alloys, especially a 50 pct iron in nickel alloy, have found wide application in the electrical industry' for choke coils and special types of transformers. The rate of the recrystallization reaction was studied at temperatures ranging from 500° to 600°C. The heat of activation for the reaction was calculated from the observed rates, and crystal orientation determinations were made before and after re-crystallization. A vacuum melted iron-nickel alloy* analyzing 49.6 pct nickel, 0.018 pct carbon, and 0.30 pct manganese was used for the experiments. The alloy was cold-rolled 98 pct into strips 1/2 in. wide by 0.002 in. thick. Specimens 1 in. long were sealed under a vacuum (approximately 10-8 mm Hg) in 1/2 in. ID pyrex glass tubes. For heat treatment these tubes were fastened to Nicrohm wires and submerged in a molten salt bath controlled to ±2°C. (The maximum measurement error was within 3°C.) The time at temperature was varied logarithmically from one sample to another, and runs were made at 500°, 525", 550" and 575°C. A molten lead bath was used for a run at 600°C which was controlled to the same accuracy. To determine the extent of recrystallization after a particular time at a given temperature, a method was used suggested by the work of Decker, Asp, and Harker.3,4 This method employs an X-ray spectrometer whereby X rays are diffracted from crystallites whose diffracting atom planes are parallel to the rolling plane of the specimen. The intensities of the diffracted rays are measured by a Geiger counter. The counter can be rotated through a range of angle, and can be motor driven through this range at a speed of 2" per min. The measured intensities are plotted vs. angle by a potentiometer recorder. In a 50 pct iron-nickel alloy it was found that only (200) reflections were obtained from the rolling plane with recrystallized material, and that only (220) reflections were obtained from the rolling plane with unrecrystallized material. No other reflections were obtained under these conditions over a range of Bragg angle of 0 to 45" using either type of specimen. As a consequence, the intensity of the (200) reflections from the rolling plane was taken to indicate the amount of recrystallized ma-terial present in a specimen, and the corresponding intensity of the (220) reflections was taken to indicate the amount of unrecrystallized material present. To insure flatness and proper alignment of the reflecting specimen in the spectrometer, the sample was laid in a slot 1/2 in. wide x 0.001 in. deep which was machined in a 3/4x2x1/2 in. steel carrier. The specimen then was taped down with cellophane tape. This steel cradle was mounted vertically in the specimen holder of the spectrometer so that the bottom edge of the carrier rested against a horizontal shoulder. Intensity data used in plotting the percent of recrystallized material vs. time at temperature were obtained by directly counting the diffracted beam of nickel filtered copper Ka radiation used throughout the work. The X-ray spectrometer was employed also in obtaining pole figure data to be used for orientation determinations. At the beginning of this investigation a transmission pole figure holder was used." As a thickness of 0.002 in. of this alloy was found to be opaque to the copper Ka radiation, it was necessary to reduce the thickness of the pole figure specimen. A cold-worked sample was etched in a 50 pct HCl solution for 5 min and was found to transmit a

Jan 1, 1951

By B. B. Wood, M. J. Rzasa, E. W. Hough

An apparatus for the determination of fluid-fluid interfacial tension by the pendant drop method has been constructed. The apparatus is refined beyond those previously described in that the samples are introduced into an evacuated windowed cell without contacting air or mercury. Interfacial tension data are reported for the water-methane system between 74 and 280°F and 15 and 15,000 psi. The variation of wettability of the stainless-steel dropper tip with pressure and temperature was observed incidentally to the measurements. INTRODUCTION The most refined previous apparatus --for the determination of fluid-fluid interfacial tension at reservoir conditions was built by Hauser and Michaels.' These authors used the pendant drop method, previously investigated by Andreas, Hauser, and Tucker.' Recently, Fordham3 has made the pendant drop method an absolute one by numerical integration of the differential equation of drop shape. It is necessary to know the difference in density of the two fluids and the acceleration due to gravity in order to evaluate the interfacial tension from measurements of drop dimensions. Contact angle is not involved in the measurements by this method. Some precise tests of the method have been made at atmospheric pressure and 77°F by Douglas," who found agreement with established values to within about 0.1 per cent for water and air, cyclohexane and air, benzene and air, and benzene and water. No values of interfacial tension from the recent pendant drop apparatusJ have been published by Hauser and Michaels. However, values for the water-benzene and water-decane systems are given in Michaels' thesis5 or Dressures from 15 .to 10,000 psi and temperatures from 73 to 268ºF. Some determinations on gas-oil and gas-water systems in the range of 15 to 4,000 psi and 78 to 178°F have been made by Hozott,' by the drop-volume method. The capillary rise method has been used to investigate the interfacial tension in gas-oil systems at 88°F to several thousand psi by Swartz.7 Other determinations in gas-oil systems at pressures under 1,000 psia have been made by Jones: and Beecher and Parkhurst.9 The scarcity of data for pressures above 5.000 psi is apparent. APPARATUS An assembly drawing of the working section of the interfacial tenqion apparatus is given in Fig. 2. 'The endanant drop (A) is housed in a pressure chamber having glass windows (B). An agitator (6) is provided. The metal parts of the apparatus coming into contact with the sample are made of AISI: Type 316 steel. Drain grooves are provided around the whdows so that the sample in contact with the O-rings is not returned to the working section when the pressure is reduced. A schematic arrangement of the apparatus is given in Fig. 3. The working section containing the pendant drop (A) is filled by samples from bombs (B) with the aid of the vacuum system (C). The O-ring sealed pistons (D) used to separate the sample from the hydraulic fluid are shown together with the pressure production and measurement system (E). A front view of the apparatus is shown in Fig. 1. Most of the hydraulic system is at the top of the panel. The pump and displacement piston are at the lower right. The Variacs for the air bath heaters. the camera. and the vacuum gauge are a1so visible on the panel. The working section together with a window closure and a drop-forming tip closure are shown in Fig. 9. The carbon arc light source with filter is visible in the upper part of the picture. Two sides of the air bath with the impeller and a part of the convection shield can also be seen. The assembled air bath is shown in Fig. 10.

Jan 1, 1951

By Gerald Sherman

MINING asbestos ore by caving at Thetford Mines, Quebec, has been described in the Transactions of the Canadian Institute of Mining and Metallurgy in papers presented by the staffs of The Asbestos Corp. Ltd., in 1934, and Johnson's Co. Ltd., in 1941. The production by caving prior to 1949, in round figures, has been 14,500,000 tons of ore by The Asbestos Corp. and 5,500,000 by Johnson's Co. They are the only caving operations in asbestos at the present time. The Jeffrey mine of the Canadian Johns-Manville Co. at Asbestos, 40 miles southwest of Thetford, is well along in preparations for caving, and The Bell Asbestos Mines Ltd. at Thetford Mines, will follow their examples later. GERALD SHERMAN, Member AIME, is Consulting Engineer, Johnson's Co. and Canadian Johns-Manville Co., New York, N. Y. AIME San Francisco Meeting, Feb. 1949, and New York Meeting, Feb. 1950. TP 2828 A. Discussion (2 copies) may be sent to Transactions AIME before May 31, 1950. Manuscript received Nov. 2, 1949. Nearly all of the asbestos produced by the western hemisphere is mined in those districts of Canada. Since it has been the subject of a number of technical papers, this contribution is offered as a review of progress. Some introductory information seems appropriate, for few members of the American Institute of Mining and Metallurgical Engineers are familiar with the mineral, or how asbestos deposits differ from others in caving characteristics. Asbestos is found in peridotite, a basic intrusive rock irregularly altered to serpentine and erratically fractured, well enough to cave in bulk, but leaving many blocks of solid ore of all sizes scattered through the deposit. Chrysotile is the variety of asbestos produced in Canada. Its fibers are the most valuable commercially, having a good proportion of long fibers, strong, silky, flexible, and suitable for spinning. The fibers occur in all lengths, from the microscopic to 3 in., or more, the shorter are more abundant. Their values depend on length combined with other qualities noted above. They occur as veinlets filling fissures in the peridotite, with crystals compactly arranged in parallel across the cracks. They are sometimes lightly frozen to the walls but the partings are distinct, often marked by particles of magnetite adhering to the fiber ends. Sometimes they break away from the rock in slabs an inch or more in thickness. In wide fissures the fibers are not always continuous, the "broken fibers" being separated by lines of parting with occasional scattered grains of magnetite, which is an injurious mineral to be elirninated when possible. In addition to the chrysotile ''Cross fiber," some "slip fiber" is found in fault fractures in limited quantities of less valuable character. With the exception of the Asbestos Corp, ore, in which the content is said to be more uniform, asbestos is eccentrically distributed through the peridotite. It is the general impression that the higher grades are more often found in hard rock. There are rich pockets and leaner areas which range down through various grades to minor barren sections. Wholesale mining is necessary to make sure of recovering all the high grade for there are no reliable indications of its existence. Ore selection only appears in constructing assay-grade limits for caving. The mill yield of the Thetford ore body in 1934 was 6 to 10 pct of the ore milled, made up of all classes, with occasional runs up to 15 pct or so. Quotations for Canadian asbestos* are as follows: Per Ton No. 1 Crude (% in. and longer) $896 @ $960 No. 2 Crude (Y8 to % in.) 350 @ 545 Spinning (less than % in.) 204 @ 378 Other classes, price diminishing with length. Shorts or "refuse" 24 @ 46 • From Eng. and Min. Jnl., Nov. 1948. Asbestos of even superior quality for electric insulation is found in Arizona limestone. It is white, containing very little combined iron, but occurs only in small quantities.

Jan 1, 1951

By Robert D. Pehlke, Weldon L. Daines

The reduction of manganese oxide from a basic slag by silicon dissolved in liquid iron at steelmaking temperatures was studied to determine the rate-controlling step for the process. The experiments were carried out under an argon atmosphere in a resistance-heated vertical-tube furnace. The two liquid phases, slag and metal, were held in zirconia crucibles. Sampling and chemical analysis of the slag phase led to time-concentration curves for the system. Results obtained by varying parameters such as temperature, stirring rate, concentration, and melt geometry defined the limiting reaction mechanism. The reduction of manganese oxide by silicon is a fast reaction and reached equilibrium within 90 min. The rate of stirring has a marked effect on the reaction rate, while, on the other hand, the effect of temperature is very small. This identifies the rate-controlling step as diffusion rather than interfaciul chemical reaction. The variation of other parameters revealed that a) diffusion of manganese in the iron phase is the slow step, b) the value of D/6 for manganese in iron at 1550°C is 7x 10-4 cm per sec, c) the energy of activation for the apparent rate constant is 7 kcal per mole, and d) increasing the amount of silicon in the iron decreases the diffusion coefficient of manganese. THE equilibrium relations pertinent to the reduction of manganese oxide from a steelmaking slag by silicon dissolved in liquid iron have been studied, but little is known about the mechanisms that control the rate of the process. The purpose of the present investigation was to study the rate of reduction by silicon and to determine the rate-limiting step. The effects of stirring rate, temperature, composition, and melt geometry were investigated as a means of identifying the controlling step. On the laboratory scale, relatively little work has been done in the area of slag-metal kinetics, primarily because of the problems confronted in attempting to contain liquid slags and metals in laboratory crucibles at steelmaking temperatures. Since graphite is an ideal material for a crucible, several investigations have been carried out on the reduction of various oxides from a slag by carbon in graphite-saturated iron and on the transfer of sulfur between a slag and graphite-saturated iron. Most of this work has been performed in induction furnaces. Studies where the container was not graphite are extremely limited. A single kinetic study of manganese transfer between molten iron and basic slags was carried out by Sano, Shiomi, and Matsushita.' These investigators prepared calcium oxide and titanium oxide crucibles to contain the two-phase system. Only a few experiments were carried out and the results had to be treated qualitatively. Despite the paucity of kinetic data, a large quantity of information is available in the literature concerning the thermodynamics2-12 and physical properties'3-21 of the slags containing manganese oxide and the dilute liquid iron alloys containing manganese and silicon which were involved in this research program. EXPERIMENTAL METHOD A vertical-tube resistance-heated furnace was used in this investigation to study the slag-metal reaction Although most previous slag-metal studies were carried out in induction furnaces, a resistance furnace has several advantages including better control of temperature and stirring rate. Hence, a resistance furnace was selected for this study, and its design is shown schematically in Fig. 1. Slip-cast, fine-grained, low-porosity, lime-stabilized cubic zirconia crucibles were used for the kinetic experiments. The crucibles were 4 1/2 in. high and 2 1/8 in. OD with a 1/16-in. wall. The crucibles were only slightly

Jan 1, 1969

By Tasuku Fuwa, John Chipman, David H. Kirkwood

Fe-C-Si alloys in silica crucibles were held at 1600°C in a controlled atmosphere of CO and Co2 and the approach to equilibrium was obsertsed. Results were not of sufficient precision to establish the activity and interaction coefficients. Observations 012 the equilibrium provided an average value for the standard free-energy change which differed by + 1.1 kcal from the earlier accepted value. The C-SL relation at other temperatures is calculated and shown graphically. IN the refining of steel in an acid-lined furnace, it is well known that, unlike in basic refining, a very significant amount of silicon may remain dissolved in the metal even when the carbon content approaches a final endpoint. Further, it is known that the residual silicon increases with increasing temperature. The quantitative relation between silicon and carbon can be calculated for equilibrium conditions from data already available on the Fe-C-O and the Fe-Si-O systems. Such equilibrium calculations could show us approximately the limitations within which the reaction would be expected to operate in practice. In the oxidation of a bath containing both carbon and silicon it is well known that at low temperatures silicon is preferentially oxidized first. It can be shown by calculation that at high temperatures carbon will be oxidized preferentially. It is the purpose of this paper to record direct observations on the Si-C relation by means of experiments carried out in an atmosphere of carbon monoxide for Fe-C-Si alloys contained in silica crucibles. When such a system is brought into equilibrium, at a fixed temperature, the composition of the liquid metal is governed by the following equations: Both K1 and K2 are known with a fair degree of accuracy, and may be calculated from the tabulated In general the percentage of CO2 is extremely small compared to that of CO and in the experiments to be described the partial pressure of the latter was generally only slightly below 1 atm. EXPERIMENTAL METHOD The induction furnace employed was the same as that described by Rist and chipman' and by Fuwa and chipman.2 It was modified slightly in view of the change from alumina to silica crucibles. The small silica crucible was surrounded by a thin susceptor of molybdenum sheet which in turn rested inside a zirconia crucible. Since the zirconia was subject to breakage, it was supported inside an alumina crucible. The charge consisted of 30 to 35 g of electrolytic iron along with the required amounts of refined silicon and graphite to produce a melt of required composition. Purified carbon monoxide and carbon dioxide were passed into the furnace through carefully calibrated flow meters. Two separate and independent series of experiments were carried out. In Series I (D.H.K.) the melt was held in the furnace at 1600°C for periods of time up to 6 hr. Samples were taken at intervals in silica tubes and were analyzed for carbon and silicon. The results of the last one or two samples taken from each experiment and the direction in which equilibrium appeared to be approached are shown in Table I. In Series II (T.F.) the silica crucibles were of less satisfactory strength and the melts were held at 1600° C for periods of only 2 to 4 hr. In heats number 29 to 45 inclusive, the metal was cooled in the furnace and the ingot itself served as a sample for analysis. In heat 46 and all subsequent heats, samples were withdrawn in silica tubes at the termination of the heat, and these samples were analyzed. The results of all of these experiments in which equilibrium appeared to be approached are listed in Table II.

Jan 1, 1965

By Pol Duwez

The effect of the rate of cooling on titanium, zirconium, and thallium has been measured. For titanium-molybdenum alloys, it has been shown that for molybdenum concentration up to 8 pct the rate of cooling has no effect on the transformation temperature, but this temperature decreases progressively with increasing molybdenum content. HIGH strength to weight ratio and remarkable corrosion resistance have often been advanced as the main reasons for the present widespread interest in titanium. For the physical metallurgist, the existence of an allotropic transformation in titanium is another important factor. As in the case of iron, an allotropic transformation may be the fundamental property upon which the response of Ti-base alloys to heat treatment will depend. It is therefore important to study the mechanism of the allotropic transformation in pure titanium and the effect of alloying elements on this transformation. The existence of an allotropic transformation in titanium was reported in 1931 by Schulze.' Later, De Boer et al.' determined the transformation temperature from a discontinuity in the electrical resistivity at 882"42O°C. According to recent work," the transformation temperature is 882.5"&1°C. The crystal structure of the room temperature a phase is hexagonal close-packed and the high temperature fi phase is body-centered cubic.' It is well known that the high temperature /3 phase of pure titanium cannot be retained by quenching to room temperature. Several investigators'-' have shown that in alloys of titanium with metals which are soluble in the @ phase (for example, nickel, molybdenum, chromium, or manganese), the /3 structure can be retained, providing the concentration of the alloying element is greater than a certain critical value. Below this critical value, the alloys transform during quenching, at least partially, into a supersaturated solid solution having the a. structure. This behavior is very similar to that found in some iron binary alloys (Fe-Ni and Fe-Mh, for example) and is an indication that the transformation might occur by means of a shear mechanism analogous to that observed in the austenite-martensite reaction. The present investigation is mostly concerned with the effect of the rate of cooling on the /3 to a transformation. In the first portion of this paper, the results obtained on pure titanium are described. Two other pure metals, zirconium and thallium, were also investigated, because they both have an allotropic transformation like titanium, from an hexagonal close-packed to a body-centered cubic structure. For purposes of comparison, the effect of rate of cooling of the y to a transformation of iron was included in the experimental program. The second part of this paper deals with Ti-Mo alloys. The equilibrium phase boundaries involving the solid state reactions on the Ti-rich side of the system were determined and a study was made of the effect of the rate of cooling on the transformation of the high temperature /3 solid solution. Experimental Procedure Since the main object of the present investigation was the determination of the effect of the rate of cooling on a transformation temperature, the best technique available was that used by Greninger" in his study of the martensite arrest in carbon steels. This method consists of heating a very small sample of metal to which fine thermocouple wires are .attached, cooling the sample rapidly by means of a blast of helium, and recording the temperature vs. time curve on an oscillograph. The apparatus used in the present experiments was very similar to that described by Greninger, except that the heating coil was made of molybdenum wire (0.040 in. in diam) and the assembly was placed in a bell jar under vacuum of 10" mm Hg or less. The specimen con-

Jan 1, 1952

By J. L. Lytton, T. E. Tietz

The recovery of flour stress of recrylstallized are-cast molyhdenum after 10 pet prestrain at 150°C was studied for recovery temperatures of 725o 801o and 901oC. The yield phenomenon observed after recovery was such that the yield drop diminished and disappeared as recovery proceeded. A fractional flow-stress recovery parameter fr, was developed and it was shown that, independent of recovery ternperatzwe, equal f, values of about 0.16 produced superposition of stress-strain curves except for the initial yield phenomenon. The activation energy for low degrees oj recovery was found to he 32,600 cal per mole, in close agreement with a reported valzie for difjzdsion of carbon in molybdenum. For values of fr greater than about 0.15 where recovery was thought to predominate over strain aging, the activation energy varied from 43,600 to 82,500 cal per mole and increased with increasing recovery and recovery temperature. A 10 pet pre-strain shifted the ductile-brittle transition temperature, as indicated by tensile elongation, from 10°C. to about -22"C. A 100-hr recovery anneal at 801oC was shown to shift the transition temperature to about —5oC, the major shift occurring during the first hour of recovery. ThE recovery of cold-worked metals is accompanied by changes in the arrangement of the dislocations generated during cold work. In recent years, the use of transmission electron microscopy has allowed direct observation of some of the dislocation configurations formed, such as "tangles", networks, cells, subboundaries, and loops. These observations have recently been reviewed by Keh and Weissmann.' Accordingly, some understanding has been gained concerning the nature of the dislocation rearrangements involved. The relationship between recovery and the decrease of mechanical strength, however, has only recently been explored. Previous efforts of the authors were directed toward investigating the decrease of tensile flow stress of aluminum and Al-1 pet Mg during recovery2 and the effects of concurrent elastic and plastic strain on the flow-stress recovery of aluminum.3 In addition, an investigation of the effect of four different solute additions on the tensile flow-stress recovery of aluminum has been made.4 Many important metals, however, have the bee crystal structure, e.g., iron, molybdenum, tungsten, vanadium, columbium, tantalum, and chromium. These metals might be expected to behave differently during recovery from those with other crystal structures for two reasons. First, the previous and present investigations of aluminum and its alloys have shown that tensile flow-stress recovery is strongly affected by foreign atoms in the lattice. In the bee metals, interstitial impurities are usually present and might be expected to have various effects upon both flow behavior and recovery. For example, heterogeneous yielding might occur during prestraining. The diffusion of these interstitials can also relock dislocations during recovery so that a yield point might be obtained upon restraining (strain aging). Second, some of the bee metals (e.g., iron, molybdenum, and tungsten) undergo the familiar structure-sensitive ductility transition and exhibit brittle fracture at suitably low temperatures. It is of considerable interest to determine the effects of recovery on the transition temperature of such metals. The purpose of the present investigation was to extend our knowledge of flow-stress recovery to the bee metals and to study the effects of recovery on the ductile-brittle transition temperature. Unalloyed recrystallized are-cast molybdenum was chosen as a test material, since it exhibits a well-defined ductility transition near room temperature. As in the previous studies, the degree of recovery was measured in terms of the decrease of tensile flow stress. Prestraining and restraining were carried out at 150°C to avoid the possibility of brittle fracture. EXPERIMENTAL PROCEDURE The are-cast molybdenum sheet used in this investigation was obtained from the Climax Molybdenum Co. of Michigan. It was 0.060 in. thick by 10

Jan 1, 1964