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By Orville R. Lyons

This paper presents a method whereby the amount of misplaced material and the difficulty of the separation can be used to compare coal cleaning equipment of all types, from effectiveness and capacity standpoints. The correlations presented do not include all types of equipment currently available, but the method can be used to evaluate any make or type of coal cleaning equipment, both old and new. THE relative performance of coal washing equipment, or the effectiveness with which any type or make of equipment removes impurities from coal, has been most difficult to evaluate in the past. The most widely used yardstick is the Frazer and Yancey efficiency formula developed in 1922,' but Yancey in a later article states that "washers treating coals of different density composition or operating at different densities of separation cannot be compared directly on the basis of this criterion."' Prior to and since 1922, a variety of other methods has been used for comparison purposes, including the distribution curve, the error area, and the "ecart probable" or probable error. Yancey and Geer in discussing these methods conclude, "Performance can be evaluated in a number of different ways, with the choice of the proper method to use being dictated by the objectives of the investigation and the data available."' It is true that performance can be evaluated in a variety of ways, but if the equipment is to be evaluated on an effectiveness basis, there should be only one universal comparison method. Varying methods have been used because one universal comparison method has not been found or developed. In the article previously quoted, Yancey and Geer state in clear terms the primary concept for a universal comparison method: "One of the simplest, and certainly one of the most obvious evaluations of washery performance is the quantity of sink material in the washed coal and the float material in the refuse. If the washery products are tested at the density at which the washing unit is operated, the sink in the washed coal and the float in the refuse represent material that has been misplaced." The quantity of misplaced material was used as a criterion of washery performance by Lincoln in 1913," by the United States Bureau of Mines in 1938,' by Hancock in 1947," and by the national French research agency Cerchar in recent years.' In 1950 Andersone proposed the use of this criterion as an efficiency value to replace the Frazer and Yancey formula. However, none of the above-mentioned investigators used the misplaced material concept in a manner that would provide universal coal-cleaning equipment comparisons. The Correlation Theory The ideal coal cleaning process would treat all sizes and would make a perfect separation at any given specific gravity. All material lower in density than the desired value would report in the coal product and all material higher in density would report in the refuse product. Unfortunately, no known cleaning process achieves this goal and there seems little likelihood that any process yet to be invented will do more than approach it. When coal is treated in volume under operating conditions, it is impossible to avoid mechanical entrapment, fluctuations in throughput and effective gravity of separation, and the creation of turbulent currents, even when a true heavy-liquid bath is used and the feed is closely sized and contains little intermediate gravity material. This being so, it is possible to appreciate the difficulties inherent in trying to obtain a perfect separation when treating a wide range of sizes and a feed containing high percentages of intermediate material, using turbulent currents to help create the effective separation gravity, under operating conditions which normally tend to be on the overload side. When coal is separated from refuse in any coal cleaning equipment, some refuse always reports to the coal and some coal to the refuse; the writer therefore assumed that there should be a relationship between the total amount of misplaced material produced by any given piece of equipment and the difficulty of separation as represented by the percentage of near gravity material in the feed. With small amounts of near gravity or k0.1 material in the feed there should be less misplacement of material than would occur with large amounts of near

Jan 1, 1953

By R. V. Wheeler, F. V. Tideswell

The suggestion, which appears to find increasing favor, that the elementary composition of coals should be used as the basis of their classification, makes it important that our methods of expressing this elementary composition should be refined. It is doubtful whether many published ultimate analyses of coals are accurate within 1 per cent. carbon. Many are probably as much as 3 or 2 per cent. in error in carbon. It is, perhaps, not always realized that the extreme variation in carbon content over the range of bituminous coals is but 16 per cent. An error of 1 per cent. in carbon on analysis corresponds therefore with one-sixteenth of the whole range, an amount which is unjustifiable even for the application to commercial needs of a system of classification based on ultimate analysis. For the purposes of scientific research such limited accuracy is of little value. The object of this paper is to examine one important source of error in ultimate analysis, that due to the occurrence in coal of inorganic material, and to emphasize the need for the quantitative measurements (of carbon more particularly) on coals to be rendered independent of their contents of such material. Adventitious Mineral Matter When a determination of calorific value, for example, is made upon a coal, it is customary to correct the measured value with respect to the ash left by the coal on incineration, recording the result as being on an "ash-free basis." It is desirable, before discussing the adequacy of the correction thus made, to demonstrate its validity, which turns on whether the ash is present as an adventitious impurity in the coal and forms no part of the coal substance itself. The inorganic matter associated with coal has originated in several ways1: 1. The "inherent" inorganic materials, occurring as constituents of the plants from which the coal-forming deposits were derived.

Jan 1, 1928

By P. A. E. Armstrong

For hardening, the steel should be heated just above the critical temperature, as a properly forged steel quenched from this heat has some toughness, with maximum hardness and density. The quenching bath should be kept at approximately a constant temperature. In drawing or tempering, longer time periods at lower temperatures are better than the shorter periods at higher temperatures. A recent questionnaire brought out the following facts: The double-arc, double-taper bit and cross bit give the best service; resharpening of the drill steels most frequently was caused by wear of gage; chipping of bit was, with one exception, the second cause; breakage of shank was with one exception, the third cause; upsetting of bit, upsetting of shank, and breakage of body were the fourth cause. The direct cause of the necessity of resharpening given was poor heat treatment first, and faulty steel and severe rock conditions second. Rock-drill Steel By N. B. Hoffman Much of the drill steel produced in America is manufactured into hollow rods. After all forging has been finished the entire bar should be thoroughly annealed and heat trcated before the point is hardened and tempered, at least where the drill has heavy duty, such as working in deep holes and hard rock. A large number of tools of like character and doing work requiring similar properties are now being made of alloy steels; for instance a steel with a carbon content of 0.75 to 0.85 per cent., manganese 0.25 to 0.35, phosphorus under 0.020, sulfur under 0.030 and vanadium 0.18 to 0.23 has proved very satisfactory for chisels. A steel of this type properly manufactured will be far superior to the ordinary carbon steels, while at the same time is one whose price range should not prove prohibitive. Much of the cutting efficiency of a drill depends on proper heat treatment. This vanadium steel has a far greater safe hardening and forging range than ordinary carbon steels, making it more nearly foolproof. It will harden nicely at 1400" I?. but will also harden well and satisfactorily at 1550" F. retaining its quality of hardness and toughness in either case. Drill Steel from Hollow Ingots BY P. A. E. ARmstrong Hollow drill steel is made from a drilled billet with a sand-filled core or the drilled, pierced, or the drilled and pierced billet, not sand-filled, is rolled down over a projectile or ball much the same as in ordinary pipe manufacture. The first is the method generally used in this country; the second is used largely in Sweden. Swedish hollow drill steel has a

Jan 1, 1922

By Horace Pops

Martensitic transformation has been studied during cooling and heating in ß-phase Cu-Zn alloys to which small additions have been made of Ni, Ag, Au, Cd, Ga, In, Si, Ge, Sn, and Sb. The start and finish of the martensitic reaction and the variation of transformation temperature with third-element content were determined by electrical-resistivity measurements from alloys which had either constant zinc contents or constant values of electron concentration. All of the third elements, except nickel, lowered the transformation temperature if the results were plotted along the lines of constant zinc contents in each ternary system. A significant difference in the rate of lowering of the transformation temperature per atomic percent of the third element was observed for elements which had the same nominal valence. No systematic variation of transformation temperature with the valence of the third element was observed. It is suggested that the observed increase in transformation temperature for nickel-bearing alloys is due to the transfer of electrons from the conduction band of the alloy to virtual bound states. However, electron concentration is not the most important factor controlling the instability of the 0 phase. The transformation temperatures of the ternary alloys can be predicted from the following approxilnate expression: Ms (°K) = +3280 - 80 Zn + 8 Ni - 30 Ag - 12 Au - 140 Cd -90 Ga- 145 In - 80 Ge -175 Sn - 120 Si - 150 Sb MOST binary ß -phase alloys based upon the noble metals copper, silver, and gold are unstable at low temperatures and transform spontaneously by a martensitic reaction. This transformation has been studied recently in the ordered bcc ß'-phase Cu-Zn binary al1oys1,2 where the transformation temperature is below the room temperature and decreases with an increase in zinc content. It has been reported that the transformation temperatures can be raised above room temperature by small additions of a third element such as silicon3,5 or gallium,4,5 but no quantitative study has been made. The transformation temperature of different binary alloys can be altered by third-element additions. For example, it was shown that nickel and copper may have a large effect on the Ms temperature of CU-Al6 and Au-cd7 alloys. The present investigation was made to determine systematically the influence of various third elements on the martensite-transformation temperature of Cu-Zn ß-phase alloys. Since these alloys have an electronic origin,' alloy compositions were chosen so that the transformation temperatures could be determined at constant zinc contents or at constant values of electron concentration. I) EXPERIMENTAL PROCEDURE Ten ternary alloy systems were obtained by adding nickel, the noble metals silver and gold, and some B-subgroup elements to a Cu-Zn matrix. These are arranged according to their rows and columns in the Period Table as follows: Each ternary alloy was prepared by melting and casting weighed quantities of high-purity metals (99.99+ pct) in sealed quartz tubes under a partial pressure of helium to make a 4-g ingot. The molten alloys were shaken vigorously and then quenched in water. Since the weight loss was negligible, the compositions of the ingots after casting and annealing were assumed to be the same as the nominal compositions. The ingots were homogenized after casting in helium-filled Vycor tubes for 24 hr at temperatures between 750" and 810 C and quenched into brine. Metallographic examination revealed that all alloys were homogeneous, poly crystalline ß-phase alloys, and that the grain size was in the range 1 to 5 mm. Electrical-resistivity measurements were made to determine transformation temperatures of the ternary alloys during continuous cooling or heating. Transformation temperatures of the ternary alloys can be determined by electrical-resistivity measurements since the resistance of the martensitic phase is much higher than that of the 0' phase. The technique has been described previously in connection with a study of Au-Zn alloys.9 The reproducibility of transformation temperature was approximately ±6°C. II)RESULTS A hysteresis was always observed in electrical-resistivity curves and was usually less than about 12°C.

Jan 1, 1967

By George Smith

THE discussion of production control at the Insti-tute's annual meeting was profitable in that it started some thinking. One pertinent question there raised was how the opening of new mines, whose now unneeded output would inevitably react on prices, could be postponed until the market was ready to ab-sorb their output. In transportation some restraint is being exercised on building new and unneeded fa-cilities, but in production no general solution of the problem has been offered. Yet in a special field of new development of increasing importance a quick and easy solution has been suggested. I refer to the leasing of mineral deposits on the public domain. The report of the United States Coal Commission says: "In its possession of approximately 60,000,000 acres of public coal lands the Federal Government has its own responsibility in restraining overdevelopment. The Department of the Interior, which has control of these lands under the leasing law, construes the law as giving the department no discretionary right to refuse applications for leases regularly made by qualified citizens or corporations for coal deposits on the pub-lic domain. It presumes, that the applicant believes there is a need for an additional coal mine in his local-ity and expects to find a market for the coal produced. Overdevelopment, however, is a fact in the West as well as in the East, and wherever present involves needless investment and excessive costs of operation. Plainly, the leasing law should be so amended as to give the executive department full discretion in its approval of the opening of a new coal mine on the, public domain, making that approval contingent on the showing that the new mine would serve the in-terest of the consuming public." A similar view of the Federal Government's re-sponsibility in restraining overdevelopment of its oil lands has been presented by the Federal Oil Conser-vation Board. The leasing law vests little or no dis-cretion in the Secretary of the Interior, who must grant a lease to a qualified applicant regardless of any obvious country-wide condition of surplus produc-tion. This wide-open invitation to exploit and dissi-pate the people's own oil resources could be restricted by simple legislation, although as remarked by the board, such a legislative remedy for a plain economic ailment "does not evoke popular support when its purpose is avowedly to put the brakes on development of natural resources."

Jan 5, 1928

By Klaus Schwerdtfeger

The equilibria in the Cr-N system have been investigated in the temperature range 1100° to 1310°C by reacting chromium powder with Nz-Hz gas mixtures. The solubility of nitrogen in chromium in equilibrium with chromium subsitvide ("Cr,N") is given by Chromium subnitride is nonstoichiotnetric; its nitrogen content is always less than that corresponding to the formula CrzN. The lattice paranzeters of quenched samples have been measured; c, and a. parameters are found to increase with increasing nitrogen content. The growth rate of the subnitride layer on chromium plates was measured by a thermogravimetric technique, using a silica spring balance. The self-diffu-sivity obtained from the theoretical analysis of the parabolic rate constant is found to decrease with increasing nitrogen content, i.e., with decreasing vacancy concentration in the nitrogen sublattice. The intrinsic nitrogen diffuivity is derived from another series of rate measurements using "CrzN" plates; the intrinsic diffusivity, DN = 3.2 X 10-a cmZ sec-' at 1200 C, is found to be essentially independent of- the subnitride composition. The concentration gradient was measured in a chromium subnitride layer by the X-ray method and found to be consistent with the derived diffusivity value. TWO chromium nitrides are known to exist:' the nonstoichiometric subnitride "CrzN" and the nitride CrN. In the present work the kinetics of the formation of chromium subnitride from chromium and nitrogen have been investigated at 1100" and 1200°C. In additional experiments the relevant equilibria have been measured. The data are used to evaluate the diffusivity of nitrogen in chromium subnitride. Since chromium nitrides are often found in chromium-containing steels, the results are expected to be helpful in the interpretation of the chemical reactions between chromium steels and nitrogen. Equilibria in the Cr-N system have been determined by several investigators.2"3 The rate of nitriding of chromium was measured by Arkharov et a1 .' in ammonia in the temperature range 800" to 1200°C. The parabolic rate law was observed. Due to the undefined nitrogen activity of the ammonia atmosphere, it is dif- ficult to interpret these rate data theoretically. An additional difficulty arises from the fact that the two-layer scale consisting of CrN and "Cr2N" was formed at the temperatures below 1030°C. The rate of nitriding of technical chromium (95 pct) was measured by Zaks in nitrogen at -1 atm in the temperature range 800" to 1300°C. EXPERIMENTAL METHODS The chromium samples were reacted with Nz-HZ gas mixtures in a vertical tube furnace, wound with Pt-10 pct Rh resistance wire. The gas-tight reaction tube was of high-purity recrystallized alumina. In nitrogen solubility measurements the nitrogen content of chromium was determined by the Kjeldahl method, on samples quenched in the cold part of the furnace. The nonstoichiometry of the subnitride and the nitriding rates were measured thermogravimetrically using a sensitive (+0.1 mg) silica spring balance. For the equilibrium measurements samples of 1 g of chromium powder contained in high-purity alumina crucibles were used. In order to remove most of the oxygen and nitrogen impurities from the chromium, the samples were initially annealed in purified hydrogen until a constant weight was obtained. The chromium plates (approximate dimensions 2 by 1 by 0.08 cm) used for the rate measurements were machined from ingots obtained by arc-melting of iodine-processed chromium. According to manufacturers' specifications the purity of the chromium powder was 99.9 pct Cr and that of the iodine-processed chromium 99.99 pct Cr. Our own spectroscopic analysis of the chromium powder yielded 0.02 pct Fe, 0.05 pct Mn, 0.05 pct Si, and 0.02 pct Ti as major impurities with all the other detectable elements below 0.005 pct. The nitrogen partial pressure of the gas phase was controlled by mixing prepurified hydrogen and nitrogen with constant pressure head capillary flowmeters. Oxygen and water vapor were removed from the mixed gas by passing it through columns of platinized asbestos (450°C) and anhydrone. The gas flowed upward in the furnace with flow rates of 300 to 500 cu cm per min (25"~). Gas tightness of the furnace system was ensured by pressure checks at regular intervals. The furnace temperature was controlled electronically in the usual manner. The reported temperatures were measured with a Pt/Pt-10 pct Rh thermocouple and are estimated to be accurate within ±$C The X-ray measurements were made with a Debye-Scherrer camera and a diffractometer using chromium radiation {\Ka = 2.29092A). EQUILIBRIUM MEASUREMENTS The experimental results of the equilibrium measurements are contained in Tables I to In. Fig. 1 shows the solubility of nitrogen in solid chromium in the temperature range 1100" to 1310°C. In Figs.

Jan 1, 1968

Chairman James F. Kemp :—I call upon Prof. Albert Sau-veur to open the general discussion on the subject of sound steel ingots. Albert Sauveur, Cambridge, Mass.:—I believe that I have the privilege, this instant, of opening one of the most important, significant and far-reaching discussions ever held before this Institute. I believe that our talk this morning will be heard far and wide, not only on this side of the water, but also on the other side. Everywhere, in short, where steel is manufactured. I welcome the opportunity of expressing my mind in regard to the production of sound steel ingots, a subject about which I entertain Rome very decided opinions. I shall not attempt to discuss the relative merits of the various processes that have been described. I shall, on the contrary, confine myself to a general consideration of the important questions before us. So intensely do I feel regarding the production of sound steel ingots that I have put my remarks on paper in my desire to present my views. I have always remembered my astonishment, not to say dismay, some twenty years ago, when I beheld for the first time a steel ingot, a rail steel ingot, that had been split open from top to bottom. There was the pipe! There was an enormous cavity extending from a short distance from the top to nearly one-third of the length of the ingot; and when I was told that manufacturers seldom discarded more than one-tenth of the total length of the ingot it became evident to me that a large proportion of piped, and therefore, defective, or at least weak and dangerous, rails were manufactured; in fact, probably one rail out of six, nearly 17 per cent., of dangerous rails.

Jan 1, 1914

By D. Turnbull

RECENTLY it has been shown that aggregates of small liquid droplets of tin,' mercury' or gallium' kept from intercommunicating by suitable films do not solidify at an appreciable rate unless the supercooling is very much greater than that usually necessary to cause a large continuous mass of the metal to solidify. For example oxide-coated tin droplets1 (1 to 10 micron diam) must be supercooled 100" to 110°C before their rate of solidification becomes rapid, although large continuous masses of liquid tin usually begin to solidify when supercooled 30" or less." These experiments have been interpreted' on the hypothesis that nucleation of crystals in large continuous metal samples is almost always catalyzed by accidental inclusions. Therefore, if the metal is dispersed into a number of isolated droplets large in comparison with the number of inclusions, most of the droplets should not crystallize until a temperature sufficiently below the thermodynamic melt- ing point, T0, has been reached to permit an ap-preciable rate of homogeneous (noncatalyzed) nu-cleation. In general this temperature is very much less than the temperature at which the rate of heterogeneous (catalyzed) nucleation is appreciable. If this interpretation is correct then investigation of the kinetics of crystallization of small droplet aggregates should be one of the most fruitful methods of obtaining information about the homogeneous formation of crystal nuclei in liquids. In this paper the results obtained on gallium and mercury are reported more fully. In addition, results on the supercooling of aggregates of liquid bismuth and lead droplets are included. Experimental Procedure Some care must be exercised in the choice of a barrier to prevent the liquid droplets from coalescing lest the barrier itself catalyze the formation of crystal nuclei. From this standpoint, the most desirable barrier is vacuum or inert gas, but, because of the difficulty of experimental arrangement, solid compounds and adsorbed films have been selected. There are two guiding principles in the selection of compounds as barriers in addition to the requirement that they prevent coalescence of liquid droplets. First, the compound should be almost insoluble in the liquid metal and second, its lattice structure should be quite unlike that of the forming metal crystal. Even so, there is no a priori assurance that the solid compound film will not catalyze to some extent the formation of metal crystal nuclei. An adsorbed protective monolayer is less likely to be catalytic than a crystal compound film, but it is not often possible to find a monolayer that is stable under the experimental conditions. Also, it is desirable, though not generally essential, that the aggregate of droplets be prepared by breaking up the massive metal rather than from a powdered compound of the metal.

Jan 1, 1951

By B. B. Rath, Hsun Hu

In wedge-shaped bicrystals of zone-refined aluminum it is observed that (111) pure tilt boundaries migrate under the driving force of their own inter-facial free energy. The boundary velocity is a power function of the driving force. The driving force exponent decreases with decreasing angle of misorien-tation. For example, at 64O°C, the exponent decreased from 4.0 for a 40 deg to 3.2 for a 16 deg tilt boundary. An evaluation of the driving force acting on the boundaries during their motion indicates that for low driv-forces, up to about 2 x l03 ergs per cu cm, the velocity is relatively independent of misorientation, whereas at higher driving forces a 40 deg tilt boundary exhibits the highest velocity. The measured activation energy for boundary migration approaches that for bulk self-diffusion at low driving forces, decreasing from 33 to 27 kcal per mole as the driving force is increased from 1 x l0 to 5 x l03 ergs per cu cm. These results are compared with current theories of grain-boundary migration. In previous experimental studies of grain boundary migration the driving force has been limited to a difference in stored energy across the boundary. This stored energy has been introduced into the crystal either by prior deformation1-3 or by grown-in lineage structure. A part of the energy stored in the deformed crystal is released by recovery either prior to or concurrently with grain boundary migration, thus introducing an uncertainty as to the magnitude of the driving force responsible for grain boundary migration. The grown-in lineage structure, though thermally stable during annealing, neither provides conditions under which different levels of energy may be stored in the imperfect crystal nor provides a control of orientation difference across the migrating boundary of a growing grain. Furthermore, because of variation in the lineage structure, it is difficult to determine accurately the energy stored in the imperfect crystal. Several investigations of grain boundary migration during normal grain growth have also suffered from difficulties in estimating the driving force because of uncertainties in the principal radii of curvature.~ In the present investigation the velocity of pure tilt boundaries in zone-refined aluminum bicrystals of selected orientation (40, 30, and 16 deg around the [Ill] tilt axis) has been measured in the absence of a dislocation density difference across the moving boundary, thus eliminating the previous experimental difficulties. The driving force for boundary migration is derived from a gradient of the total interfacial free energy of the migrating boundary in wedge-shaped bicrystals. A similar method was attempted by Bron and Machlin in a study of grain boundary migration in silver. However, they found that one of the crystals was deformed and consequently the motion of the boundary was partly due to a difference of stored energy across the boundary. The observed behavior of boundary velocities as affected by the driving force is examined in the light of the predictions of the current theories of grain boundary migration.7"10 The effect of boundary misorientation on velocity is compared with the theory of " which is based on a dislocation core model for high-angle boundaries. EXPERIMENTAL METHOD Seed-oriented bicrystals of zone-refined aluminum, 2.5 cm wide, 0.5 cm thick, and 12 cm long, containing tilt boundaries with a common (111) axis, were grown from the melt in the direction of this axis. Spectro-graphic analysis, reported earlier,'' indicated the purity of the crystals to be 99.999+pct. Three such bicrystals containing 16, 30, and 40 deg tilt boundaries were used. Wedge-shaped specimens were prepared from these bicrystals by spark cutting followed by electrolytic polishing. The angle of the wedge was usually 40 deg and the specimens were usually 0.25 cm thick. The intercrystalline boundary was located within 0.2 to 0.5 cm from the tip of the wedge. Fig. 1 shows a section of an oriented bicrystal containing an outline of a wedge-shaped specimen. The crystallographic directions shown in Fig. 1 represent the orientation of one of the crystals (the larger section of the bicrys-tal); the orientation of the other crystal differs only by rotation around the common [lil] axis. The parallel faces of the wedge always corresponded to the common (171) planes in both crystals, whereas the orientation of the side faces varied, depending on the misorientation angle. The bicrystal orientations were determined

Jan 1, 1970

By George K. Burgess

The suggestion has often been made that it would be highly desirable, at least for certain grades of steel, to be able to control more certainly, by pyrometric measurement or otherwise, the temperatures of the operation of the open-hearth and other furnaces used in the manufacture of steel and iron. That the properties of the ingot and finally of the finished steel product are intimately related to the final temperatures attained by the metal in the furnace and to the temperature of casting, has been recognized by metallurgists for a long time, and this question has been very thoroughly treated, especially by Prof. Henry M. Howe, and by A. W. and H. Brearley.1 It is the object of this paper to demonstrate that in so far as casting temperatures of furnaces, steel ingots, and similar operations involving the temperatures of streams of iron and steel are concerned, well-known pyrometric methods may be easily applied—if certain recently determined corrections are made and precautions taken—with a relatively high degree of accuracy. Greater but not insurmountable difficulties will be encountered in the case of open-hearth furnace temperatures, while for those of the converter type a ready solution does not seem practicable. It appears that certain essential data have always been lacking in reporting observations such as those of Le Chatelier in 1892. Thus the correction to be applied for the characteristic radiation—or emissivity— has hitherto not been adequately taken into account nor the limitations of the several possible pyrometric methods clearly recognized. In view of the fact that in open-hearth practice the temperature to be measured in the furnace itself, in the stream of metal and slag and in the teeming into molds or ingots is usually above 1,500" C. (2,700" F.)

Jan 1, 1917

By L. D. Mullins, E. B. Elfrink, W. L. Wahl

In the past few years several articles and papers presenting results of solution gas-drive depletion calculations have appeared in the lit-erature. Such calculations are of interest to the oil industry, for investment decisions must often be made before much is known about a reservoir. At other times, an estimate of the possible benefits to be realized from alternate production methods is desirable, and theoretical depletion calculations can serve as a floor or reference level from which to work. In any case, an estimate of ultimate oil recovery based upon engineering data is commonly required. An engineer confronted with the problem of obtaining, for a specific reservoir system, an estimate of ultimate oil recovery by solution gas-drive depletion usually will be forced to perform the calculations himself. This is despite the quantity of data in the literature. Rarely will either experience or the literature provide results from a reservoir system similar in all important respects to the one under consideration, and calculated results are not so plentiful that satisfactory interpolation procedures can be devised. Performing the calculations, however, is a tedious, time-consuming task unless an electronic computer is available, and, in practice, time and manpower are not always available for this purpose. A quick, simple, consistent method was needed for reducing the uncer- tainty in estimated oil recovery from solution gas-drive reservoirs when only minimum information about the reservoir system is available. PROCEDURE Method of Calculation The usual requisite assumptions were made so that the material balance equation could be used to calculate data for the charts. The following assumptions were made: (1) the reservoir is homogeneous and isotropic; (2) oil recovery is due entirely to solution gas drive and neither a gas cap nor a water drive nor gravity drainage is present; (3) the initial reservoir pressure is the bubble-point pressure of the reservoir fluid; (4) initial total liquid saturation is 100 per cent of pore space; (5) interstitial water saturation remains at the initial value as the reservoir pressure declines from the bubble-point pressure to atmospheric pressure; (6) equilibrium gas saturation is 5 per cent of pore space; and (7) oil and gas saturations are uniformly distributed throughout the reservoir at all times. There are no saturation gradients due to a wellbore, nor is the geometry of the reservoir system considered. The material balance equation was written in the form of a differential equation' which was integrated to determine the change in oil saturation for an assigned pressure drop. Formal integration was not possible, so recourse was made to the Runge-Kutta method' of numerical integration. All computations were performed on IBM equipment. Numerical integration yielded the change in oil saturation within the reservoir as the pressure. declined from the bubble-point pressure to atmospheric pressure. The initial oil saturation minus the change in oil saturation yielded the oil saturation at atmospheric pressure. The oil originally in place was obtained by dividing the initial oil saturation by the initial formation volume factor (differential liberation) while the oil in place at atmospheric pressure was obtained by dividing the final oil saturation by the formation volume factor at atmospheric pressure. U1timate oil recovery, expressed as a percentage of the initial oil in place, was obtained by dividing the difference between the oil initially in place and the oil in place at atmospheric pressure by the oil initially in place and multiplying by 100. PVT Data Charts were based upon 135 solutions to the material balance equation. PVT properties of the reservoir fluids were variables in this equation and had to be known as functions of pressure. The required PVT data might have been obtained from either actual reservoir fluid systems or correlated data. However, correlated PVT data were developed and employed in the calculations for the following reasons: (1) it is doubtful that 135 sets of PVT data could have been obtained for the values of variables investigated in this study, and (2) results of recovery calculations from randomly obtained PVT data could not be correlated as well as rssults from selected PVT data. The PVT data used to develop the

By F. N. Lewis

HOPPERS may be divided into three classes accord-ing ing to type, namely; round, suspension, and square or rectangular. The round hopper is the most simple of the three and is the most easily understood by laymen. The bursting pressure in these hoppers is resisted by the shell as a round tank resists the pressure of water. They are perfectly stable as long as they retain their cylindrical shape. In holding water or other liquids the pressure on the sides is always the same, but with ore this is seldom true. These round hoppers for ore are usually filled from one side and frequently drawn off from the opposite side. A considerable amount of fine material is deposited on the side of filling, while the coarse particles run over to the far side. This con-dition of filling causes the ore to flow out unevenly, sometimes drawing out completely on one side. After a time the remaining ore breaks down and the shock of this falling ore, sometimes many tons, is very severe and has a tendency to distort the hopper from its true shape; consequently a round hopper holding ore should be stiffened to resist such shocks as this. For these reasons round hoppers have never been popular with engineers for holding ore. They also have a trait of welling. up worse than the square hoppers. I have known of round hoppers holding fine ore, 25 ft. diam-eter, that had a tendency to well up until there was ,only a round hole about 6 ft. diameter leading to the gate. The suspension or catenary type of hopper is well adapted to conditions requiring only small capacities. The selection of this type of hopper should be done with considerable care. The theory of the design is that the loads are carried entirely in tension in the bottom plate, just as a hammock carries a load, and it must be remembered that the bottoms of these hoppers are con-stantly changing form with the changing loads. Take a hammock, for example, if a person stands up in it, the bottom goes down and the sides approach a straight line; if the person lies down then the bottom comes up and the sides go down. It is this constant changing of the form that has a tendency to break the connections.

Jan 7, 1928

By P. R. Swann, R. M. Fisher, W. R. Duff

Observations by transmission electron microscoPy of the microstructures in quenched Fe-A1 alloys provide the basis for revision of the Fe-A1 phase diagram. The diagram is characterized by first-order transitions between the disordered ferromagnetic a! solid solution and the two ordered phases FeAl and Fe3Al. However, second-order transitions showing corre-luted fluctuations in the degree of order near the critical temperature occur between the paramagnetic a! and FeAl and the FeAl and FeAl phases. ELUCIDATION of the true nature of the Fe-A1 equilibrium diagram has proven to be an extremely challenging problem as evidenced by the fact that six quite different versions of the iron-rich portions have been published since 1958.&apos;-6 The most recent diagram shown in Fig. 1 was proposed by Rimlinger6 based on X-ray diffraction analysis and dilatometry. He shows the single-phase fields of the ordered phases of the types FeAl(L20) and Fe3A1(D03) and the disordered bcc Fe-A1 solid solution, a, separated from one another in every case by large two-phase regions. Thus, he implies that in Fe-A1 alloys ordering always occurs via a classical phase transition. In contrast to this view, theoretical studies3&apos;" have concluded that the transitions a — FeAl and FeAl — Fe3A1 should be second- or higher-order in character. A further revision of the Fe-A1 diagram of state is presented in Fig. 2 based on our previous8,9 and current observations by transmission electron microscopy of the microstructures found in quenched specimens. The average quenching rate was 50,000°C per sec and although this was sufficiently rapid to prevent Fe3A1-type ordering below -550°C it was not sufficient to suppress ordering of the type FeAl in the temperature range -650" to 800°C. However, ordering during the quench is easily recognized since the FeAl phase formed during holding at temperature rapidly develops large antiphase domains whereas FeAl formed during quenching has a very fine domain size, -50Å diam. Each symbol in Fig. 2 indicates the result of a micro-structural examination of specimens heat-treated (-1 hr) until equilibrium was established. The results will be discussed in the order of the points, A - Dl marked on both diagrams with reference to the phase fields, critical temperature, and "transition" ranges, respectively. The dark-field micrograph corresponding to point A is shown in Fig. 3 and was obtained using a 111 reflection of the Fe3A1 structure which reveals both a&apos;/4<111> and a&apos;/2<100> antiphase domain boundaries.10 The structure observed in this specimen is typical of single-phase Fe3A1 and shows no evidence for FeAl or any other additional phase. The dark-field micrograph shown in Fig. 4 corresponds to point B on the two diagrams. It was taken with a 100 FeAl reflection which reveals a/2(111) domains and exhibits the typical structure of single-

Jan 1, 1970

By Dr. O’Neil Thomas J., Donald W. Gentry

"There are two characters to the value of mining properties-one mine may have a value, owing to its real intrinsic worth; an- other (having no intrinsic value) may have a value by being so situated as to harass the working of the really valuable mine-in mining camps one is looked upon as much of ligitimate enterprise as the other.&apos;&apos; -Victor Clement to Simeon Reed, Wardner, Jan. 11, 1890, Reed Mss. APPROACHES TO VALUATION Value and costs are two terms basic to financial analysis which cause, perhaps, the greatest amount of misunderstanding in mine valuation studies. When discussing the value of an asset or the costs of producing a product, it is absolutely essential that the speaker define the terms being used. There are many types of value and many categories of costs. Without specific accompanying definitions, value and cost data are not very meaningful. Value refers to a measure of the desirability of ownership of property. Some of the types of value which might be encountered in a mine valuation study are: 1) Market value 2) Full cash value 3) Salvage value 4) Replacement value 5) Capitalized value 6) Book value 7) Assessed value 8) Insured value Each of these has a specific meaning which can be applied to determine a monetary amount in a specific situation. The major item of interest here is the broader question of "what is the value of the mine?" or "what is the mine worth?" In this case, we are speaking of market value. Market value is the value (price) established in a public market by ex- changes between a willing buyer and a willing seller when neither is under duress to complete the transaction. Market value fluctuates with the degree of willingness of the buyer and seller and with the conditions of the sale. The use of the term market suggests the idea of barter. When numerous sales occur on a market, the

Jan 1, 1984

By E. Douglas Sethness

The uranium industry is booming. In Texas alone, there are about 22 different companies with active exploration programs. Twelve solution mines have been permitted; three surface mines have been authorized; and two mills are currently in operation. However, the industry also has a problem, and that is the disposal of radioactive wastes. Over the past several years, stories concerning nuclear wastes have appeared frequently in the news. One of the most frequently cited cases occurred in Grand Junction, Colorado. In 1966, after ten years of investigations, the U. S. Public Health Service (PHS) discovered that tailings from a uranium mill were being used as fill material and aggregate for local construction purposes. It was estimated that between 150,000 and 200,000 tons of material had been removed and used under streets, driveways, swimming pools, and sewer lines. In addition, tailings had been used under concrete slabs and around foundations of occupiable structures. Further studies prompted the Surgeon General to warn that the risk of leukemia and lung cancer could be doubled at the measured radiation levels. More recently, the L. B. Foster Company discovered that its building site in Washington, West Virginia, was radioactive. While digging a foundation, the ground erupted and a ball of fire 30 feet high shot out. Evidently, the dirt was laced with radioactive thorium and zirconium, a potentially explosive mixture contained in a Nigerian sand which had been used by the previous site owners in the manufacture of nuclear fuel rods. Just this month we have read about legal suits to stop exploration for a nuclear waste disposal site in Randall County, Texas. The U. S. Department of Energy is trying to locate a deep underground nuclear waste depository for final burial of over 76 million gallons of high-level wastes. The problem is acute, the wastes are accumulating at a rate of about 300,000 gallons per year. Nor do these numbers include the spent fuel elements from nuclear power plants that are in temporary storage facilities. Fortunately, public awareness of these and other related issues is high. Unfortunately, the differences in the waste products from the nuclear fuel cycle are not always apparent to the general public. There are two distinct types of radioactive wastes: "high-level", which consist of spent fuel or wastes from the reprocessing of spent fuel; and "low-level", which, in general, are by-product wastes. There are numerous non-technical definitions that can be applied to help the layman differentiate between high-level and low-level wastes. For this latter purpose, it is best to think of them in terms of what we can see and feel. In general, high-level wastes are physically hot and can cause acute radiation sickness in a short period of time. Low-level wastes are not hot, but may cause chronic health effects after long exposure. The wastes which we are concerned with in the uranium mining and milling industry are low-level wastes. As recently as ten years ago, there were very few controls or regulations governing tailings disposal methods. At the same time, mine reclamation was not enforced through either state or Federal laws and the long-term viability of abandoned tailings ponds was not assured. The regulatory climate has changed significantly in the last decade, however. The low-level radioactive wastes generated by uranium mining and milling are generally contained in a tailings pond. Approximately 85-97% of the total radioactivity contained in uranium ore is present in the mill waste that goes to such tailings ponds. The isotope Radium-226 is probably the most potentially harmful radioactive parameter in the ponds. Radium emits gamma radiation and is also an alpha particle emitter. Because gamma radiation is very penetrating, it presents a potential health problem when a source is located external to the body. Gamma radiation will go through the body, causing damage to each cell encountered on the way. Although alpha particles have very little penetration capability, they can cause extensive cell damage. For this reason, alpha particles are a problem after inhalation or ingestion. Radium creates a health hazard by both of these mechanisms. Radium decays to radon gas which can be inhaled and serve as an alpha particle emitter. Additionally, radium is very soluble and readily enters the natural hydrologic cycle if allowed to leach from a tailings pond. With a half-life of 1620 years, radium has plenty of time to be taken into the food chain and end up in our bodies, emitting alpha particles. Because the potential health problems are better understood today than ten years ago, and because the Nuclear Regulatory Commission (NRC) has developed increasingly stringent government regulations, the uranium mining industry applies a high level of technology to the disposal of nuclear wastes. In most cases, low-level radioactive wastes are disposed of at or near the site where they are produced. There are six commercial burial grounds for low-level wastes, but it would not be economical to ship all mine or milling wastes to these sites. The on-site disposal methods most often used are ponding

Jan 1, 1979

By J. Weertman

It is shown that conditions suitable for the conversion of straight dislocations into helices are common in crystals hardened either through long-range dislocation interaction or by jog formation on dislocation lines. For dislocations other than screw dislocations or dislocations nearly screw in orientation, helices can be formed with the aid of core difhsion at temperatures low relative to the melting point of the material. It is assumed that, once formed, helices will degenerate into dislocation tangles. THE dislocation tangle is one major dislocation arrangement observed by the transmission electron microscope which had not been predicted previously from dislocation theory. Since first being observed, the tangle itself has become important in applied dislocation theory. The Cambridge school (P. Hirsch) bases part of its theory of work hardening upon this phenomenon.&apos; In spite of the fact that the tangle obviously is a structure of prime importance to a cold-worked crystal, until recently relatively little attention has been paid to the question of exactly why or how tangles arise. It was generally assumed at first that cross slip and perhaps jog formation will lead to tangles. However Wilsdorf and schmitz2 concluded that cross slip cannot markedly contribute to the tangles. The Wilsdorfs and Maddin3&apos;* have proposed that tangles arise by means of prismatic dislocation loops and superjogs which are produced by a precipitation of point defects. Dislocations with superjogs or which have joined with prismatic loops are themselves dislocation sources and they can produce new dislocations, a process they call "mushrooming", which then tangle with each other. There is no question that the Wilsdorfs et al. mechanism as well as the cross slip one can pro- duce tangles. However these may not be the only paths leading to tangle formation. In this paper we wish to consider another way tangles can form. Our work will be based on the theory of dislocation screw We wish to show that whenever a crystal is cold-worked there always is present a "driving force" leading to helix formation and thus to tangle formation. We do not expect to observe well-developed, regular helices in cold-worked metals. Rather we would expect that after a helix is formed random interactions with nearby dislocations will cause a well-developed helical dislocation to go over into an irregular form. A dislocation tangle is, of course, a grouping of dislocations with irregular forms. Therefore we feel that, in accounting for the formation of dislocation tangles, it is sufficient to show that many dislocations in a cold-worked crystal will take on helical form. The crucial step is helix formation. In order to have a straight dislocation assume a helical form, diffusion of vacancies, or interstitials, must occur. Our discussion of tangle formation therefore will be limited to crystals which have been deformed at temperatures ranging from about one fifth to half their absolute melting point. In this tem per a ture range appreciable core diffusion is possible. (There is no problem in explaining tangle formation at temperatures greater than half the melting point. At such temperatures dislocation climb through a bulk diffusion of point defects gives an extra degree of freedom to dislocation motion. This extra freedom of motion makes tangle formation easy.) At temperatures much below one fifth the absolute melting temperature diffusion processes are sufficiently slow that helix formation does not occur. Of course helix or tangle formation could occur in material deformed at very low tempera-tures which was subsequently warmed in order to make the thin specimen necessary for electron-microscope examination. Since within the temperature range of approximately one fifth to half the melting temperature the diffusion of point defects in the core of the dislocation is rapid, we can reasonably assume that within

Jan 1, 1963

By Robert M. Dreyer

Diamond is the hardest known material. The diamond industry is separated into two major segments: (1) industrial and (2) gem. The major industrial use of diamonds is as a high-grade abrasive in a wide variety of applications for which no effective low-cost substitute is available. Table 15.4C.1 gives data on diamond production and on United States diamond imports. In both gem and industrial diamond production, but especially in the case of gems, the quantity of stones produced is not directly correlated to value, for the determination of value is markedly affected by quality. South Africa long has been the world leader in gem diamond production with an estimated production in 1968 (including South-West Africa) of about 4 million carats, equal to 43% of world production. During the past decade there has been a marked rise in gem diamond production in the USSR, so that nation is now the second-ranking world producer, with Angola (Portuguese West Africa) third. In natural industrial diamonds, the major production comes from Congo-Kinshasa (former Belgian Congo) which had a production of 15 million carats in 1968, equal to 47% of world production. In the past six years, there has been a remarkable increase of industrial diamond production in the USSR, which now ranks second with a production of 6 million carats, equal to about one-fifth of world production. South Africa now ranks third followed by Congo-Brazzaville and Ghana. No natural diamonds are produced in the United States. Since the discovery of an economical method for synthesizing small diamonds, there has been a rapid increase in synthetic production from 4 million carats in 1962 (1 carat equals 200 mg) to 11.5 million carats in 1968. Table 15.4C.1 gives the gradually increasing imports of industrial diamonds and indicates a relatively stable price. In 1968, the imports of industrial diamonds were 13.5 million carats with a value of $59.2 million. Table 15.4C.2 shows the rapidly increasing imports of cut and uncut gem diamonds into the United States. In 1968 the United States imported 2.5 million carats of uncut stones with a value of $166 million and 1.5 million carats of cut stones with a value of $175 million. Diamond marketing is dominated by DeBeers Consolidated Mines, Ltd., which, through subsidiaries, controls many of the major African diamond deposits. More important economically, however, through its Central Selling Organization, DeBeers markets 86% of the non-Communist world&apos;s natural diamonds, including most of those derived from deposits outside of corporate control. In 1967 the mines owned by DeBeers produced nearly 8 million carats, and the Central Selling Organization sold over $600 million in gem and industrial diamonds. The dominance of diamond sales by

Jan 1, 1976

By Carl Henrich

While the size of the blast-furnaces used for smelting leadand copper-ores has constantly increased, during late years, the manner of removing the slag from the furnace to the slag-dump has (until quite recently, at least) practically remained the same. The only appliance in practical use has been the two-wheeled slag-pot, holding about 300 to 400 pounds of slag, or as much as one man could manage to pull or push on a smooth, hard surface. When these slag-pots on wheels were first introduced, the daily capacity of a lead or copper blast-furnace hardly reached 30 tons of smelting-charge. With furnaces of this size, and even when the capacity had been doubled, this kind of slag-pot answered well enough. All the improvements made were in the construction of these pots and not in an increase of their size, which was limited by the weight which one man could handle. But the size of the furnace went on increasing. The daily consumption of one furnace rose gradually to 75 tons, 100 tons, and even more. The writer had occasion to construct some of these furnaces, which, in their day, were among the largest, if not the largest, of the water-jacketed kind then in existence. A water-jacketed furnace will always smelt more in the same time than a furnace with brick walls. Already, in those days, the slag-removal from these enlarged furnaces received due attention by the writer. The removal of 60 tons of slag in 24 hours, in pots holding only 300 pounds each, will require 400 pots to be filled in 24 hours; and when the amount.of slag increases to 90 tons, 600 pots will be required. That is to say, in the last case, 25 pots in each hour, or 1 pot each 24 minutes, will have to be put under the slagspout, filled, pulled away, and replaced with another. This will involve too much hurry for any degree of comfort, convenience

Jan 1, 1896

By Hsun Hu, R. S. Cline

The annealing behavior and the mechanism of re-orientation during recrystallization of iodide titanium cold-rolled 94 pct have been studied in detail. Results indicate that recrystallization occurs by the nucleation and growth of new grains, as in other common metals. Recrystallization nuclei form by the coalescence of subgraim, and the change in texture as a result of recrystallization is largely due to selective growth among the nuclei formed. The annealing of titanium is characterized by a wide range of overlap of the various stages of the annealing process, which may be responsible for a range of activation energies observed, and for the apparently gradual change in the annealing texture as a function of time or temperature. The deformation and recrystallization characteristics of titanium and zirconium are very similar. In cold-rolled strip, the deformation texture consists of two symmetrically oriented components, each having the basal plane laterally tilted at about 30 deg from the rolling plane and the [1010] direction parallel to the rolling direction. Upon annealing for recrystallization, the change in texture can be described, for simplicity,* as rotations around [0001].2&apos;6&apos;8 According to McGeary and Lustman,&apos; recrystallization occurs in zirconium through normal growth of the subgrains, which they called "domains", without the nucleation of new grains; and the magnitude of rotation around the [0001] axis increases gradually during the progress of recrystallization. If these conclusions were true, the mechanism of recrystallization in zirconium would be basically different from that in most metals, since it is commonly known that recrystallization with reori-entation always involves the migration of high-angle boundaries. In an attempt to clarify the situation, the mechanism of reorientation during recrystallization in iodide titanium cold-rolled 94 pct was studied in detail. The structural and textural changes upon annealing at various temperatures were examined by optical and transmission-electron microscopy, X-ray pole figures, pole density distribution measurements, and micro-beam techniques. EXPERIMENTAL PROCEDURE Material and Specimen Preparation. An iodide titanium crystal bar was are-melted and solidified in a cold-hearth crucible under a purified argon atmosphere. The solidified ingot had dimensions of approximately 3 by 1/2 by 3 in. One face of the ingot was somewhat uneven, but was as clean and shiny as the remaining parts of the ingot. Large grains with a Widmanstatten internal structure were clearly shown on the shiny surfaces, indicating the occurrence of P — a transformation upon rapid cooling from the melt. Analysis of the are-melted ingot indicated C 0.033, N 0.010, H 0.013, 0 0.002 in weight percent, and traces of iron, copper, and silicon as detectable impurities. The ingot was cold-rolled -40 pct to 0.300 in. thick with a reduction of 0.005 in. per pass. The defects on the uneven side of the ingot were then removed by machining. This reduced the thickness to 0.285 in. The piece was then recrystallized by annealing at 800°C for 1 hr in a fused silica boat charged into a fused silica tube furnace under a vacuum of 10~5 mm Hg. To refine the grain size, the recrystallized metal was again cold-rolled 40 pct to 0.170 in., then annealed at 700°C for 1 hr. These treatments yielded a strip with a uniform equiaxed grain structure, having a penultimate average grain diameter of 0.04 mm and a hardness of approximately 90 Dph. Final rolling reduced the thickness from 0.170 to 0.010 in., corresponding to a reduction of 94 pct. The strip was rolled in both directions by reversing end for end between passes. Surface lubrication was provided by oil-soaked pads attached to both rolls. Specimens of 1 in. length (for X-ray examinations) and +in. length (for hardness and microstructure examinations) were cut from the rolled strip, and a width of & in. was cut from the edges of each specimen by a jeweler&apos;s saw. These specimens were then etched in a solution of 10 cu cm HN03, 5 cu cm HF, and 50 cu cm H,O to 0.008 in. thick to remove the surface metal, as well as the distorted metal at the saw cuts, prior to annealing or measurements. To minimize any surface reaction with the atmosphere, all specimens were kept in an evacuated desiccator. Isothermal Anneals. All annealing treatments were conducted in vacuum in a fused silica tube furnace as described earlier. The temperature of the furnace was controlled to within *2"C. The specimen was placed in a fused silica boat, then pushed into the hot zone of the furnace. It took about 5 to 6 min for the specimen to reach the furnace temperature. After the specimen was held at temperature for a desired length of time the boat was pulled to the cold zone of the furnace; the heating-up period was excluded from the isothermal annealing time. Thus, the uncertainty in annealing time is higher for very short anneals, but negligible for long anneals.

Jan 1, 1969

By K. P. Gupta, P. C. Panigrahy

The stabilization of the R, a-Mn, and 0-Mn phases have been studied in the Mn-Ti-Fe and Mn-Ti-Co systems. Iron and cobalt both appear to stabilize the (Mn-Ti) R phase to almost the sarne extent. The R-phase region was found to extend from the lowest e/a to slightly beyond the maximunz e/a limit known for this phase. But, while iron appears to stabilize the a-Mn phase, cobalt tends to stabilize the p-Mn phase. In the two systems manganese appears to get replaced by iron and cobalt in each of the mentioned phases. The instability of the a-Mn phase in the Mn-Ti-Co system and the /3 -Mn phase in the Mn-Ti-Fe system cannot be explained on the basis of adverse size effects because atomic diameters for both iron and cobalt (C.N. 12 at. diam) are ziery similnr and not much different from manganese which they replace. Qualitatively, the reason for the stability of the a-Mn and the p-Mn phases can be traced to the more favorable e/a ratio prevailing in the respective systems and to a competing tendency between the two phases. In transition metal alloy systems the o, p,P, R, a- Mn,&apos; and p-Mn2 phases have been claimed as electron compounds. A large volume of work has been done to establish the criterion for the formation of the o phase but until very recently practically no systematic work was done on the a-Mn and the /3-Mn phases. A recent investigation on the P-Mn phase3 indicates the e/a criterion for p-Mn phase stabilization. Since the R phase was first known to appear only in certain ternary systems1 no detailed work was then possible for this phase. The R phase has been recently discovered as a binary intermetallic compound in the Mn-Ti~ and Mn-si~-&apos; binary systems. The existence of binary R phases opens up the possibilities of studying the effect of alloying elements on the stabilization of the R phase. Of the two binary systems possessing an R phase, the Mn-Ti system appears to be more interesting because at a suitable high temperature it is possible to find the three electron compounds, the a-Mn, p-Mn, and R phases, side by side and it is possible to study the effect of a third transition element on these three electron compounds. For the present investigation iron and cobalt, so called B elements for the formation of electron compounds, have been used as the third element to study the stabilization of the a-Mn, P-Mn, and R phases. EXPERIMENTAL PROCEDURE The alloys were prepared by using 99.9 pct pure electrolytic Fe and Mn, 99.5 pct Co, and crystal bar titanium, supplied by Semi Elements Inc., New York and Gallard Schelsinger Mfg. Co., New York. Weighed amounts of the components were melted in recrystal-lized alumina crucibles in an inert atmosphere (argon) high-frequency induction melting unit. Titanium was made into fine chips for easy dissolution and a special charging procedure was adopted to avoid contacts of titanium chips with the alumina crucibles. Up to 20 at. pct Ti, the maximum titanium content in the investigated alloys, there was no visible sign of reaction of titanium with the alumina crucibles. With a careful control of melting time and temperature the losses were minimized and were always found to be below 0.1 pct. Because of such small and almost constant weight losses, the alloys were not finally analyzed. The alloys were wrapped in molybdenum foil and annealed in evacuated and sealed silica capsules at 1000" * 2°C for 72 hr and subsequently quenched in cold tap water. Annealed samples were examined metallographically and by X-ray diffraction. For all high manganese alloys oxalic acid solutions of various concentrations and 1.0 pct HN03 solution were found suitable as etching reagents. Best contrast between the a-Mn and the R phases could be obtained by using freshly prepared 60 pct glycerine + 20 pct HN03 + 20 pct HF solution. For high iron and cobalt containing alloys, especially for alloys containing the a-Fe, y-Fe, and P-Co phases, 15 cc HNOJ + 60 cc HC1 + 15 cc acetic acid + 15 cc water solution was found to be the best etching reagent. All X-ray diffraction work was carried out (using specimens prepared from annealed powders) with a 114.6 mm diam Debye-Scherrer camera using unfiltered FeK radiation at 25 kv and 10 ma. All calculations for X-ray diffraction work were carried out using an IBM 7044 digital computer RESULTS AND DISCUSSION The two ternary systems, MnTiFe and MnTiCo, were investigated near the manganese rich end, Figs. 1 and 2, and show some common features. In both alloy systems large extensions of narrow R phase regions occur at almost constant titanium contents. At titanium contents higher than that of the single phase R-phase alloys, the same unidentified X phase was found in both ternary systems. The extensions of the X phase close to the Mn-Ti binary indicate that this phase could be the TiMns phase. Too few X phase diffraction lines were present in the diffraction patterns to make positive identification of the X phase. In contrast to this similarity the two systems show opposite behavior in the extensions of the a-Mn and 8-Mn phase regions; while iron tends to stabilize the a-Mn phase, cobalt

Jan 1, 1970