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By R. W. Diamond

THREE mills are now operated by the Consolidated Mining & Smelting Co. of Canada, Ltd.: (1) The Sullivan Concentrator, Chapman Camp, B. C. (near Kimberley, B. C.), (2) the, St. Eugene Concentrator, Moyie, B. C., and (3) the. Tadanac (Customs) Concentrator,Trail, B. C. Differential flotation is the only concentration method employed in all three plants. Tables are used for pilot work only. Values in all ores treated are as minerals of some or all of the metals -lead, zinc, gold or silver. Occasional customs ores contain small amounts of copper. Sulfide iron content of ores treated varies from a FIG. 1.-THE SULLIVAN CONCENTRATOR, CHAPMAN CAMP, B. C. (NEAR KIMBERLEY, B. C.). low figure in some ores to 38 per cent. iron in Sullivan ore. Gangue minerals vary considerably. A few customs ores are somewhat oxidized and as a result present difficulties. Flotation equipment and control are similar in the three mills, and the flow sheets at the St. Eugene and Tadanac mills, while much simpler, are modeled after that of the Sullivan Concentrator. In this paper I will not detail treatment nor equipment in any of our plants but will describe special features in, or in connection with, all of them. I will then deal with flotation reagents, describing our use of each of them, particularly where differing from general practice. I will conclude with comments on special subjects related to differential-flotation.

Jan 1, 1927

By David J. Mack, A. H. Kasberg

The transformation characteristics are found to resemble a similar binary alloy. The differences are due to the alpha iron particles. While strength properties of the isothermally transformed alloys are high, ductility is low, resulting in high notch sensitivity. The elastic modulus can be varied widely and correlates with strength. Abnormal grain growth sometimes occurs. A CONSIDERABLE amount of work has been done on the isothermal transformation of high purity aluminum bronzes using the methods pioneered by Davenport and Bain.' The first isothermal study of these alloys was made by Smith and Lind-lief,' but more recently Mack", " and Klier and Grymko5 have continued the work. In commercial aluminum bronzes, one or more elements are always added to refine the ß grain size and to improve the mechanical properties. The effect of these elements on the isothermal transformation has not been reported in the literature. No systematic study of the properties of isother-mally transformed aluminum bronzes has apparently been reported, although adequate information is available on their properties after more conventional heat treatments. The fact that a eutectoid alloy will reject face-centered cubic a during isothermal transformation at appropriate temperatures2'" warrants investigation of the properties of such structures. Another cogent reason for determination of properties is the fact that the modulus of elasticity of the eutectoid alloy varies widely with the nature of the heat treatment." Preliminary experiments7 indicated that a high purity eutectoid alloy would not give reproducible results in property determinations because of ex- cessive ß grain growth during heat treatment, a single ß grain frequently occupying the entire cross-section of a standard test bar. For this reason a commercial aluminum bronze containing iron as a grain growth inhibitor was used. The isothermal transformation diagram of the alloy was first determined. On this basis, the heat treatments necessary to produce representative structures were selected and the properties determined. Material A commercial alloy (Ampco Grade 20) of the following composition was used: Cu, 83.68 pct; Al, 12.31 pct; Fe, 3.68 pct; Ni, 0.28 pct; and Mn, 0.03 pct. This alloy proved to be very close to the eutectoid composition and may be compared with the accepted value of 11.8 pct A1 in the binary Cu-A1 system. Comparison is valid only because of the "negative replacement effect" of the iron? Since traces of excess y² are present, the alloy is slightly hyper-eutectoid. The material was received in the form of 3/4 in. round, hot extruded bars. Specimens of approximately ? in. thickness were cut from this stock and drilled to allow suspension in the salt bath. The structure of the as-received material varied from 5 to 40 pct a + y2* in a matrix of ß. The amount of a + y² present was dependent upon the ß grain size of the extruded material. The coarse grained portions did not form as much a + Y, as did

Jan 1, 1952

By Henry G. Petrow, Edward S. Porter

MAJOR occurrences of lignite with significant uranium concentrations have been reported in western parts of the Dakotas, especially in Harding County, S.D., and Billings and Slope counties, N.D. Samples containing as much as 2 pct U3O8 were received at the Winchester Laboratory, but those deposits reportedly of major significance have generally contained 2 to 4 lb of uranium per ton of lignite. Most of the uranium appears to be associated with the carbon in the lignite as a urano-organic complex, although in most instances analysis of inorganic material has shown uranium values. Molybdenum is similarly associated with the carbon, and its concentrations generally parallel the uranium. The noncarboniferous material is mainly composed of gypsum, jarosite, bentonite, sandstone, and miscellaneous carbonates and silicates. The heating value of these lignites has been determined as about 3500 Btu on a dry lignite basis. The ash content of lignite has varied widely, the important samples containing 30 to 60 pct ash on a dry lignite basis. Because of the significant upgrading of uranium values obtained by ashing the lignite, and because recovery of uranium from raw lignite presented formidable processing problems, major emphasis at Winchester has been placed on devising methods to recover uranium from lignite calcines. Proper roasting removes the problem of organic fouling of solutions, reduces the size of the leaching plant for an equivalent uranium throughout, and generally provides equal or better acid leaching efficiency with less acid than can be obtained from the raw lignite.

Jan 9, 1957

By W. O. Philbrook, R. S. Kaplan

Rates of oxidation were studied for levitation-melted samples of Fe-C, Fe-Si, Fe-C-Si, Fe-C-P, Fe-C-Mn, and Fe-C-S alloys oxidized by mixtures of either 5 or 10 pct O2, in helium at several temperatures in the range 1850o to 2050°C. The concentrations of the oxidized elements decreased linearly with time. Mass transport in the gas phase was found to be the rate-controlling step. Ilowever, a mulliple -film model was necessary to account for the flux of iron and manganese vapor counter to the oxygen flow. THE advent of the basic oxygen and spray-refining processes of steelmaking has generated considerable interest in the kinetics of reactions between oxygen-bearing gases and liquid iron alloys. The object of this research was the determination of the role of mass transport in the gas and metal phases and of surface chemical reactions in the overall oxidation process. Of the reactions for elimination of elements dissolved in liquid iron, the oxidation of carbon has been investigated most extensively. Studies by Fuji and Ura,1 Fuji,' and Gunji et a1.,3 done with pilot-scale basic oxygen furnaces, indicated that mass transport of gaseous oxygen was rate controlling when the metal carbon content was greater than 0.15 to 0.30 pct. When the carbon content was less than this critical amount, diffusion of carbon dissolved in the iron was thought to be rate limiting. L. A. Baker et a1.4,5 levitated samples of iron-carbon alloys and oxidized them in CO-CO2, CO-C02-He, and O2-He atmospheres at 1660°C. These investigators found that the metal carbon concentration decreased linearly with time and their samples exploded when the carbon content fell below about 0.8 pct. They concluded that the rate-limiting step was the transfer of oxidant to the gas-metal interface. A model based on this rate-limiting step predicted decarburization rates within a factor of two of those experimentally observed. The explosion of the samples was attributed to a shift in rate-controlling step from gas-phase mass transport to transport of carbon dissolved in the metal. L. A. Baker and R. G. ward6 and R. Baker7 used freely falling droplets to investigate oxidation kinetics. In both papers, the authors speculated that flame fronts may have been partially responsible for the lack of better agreement between experiment and theory. Swisher and Turkdogan8 oxidized liquid iron having an initial carbon content of 0.97 pct with CO-CO2 and CO-C02-Ar mixtures at 1580°C. They concluded that a surface reaction was the rate controlling step. Iron-silicon alloys have been oxidized by R. Baker7 and Filipov and Martynov.9 R. Baker 7 also studied iron-carbon-silicon alloys falling as liquid droplets through oxygen. He found that the rate of both carbon and silicon oxidation varied inversely with the drop radius. When the initial silicon concentration was greater than 6 pct, the samples cooled during their fall. Manganese oxidation has been studied by R. Baker,7 and by Filipov and Martynov.9 ward10 investigated the loss of manganese from liquid iron under vacuum. He found that mass transport of manganese vapor was rate limiting in argon atmospheres of 0.02 torr to 700 torr pressure. Iron-sulfur alloys were desulfurized with H2-He and H2-Ar mixtures by Peters et al.11 A model based on the transport of H2S from the gas-metal interface into the bulk gas gave fair agreement between the predictions of the model and the experimental results. EXPERIMENTAL TECHNIQUE Essentially, the technique consisted of using helium-oxygen gas mixtures to oxidize levitated samples of liquid binary and ternary iron-base alloys. The fused silica reaction tube is shown in Fig. 1. The coil design is similar to that of L. A. Baker et al.4'5 A Tef-lon-sheathed, ferritic stainless steel lifter (B) and an alumina pedestal (C) were used to support the sample in the coil prior to levitation and for the recovery of the sample. Samples were inserted into

Jan 1, 1970

By J. P. Morgan, J. R. Dawson

A review of Jones & Laughlin's Vesta coal preparation plant flowsheet in 1959 indicated that, of the three sizes making up the company's metallurgical product, the -48-mesh fraction offered the greatest potential for improvement. Efforts to improve quality and yield were directed toward the use of froth flotation. After laboratory and plant tests were completed, it was decided that treatment of the entire -@-mesh fraction of the coal mix by flotation was a significant step forward in providing better coal quality and recovery. Reduction of ash in the metallurgical coal from 6.75 to 6.25% provided material savings at the coal washer, coke plant and blast furnace. The Vesta preparation plant of the Jones & Laughlin Steel Corp. is located on the Monongahela River, eight miles south of Brownsville, Pa. This plant is capable of cleaning 2400 tph of a mixture of Pittsburgh seam high volatile coals from the Vesta and Shannopin mines to produce a metallurgical coal of 6.75% ash and a steam coal containing 14% ash. Heavy media vessels are utilized to upgrade the coarse coal while launders and tables handle the -1/4-in. fraction of the raw coal feed. Cleaned -1/4-in. coal goes to boot classifiers where the 48 mesh material settles and is removed by bucket elevators. This 1/4x48 mesh material is then sent to centrifugal dryers for dewatering. As designed, the -48 mesh fraction overflows the boot as a 7 to 8% solids slurry and is recovered by cyclones and a thickener and then dewatered by filtration. The products of the heavy media circuit, the launder-table-boot-dryer circuit and the cyclone-thickener-filter circuit, are combined to constitute the shipped metallurgical coal. The Vesta-Shannopin mix is combined at Jones & Laughlin's Pittsburgh and Aliquippa Works with purchased low and medium volatile coals for the production of metallurgical coke. A review of the preparation plant flowsheet in 1959 indicated that of the three sizes making up the metallurgical product, the -48-mesh fraction, containing 16.7% ash, offered the greatest potential for improvement of coal quality. This fraction was considered low, averaging 55%, while the rejects averaged only 33% ash. It was evident that while the boot overflow fraction represented only 6% of the product mix, a reduction in ash to 6 or 7% would reduce the ash of the metallurgical coal by 0.5%. The original plant flowsheet handling the boot overflow (Fig. 1) consisted of ten 14-in. cyclones treating 70% of the overflow and producing a 48x150-mesh underflow and a -150-mesh overflow. The cyclone overflow was combined with the remaining 30% of the boot overflow to constitute the feed to a 150-ft thickener. The cyclone underflow and the 48x325-mesh thickener underflow were fed to vacuum disk filters for dewatering prior to combination with the coarse metallurgical coal. The thickener overflow stream containing the -325-mesh material was split, 40% being recycled directly to the plant water system and 60% being flocculated and clarified in an 85-ft thickener. The underflow of this thickener was pumped to waste. Efforts to improve quality and yield on the -48-mesh coal were directed toward the use of froth flotation early in the investigation. A review of the fine coal flowsheet indicated that three streams should be considered for treatment by froth flotation. These were the -325-mesh thickener overflow, the -150-mesh cyclone overflow and the -48-mesh boot overflow. Flotation of the thickener overflow was considered basically to improve recovery by upgrading the reject stream so that it could be fed to the filters along with the thickener and cyclone underflows. The overflow from the thickener contained 3.7% solids and had an ash analysis of 33%. Flotation of the cyclone overflow and boot overflow were considered to upgrade coal quality as well as yield. Treatment of the cyclone overflow, a 5.4% solids slurry averaging 27% ash, would eliminate the coal thickener and provide a feed of flotation concentrate and cyclone underflow to the filters. Treatment of the boot overflow would eliminate both the cyclones and thickener and feed only flotation concentrates to the filters. Flotation feed in this

Jan 1, 1962

By Wittenau, E.

OPERATION of a pilot mill of 100 tons' daily capacity during 1930 and 1931 proved that the copper minerals of the Colorado and Clay sections of the Morenci ore body could be successfully concentrated by the flotation process. Economic conditions caused postponement of further testing work until late in 1938 when the former No. 6 Concentrator at Morenci was converted into an experimental mill of 1500 tons' daily capacity which was known as the Test Concentrator. The Test Concentrator began continuous operation on April 3. 1939, using such crushing and grinding equipment of the No. 6 Concentrator as was appropriate. Modern classifiers and certain flotation machines were installed and later other types of flotation machines, a large grate-type grinding mill. mechanical screens, a short-head cone crusher, and a disk filter were added to the test equipment. Operation was uninterrupted until Dec. 31, 1941.

Jan 1, 1942

By J. Parke Channing

WE are accustomed to think that we are efficient in the United States, particularly with respect to such things as mining and manufacturing. The conduct of the war has demanded in England and in France a complete readjustment of manufacturing methods and plans, and today England is probably as efficient a country as there is in the world, not even excepting Germany. This is all the more remarkable because it has been notorious for years that England has been inefficient in her manufacturing and the country has been flooded with things "Made in Germany." Today England is almost a socialistic community and the State is doing almost everything. England is now in such a position that practically everyone in the country is engaged in industry necessary for the conduct of the war, and this has been accomplished by increasing the efficiency both of her tools and of her man power. In the United States we certainly have been efficient so far as machines and perhaps so far as methods have been concerned, but we have not been efficient iii the utilization of our man power. Before the war, our labor was undoubtedly far more efficient than that of England, but it certainly was not so highly efficient as it should have been, and the problem that confronts us today, and will all the more confront us after the war, is to make our man power efficient. England has had a taste of what you may call State Socialism, and her laboring men are not going to be content to return to the old order of things. There is one feature of the labor problem in England which has permitted her to reach this condition of state socialism with comparative rapidity; this is that practically all of her laborers are English; she has little or no foreign population. While an Englishman may he a strong union man and ready to fight his employer tooth and nail, at heart he is still a British subject, and when his country was in danger he rose to the occasion.

Jan 5, 1918

By Charles Will Wright

ITALY is by- far the poorest in mineral resources of the so-called great pou7ers of Europe. Before the World War this shortage was not so serious as the essential minerals that could not be mined domestically at a profit were bought in the world markets. In those days tourist trade and remittances from Italians abroad helped to provide Italy with foreign exchange. During the World War, when certain mineral products were scarce and when prices rose in some instances to several times normal, the Italian Government began to regulate consumption and imports. It was not until later, when the Fascist regime established the corporate state, that Italy seriously considered adopting a of self-sufficiency. Private industry was grouped under separate corporations, for the principal industries; including mining and metallurgy and the heads of the large producing companies, became members of these corporations. Industrialists were thus able to regulate their own affairs and to advise the Government as to the measures necessary to meet domestic or foreign trade demands upon their industry. Labor was or-

Jan 1, 1939

By W. Crafts, J. L. Lamont

Electrolytic extraction of carbides from quenched and tempered steel and their examination under the electron microscope were found to be helpful in understanding the mechanism of secondary hardening in alloy steels1 and the same technique has been applied to isothermally transformed steels. A preliminary survey of the utility of the method has indicated that it has promising possibilities for investigating the characteristics of pearlite and bainite. The examination of a few carbon and chromium steels has suggested that pearlite is formed with carbides of two varieties representing the lamellar and nonlamellar forms and that bainite appears to form with a structure analogous to, but significantly different from, martensite. Procedure The method used for examining the carbides was similar to that described previously.1 The chromium steels used in the investigation were made from an Armco iron base in an induction furnace and were not treated with grain-refining deoxidizers. The composition of the steels is given in Table 1. Three-inch square ingots were forged and rolled to 1/8-in. thick sheet for use in the isothermal studies. After rolling, the sheets were sandblasted to remove the rolling scale. Specimens approximately 3/4-in. wide and l1/2-in. long were prepared from the 1/8-in. sheet. They were heated in a salt "bath at 2100°F (1150°C) for 1/2 hour, transferred to other salt baths at 1300°F (704°C), 1000°F (538°C), 800°F (426°C), or 600°F (315°C), held for various periods up to 100 hr and finally quenched in water. After quenching, the specimens were cut in half, one half being submitted to microscopic examination and the other half being used for the electrolytic extraction. The isothermally transformed specimens were submitted to examination under the light microscope to determine the degree of transformation and type of structure and were photographed at 2000 X. Carbide residues were obtained by electrolyzing in 10 pet hydrochloric acid and collecting the residue in glycerine to minimize attack by the acid. As pointed out in the earlier paper, the carbides are attacked to some degree during electrolysis. This attack is relatively minor on the coarser carbides but may be quite severe on the fine bainitic carbides. The residues were washed free of the glycerine with water and alcohol. The alcohol was removed by repeated washings of amyl acetate, and the amyl acetate suspensions were transferred to stoppered vials. Debye-Sherrer X ray diffraction patterns of the residues were prepared using a chromium target. A selection based on the microstructure and X ray data was then made of samples considered most illustrative of the progress of transformation. Electrolytically extracted residues from these samples were submitted to examination by W. D. Forgeng and A. C. Jenkins on the electron microscope at the Research Laboratories of The Linde Air Products Co. and electron micrographs were made at X 5000 and enlarged to 25,000 X. The micrographs obtained with both the light and the electron microscopes are mounted side by side in the accompanying figures. Although no effort was made to establish the T-T-T diagram for the chromium steels, the times for holding in the salt bath were adjusted with the intention of making the samples represent different stages in the transformation at temperatures of 600, 800, 1000, and 1300°F. The degrees of transformation and results of X ray examination are shown in Tables 2—4. The carbide designations refer to orthorhombic or trigonal carbide types and are not meant to imply that the compositions are exactly as indicated. X ray Diffraction Data The occurrence of Fe3C and Cr7C3 in these steels was found to be consistent with published data on comparable chromium steels. Steels A and B with relatively low ratios of carbon to chromium contained Cr7C3 in the pearlitic structures and FeuC in the

Jan 1, 1950

By L. S. Darken

IT has been known for sometime that in an ionic lattice, such as that of Ag2S or FeO, the migration velocity of the anion may differ markedly from that of the cation, the cation being usually the more mobile. It has been shown, for example, that during the formation of silver sulfide tarnish on silver, the tarnish layer grows principally by the motion of silver cations through an essentially stationary silver sulfide lattice. Similarly, the growth of wüstite scale on iron proceeds principally by motion of iron ions, this being accomplished with the aid of vacant iron lattice sites into which the iron may readily move. Likewise investigations of solid solutions of various salts show that there is usually considerable difference in the mobility of the different cations. Measurements of the mobility of various metals in liquid mercury indicate that the various metals move at different rates relative to the mercury. In view of the foregoing it seems reasonable to suppose that when any single phase solution is acted upon by a field, the atoms of different elements respond in different ways; in particular, that the force arising from a particular concentration gradient in a binary alloy would cause the atoms of one component to move with a drift velocity which is in general different from that of the atoms of the other component. The recent experiments of Smigelskas and Kirkendall1 support this viewpoint, since they found, in essence, that during diffusion in alpha brass zinc moves appreciably faster than copper. This was evidenced by the relative motion of inert molybdenum wires. Such a phenomenon is incompatible with the place interchange theory of diffusion proposed by several authors, most recently by Birchenall and Mehl.2 This theory has, however, also been questioned3.4 for other reasons, principally because of the high energy required for direct atomic interchange. The rather large difference between self-diffusivity, as determined by tracer technique, and mutual diffusivity, as determined with the aid of another element at low concentration, might also be interpreted as favoring the above idea; however, Birchenall and Mehl have pointed out that there is another possible interpretation of this difference; this is discussed later. A rather interesting and pertinent observation was made by Schwartz,5 who noticed that tungsten lamp filaments, heated by direct current, commonly fail by necking down in the region at the temperature gradient near the positive end, with a concomitant swelling at the corresponding position near the negative end. He introduces considerable auxiliary evidence to show that this phenomenon is occasioned by the motion of tungsten cations which migrate in the high temperature region from the positive end toward the negative end. The mechanism of this migration is not clearly understood. If, however, this experimental evidence be accepted, we are forced to conclude that alloying elements in the tungsten may migrate by the same or a similar mechanism, whatever it may be, at a greater or lesser rate, thus giving rise

Jan 1, 1948

By W. D. Forgeng, P. F. Wieser

MnSi alloys containing from 2 to 24 at. pct Si have been investigated by metallographic and X-ray methods. Contrary to published data, the temperature of the ß-manganese to a-manganese transformation is lowered by the addition of silicon. The existence of two additional intermediate phases E and ( is demonstrated. E, ranging in composition from 12 to 15.5 at. pct Si, forms by a peritectoid reaction between ß manganese and the second new phase, (. The latter phase is homogeneous at silicon contents from 16.2 to 18 at. pct and forms peritectically from the 0 phase and liquid. The eutectic reaction occurs between the ( phase and MnsSi rather than the ß phase and Mn3Si. As a result of this investigation, a modification of the manganese-rich section of the MnSi equilihrium diagram is proposed. The binary diagram of Mn-Si, according to Hansen,' Fig. 1, exhibits in the low-silicon region (below 24 at. pct Si) a peritectic transformation at 900°C of 0 manganese + Mn3Si to a manganese as determined by R. Vogel and H. Bedarff2 as well as a thermal effect at 980°C between 16.2 and 22.6 at. pct Si reported by the same authors. The nature of the latter transformation was unknown. K. Amark, B. Boren, and A. westgren3 reported indications that possibly four more phases existed in the range of 10 to 27 at. pct Si. The present investigation deals with the phase relations in the range of 2.1 to 17.7 at. pct Si. A few spot checks were conducted up to 23.84 at. pct Si. EXPERIMENTAL PROCEDURE Small experimental heats were prepared by induction melting in a vacuum furnace using argon as a protective atmosphere. The crucible was of commercial-quality magnesia. The furnace charge consisted of silicon metal and dehydrogenated electrolytic manganese. The typical analysis for these materials is shown in Table I. The molten alloy was tapped into 2-in.-square, 7-in.-long ingots. After visual inspection, a small quantity of alloy in the middle portion of the ingot was separated for use in this investigation.MnSi alloys containing from 2 to 24 at. pct Si have been investigated by metallographic and X-ray methods. Contrary to published data, the temperature of the ß-manganese to a-manganese transformation is lowered by the addition of silicon. The existence of two additional intermediate phases E and ( is demonstrated. E, ranging in composition from 12 to 15.5 at. pct Si, forms by a peritectoid reaction between ß manganese and the second new phase, (. The latter phase is homogeneous at silicon contents from 16.2 to 18 at. pct and forms peritectically from the 0 phase and liquid. The eutectic reaction occurs between the ( phase and MnsSi rather than the ß phase and Mn3Si. As a result of this investigation, a modification of the manganese-rich section of the MnSi equilihrium diagram is proposed. The binary diagram of Mn-Si, according to Hansen,' Fig. 1, exhibits in the low-silicon region (below 24 at. pct Si) a peritectic transformation at 900°C of 0 manganese + Mn3Si to a manganese as determined by R. Vogel and H. Bedarff2 as well as a thermal effect at 980°C between 16.2 and 22.6 at. pct Si reported by the same authors. The nature of the latter transformation was unknown. K. Amark, B. Boren, and A. westgren3 reported indications that possibly four more phases existed in the range of 10 to 27 at. pct Si. The present investigation deals with the phase relations in the range of 2.1 to 17.7 at. pct Si. A few spot checks were conducted up to 23.84 at. pct Si. EXPERIMENTAL PROCEDURE Small experimental heats were prepared by induction melting in a vacuum furnace using argon as a protective atmosphere. The crucible was of commercial-quality magnesia. The furnace charge consisted of silicon metal and dehydrogenated electrolytic manganese. The typical analysis for these materials is shown in Table I. The molten alloy was tapped into 2-in.-square, 7-in.-long ingots. After visual inspection, a small quantity of alloy in the middle portion of the ingot was separated for use in this investigation.

Jan 1, 1964

By J. B. Umpleby

The function of gas in the development and production of oil has far-reaching consequences that should be emphasized. The technical aspects of the subject have recently had a great deal of attention but less thought has been given to fundamental concepts underlying oil and gas law and the many customs and practices that have grown up in the industry. I shall review briefly the advance in knowledge of the subject and attempt to relate these concepts to legislative, judicial and operating problems. Early writers on oil and gas in Pennsylvania noted their close association in oil pools and their contemporaneous decline in output; also, that as gas pressure decreases so does production, that new offset wells reduce the output of existing wells, and even that tubing a well lengthens its flowing life. Expanding gas was deduced as the principal expulsive force. These fundamental observations had been made at least as early as 1880.1 They were treated, however, as academic considerations and did not greatly influence law or practice for many years except to instill into all operators an urgent desire to hasten offset drilling in order to obtain as much of the flush as possible. The most noteworthy later advance in knowledge came in 1926 when Beecher and Parkhurst published their paper on Effects of Dissolved Gas upon the Viscosity and Surface Tension of OiL.² This paper added to the old idea of an expelling force the concept of a modifying agent that actually changes the physical characteristics of the oil and facilitates its movement through the pore spaces of the reservoir rock. Largely since the publication of this paper, intensive study and concerted effort have been given to means of using to a maximum, in producing oil, these combined properties of gas. This is shown clearly by the number of papers dealing with the subject, and the new and revived terms now widely used to express thought, such as the expressions gas energy, reservoir pressure, reservoir energy, gas-oil ratio, bottom-hole pressure, beaning, tapered tubing and back-pressure. These terms, some of which are

Jan 1, 1932

By S. E. Haszko

Fused alumina or silica crucibles were used as the containing vessel. X-ray powder photographs were taken with CrKa radiation and the use of Straumanis type Norelco cameras of 114.6 mm diam. Crystallographic data for these compounds are shown in Table I. These results are in agreement with others of this series reported previousiy by Wernick and Geller.l The abnormally large atomic volume of Eu and Yb metal3 is manifested in EuA12 and YbAl2, when one compares them with the compounds reported in Ref. 1. However, the behavior of Yb in the Ni compound is the same as that observed in Ni5Yb having the Cu5Ca structure.4 Thus, it appears that Yb has a valence closer to 3 only in the Ni compounds; i.e.. the 2s and Id electrons are lost to the conduction band. 1,4 Europium in EuAl2, appears to have a valence closer to 2. In previous work.' X-ray patterns of ingots of NdCo2 indicated a multiphase material with faint indication of the presence of a phase having the MgCu2 structure. It was thought that NdCo2 formed peri-tectically from the melt, resulting in a multiphase system when cooled under nonequilibrium conditions. An X-ray pattern of a sample from a melt cooled more slowly than in past work showed that the major phase was NdCo2. The lattice constant, Table 1, is in excellent agreement with the value one can predict from the data of Ref 1. One of the other phases was identified as Co,Nd. The author wishes to thank J. H. Wernick for his guidance in this work, and D. Dorsi for the preparation of these compounds. 'J. H. Wernick and S. Geller: AIME Trans., in press. 'J. H. Wernick and S. E. Haszko: to be published. 'F. H. Spedding and A. II. Daane: Progress in Nuclear Energy, Series 5, Eds.: H. M. Finniston and J. P. Howe, McGraw-Hill Book Co., 1956. 'S. E. Haszko: to be published. 'J. H. Wernick and S. Geller: unpublished work. 6J. 11. Wernick and S. Geller: Acla Cryst., 1959, vol. 12, p. '162. The Influence of Hydrogen on the Lattice Parameters of Ti-Sn Alloys A. Coucoulas and H. Margolin A reinvestigation of the lattice parameters of Ti-Sn alloys by the present authors has revealed consistently higher c and a values than those obtained by Worrier. The difference is most pronounced in the c-parameters which are, therefore, used as a basis of comparison. This discrepancy was not due to air contamination, since precautions were taken to exclude this possibility. Specimens for Debye-Scherrer photograms were prepared by arc-melting 15-g buttons of Bureau of Mines titanium and 99.99 pct Sn, with compositions of 6, 8, 10, and 12 at. pct Sn. These buttons were forged under an oxygen-acetylene torch to 1/8-in. sq rods to produce relatively fine grained material. Scale and some surface were ground off, and the rods, after wrapping in molybdenum and titanium sheet, were vacuum annealed 15 min at 800 °C to remove hydrogen. Heat treatment of one of the samples at 875°C to produce ß and reveal a at the surface showed no surface contamination. Following the vacuum treatment, the specimens were wrapped in molybdenum and annealed for 1week at 800°C under two atmospheres of argon. After water-quenching, the rods were ground so that the thickness tapered to 0.1 mm at the tips. The ground specimens were then recapsuled as previously and reannealed 1 hr at 800°C to remove cold work. The

Jan 1, 1961

By H. Margolin, A. Coucoulas

Fused alumina or silica crucibles were used as the containing vessel. X-ray powder photographs were taken with CrKcll radiation and the use of Straumanis type Norelco cameras of 114.6 mm diam. Crystallographic data for these compounds are shown in Table I. These results are in agreement with others of this series reported previousiy by Wernick and Geller.l The abnormally large atomic volume of Eu and Yb metal3 is manifested in EuAl2, and YbAl2, when one compares them with the compounds reported in Ref. 1. However, the behavior of Yb in the Ni compound is the same as that observed in Ni5Yb having the Cu5Ca structure.4 Thus, it appears that Yb has a valence closer to 3 only in the Ni compounds; i.e.. the 2s and Id electrons are lost to the conduction band.1,4 Europium in EuAl2 appears to have a valence closer to 2. In previous work.' X-ray patterns of ingots of NdCo2 indicated a multiphase material with faint indication of the presence of a phase having the MgCu2 structure. It was thought that NdCo2 formed peri-tectically from the melt, resulting in a multiphase system when cooled under nonequilibrium conditions. An X-ray pattern of a sample from a melt cooled more slowly than in past work showed that the major phase was NdCo2. The lattice constant, Table 1, is in excellent agreement with the value one can predict from the data of Ref. 1. One of the other phases was identified as Co5Nd. The author wishes to thank J. H. Wernick for his guidance in this work, and D. Dorsi for the preparation of these compounds. 'J. H. Wernick and S. Geller: AIME Trans., in press. 'J. H. Wernick and S. E. Haszko: to be published. 'F. H. Spedding and A. II. Daane: Progress in Nuclear Energy, Series 5, Eds.: H. M. Finniston and J. P. Howe, McGraw-Hill Book Co., 1956. 'S. E. Haszko: to be published. 'J. H. Wernick and S. Geller: unpublished work. 6J. 11. Wernick and S. Geller: Acia Cryst., 1959, vol. 12, p. '162. A reinvestigation of the lattice parameters of Ti-Sn alloys by the present authors has revealed consistently hiher c and a values than those obtained by Worrier. The difference is most pronounced in the c-parameters which are, therefore, used as a basis of comparison. This discrepancy was not due to air contamination, since precautions were taken to exclude this possibility. Specimens for Debye-Scherrer photograms were prepared by arc-melting 15-g buttons of Bureau of Mines titanium and 99.99 pct Sn, with compositions of 6, 8, 10, and 12 at. pct Sn. These buttons were forged under an oxygen-acetylene torch to 1/8-in. sq rods to produce relatively fine grained material. Scale and some surface were ground off, and the rods, after wrapping in molybdenum and titanium sheet, were vacuum annealed 15 min at 800 °C to remove hydrogen. Heat treatment of one of the samples at 875°C to produce ß and reveal a at the surface showed no surface contamination. Following the vacuum treatment, the specimens were wrapped in molybdenum and annealed for 1 week at 800°C under two atmospheres of argon. After water-quenching, the rods were ground so that the thickness tapered to 0.1 mm at the tips. The ground specimens were then recapsuled as previously and reannealed 1 hr at 800°C to remove cold work. The

Jan 1, 1961

By J. W. Fredrickson

VERY little information is available in the literature concerning the solubility of magnesium in titanium. Aust and Pidgeon&apos; report the solubility of titanium in magnesium to be 0.0025 pct at 650°C and 0.015 pct at 850°C. X-ray diffraction studies failed to detect any evidence of solid solubility. Based on the fact that the addition of titanium to magnesium increases the c/a ratio of magnesium, Busk&apos; suggested that the Mg-Ti system, like the Mg-Zr, is a peritectic. Eisenreich- eported the solubility of titanium in magnesium to be 0.003 pct at 655°C and 0.064 pct at 760°C. Removal of hydrogen from the melt increased the solubility of titanium to 0.115 pct at 800°C. Eisenreich proposes a peritectic-type diagram for the high magnesium side of the system. Recent work&apos; with Ti-Mg diffusion couples indicate that the solid solubility of titanium in magnesium is extremely small and that there is limited solid solubility of magnesium in titanium. Experimental work on the Ti-Mg system is complicated by the fact that the boiling point of magnesium is considerably below the melting point of titanium. Consequently, the usual alloying methods are impractical. Alloying was accomplished in this investigation by sheath-rolling compressed compacts." The alloys were made from titanium sponge and high purity magnesium. Chemical analysis of the sponge showed 0.04 pct Mg, <0.005 pct Al, <0.001 pct Cu, 0.05 pct Fe, 0.022 pct Mn, <0.0025 pct Ni, 0.007 pct Si, <0.005 pct V, and <0.01 pct Zn. All heat treating was conducted in a protective atmosphere by encapsulating the specimens in fused silica ampules under a partial pressure of helium. Equilibrium was achieved by holding the specimens at temperature for periods of 24 to 120 hr. The alloys were hot rolled, cold rolled, and annealed prior to solution heat treatment. Fig. 1 shows the constitutional diagram obtained for the Ti-Mg system up to 1.5 pct Mg. Oxygen was known to be present in the alloys; therefore, it is necessary to consider the alloys as ternary rather than binary and the diagram should be viewed as a plot of a section through the Ti-Mg-0 system. Other investigatorso have shown that the addition of oxy- gen to titanium stabilizes the a phase. The data presented in Fig. 1 indicate that the addition of magnesium to titanium tends to stabilize the p phase. Armour Research Foundation7 has shown that the addition of magnesium to zirconium lowers the a transus and is classified as a p stabilizer in that system. Magnesium is soluble in solid titanium to the extent of at least 1.5 pct in both the a and /3 phases. No intermetallic compounds or magnesium-rich phases were observed either by light microscopy or X-ray diffraction. However, this does not negate the possibility that such phases exist at higher alloy concentrations.

Jan 1, 1956

By John S. Alexander

By Don Masson

MUCH has been written and many prophecies made on the subject of plastics as a replacement for metal, and the extent to which these materials will compete with each other for peace- time markets. (Metal is used in a broad sense to include all types.) But although certain synthetic resin materials have made considerable headway in war production, this change to plastics can hardly be regarded as a serious threat to the mining and metallurgical industries. Even today, the total yearly tonnage of plastic materials is approximately only 1 per cent of that of steel, and only a small fraction of this might be considered as actually replacing steel. However, these figures must not be taken as indicating that the importance of plastics in fields hitherto considered exclusively the territory of metals is negligible-far from it. Considering the comparatively minute tonnage of plastics now being produced in comparison to the output of metals, the number of applications where plastics have augmented or aided the sale of metals is considerable.

Jan 1, 1944

By R. E. Old

THROUGH the use of pressure and production records, formation properties and bottom-hole sample data, the performance of an oil reservoir may be studied analytically to define and evaluate the natural forces acting on and within the reservoir. Use is made of material-balance methods and the equations for natural water encroachment to estimate reserves. Methods for calculation of pressures are given. A study of the Jones sand, Schuler field, Arkansas, has been made and is given here to illustrate the value of the methods discussed.

Jan 1, 1942

By J. R. Long, C. E. Armantrout, A. H. Roberson, T. R. Graham, R. S. Dean

THE general program of the Federal Bureau of Mines on the study of alloys made with electrolytic manganese has been extended to copper-manganese-aluminum alloys. The initial results of the work are presented in this report and concern a determination of the mutual solubilities of these three metals to form the alpha solid solution. The principal studies of this system, reported in the literature, have been focused on the Heusler alloys containing 10 to 30 per cent Mn with 9 to 15 per cent Al. These investigations concern the ferromagnetic properties and the determination of the crystallization contours, with little emphasis on the phases present in the solid state outside of this area. Heusler1 recognized a breakdown of the beta solid solution of the magnetic alloy, which he says decomposes into the alpha, beta, and gamma phases of the copper-aluminum system, somewhat modified by the presence of manganese. Rosenhain and Lantsberry2 examined alloys ranging up to 10 per cent Mn and io per cent Al but were unable to reach definite conclusions concerning the phase boundaries in that area. Krings and Ostmann3 made magnetic studies and examined the system by thermal analysis. They present contours of the liquidus surface but do not place any of the phase boundaries. Corson4 [ ] has drawn a hypothetical boundary for the alpha field with few data to support such construction. The binary systems involved have received considerable attention by numerous investigators, and their general features are well established. The hopper-aluminum system (Fig, I) has been thoroughly examined by Stockdale,5 who places the limit of solubility of aluminum in copper at 7.4 per cent at the liquidus and 9.5 per cent at room temperature. Smiths established the eutectoid temperature at 570°C. by a study of isothermal decomposition of the beta phase. The copper-manganese system (Fig. 2) has been recently reviewed by Dean and co-workers.7 Using high-purity alloys, they place the solubility of alpha man-

Jan 1, 1946

By I. J. Hegyi, J. J. Tietjen, H. Nelson, J. I. Pankove

The fabrication of p-n junctions in GaAsl-,P, alloys by a vapor-phase gowth technique has for the first tirne resulted in room-temperature injection lasers capable of operating over a broad range of wavelengths extending into the visible region of the spectrum. The shortest wavelength achieved to date is 6750A at room tetnperature. In addition, at 78°K the threshold current density values for these lasers are generally the lo~vest reported, and the emitted radiation extends to the lozc,est wavelength ever attained (6350A). With lasers fabricated from material containing 14 pct Gap, quanticm efficiencies of 26 pct and peak power outputs of 25 zu were obtained at room temperature. ALTHOUGH room-temperature operation of GaAs injection lasers has been well-documented,&apos;-5 the operation of GaAsl-,P, (x > 0) laser diodes has been restricted to relatively low temperatures.8-&apos;0 This has been previously attributed5, 7, 10-12 partially to the difficulty of preparing single-crystalline GaAsl-& alloys having a high degree of chemical homogeneity and purity. Also, with these materials it has been difficult to prepare high-quality, abrupt p-12 junctions by diffusion techniques; and, in turn, this has made it difficult to obtain optimum electrical properties for room-temperature operationL3 in the resulting laser diodes. As a result, GaAsl-,P, laser diodes have not been efficient enough to permit operation at room temperature. For example, using diffused structures, only a few diodes were obtained which could be operated even close to room temperature (255K)." Recently, a vapor-phase growth method of preparing epitaxial deposits of GaAsl - .P, alloys has been described,14 and the high-purity and homogeneity of these materials has been previously demonstrated. Of special significance, with this technique, n- or p -type doping can be initiated or discontinued at any time and at almost any rate during the crystal growth so that the donor and the acceptor concentrations can be easily controlled to obtain desired impurity profiles. This permits high-quality, abrupt p-TZ junctions to be vapor-phase grown entirely during the crystal growth process, so that diffusion or other p-n junction fabrication processes are unnecessary. After growth, the device does not have to be heated to elevated temperatures, which avoids the possible unwanted introduction and motion of both impurities and lattice defects. Using the vapor-phase growth method cited above, over 300 room-temperature injection lasers have been prepared from GaAsl-,P, alloys having compositions in the range of 0 5 x 5 0.41. These lasers have emitted cohe~ent radiation in the spectral range of 8350 to 6350A at 78°K or from 9000 to 6750A at room temperature. The threshold current densities of the best lasers are independent of the alloy composition over the range 0 < x < 0.2 and compare favorably with values for good GaAs lasers.&apos; MATERIAL PREPARATION Multilayer, epitaxial deposits of GaAsl-,P, alloys are prepared by a vapor-growth technique described elsewhere.14 With this technique, the individual layers which comprise the multilayer structure are prepared sequentially in the deposition apparatus without interrupting the crystal growth. The epitaxial layers are deposited on GaAs substrate surfaces oriented normal to the ( 100) direction. The substrate wafers employed in this study were usually doped with tellurium to an electron concentration of approximately 2 x 10" per cu cm. To avoid strains, the first 10 to 15 p of the deposited material is uniformly graded from pure GaAs to the specific GaAsl- ,P, alloy composition of interest. The GaAsl - ,P, alloy growth is then continued to form a layer of constant composition having a thickness in the range of 25 to 75 p. Both the graded region and the layer of constant composition are doped with selenium to an electron concentration of about 2 x 10" per cu cm. The p region of the diode is then incorporated in the crystal by abruptly changing the dopant concentrations in the vapor phase to facilitate doping with zinc. This layer has a hole concentration of approximately 3 X 1019 per cu cm and typically is 50 p thick. DEVICE FABRICATION In general, the GaAs substrate and the region of graded composition are removed. Ohmic contacts are made to the n-type side by tin evaporation and to the p-type side by an electrodeless nickel deposition. This is followed by an electrodeless deposition of gold on both sides. The crystal wafer is then cleaved along (110) planes and sawed into rectangular parallelepipeds. Typical dimensions are 100 by 300 for the junction area and 100 µ for the diode height. The diodes are either soldered to a copper stud or pressure-mounted in a copper clip. RESULTS AND DISCUSSION Approximately 400 laser diodes having compositions in the range of 0.41 have been prepared by the method described above. Each laser was routinely tested at liquid-nitrogen temperature. The lasers were operated with l-psec current pulses at a repetition rate of 60 pulses per sec. The parameters of greatest practical interest are the photon energy or wavelength of the laser output and the threshold current density. Fig. 1 shows the variation of photon energy with alloy composition at 78°K. The composition was determined from the lattice constant of the material obtained by X-ray back-reflection measurements. Although there

Jan 1, 1968