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By A. B. Kinzel

IT is with sincere appreciation and a deep sense of responsibility that I accept the honor of delivering the Howe Memorial Lecture. In our time metallurgical research has delved into phenomena ever more complex in character, so that much of this work must be done by organized teams. We, of the Union Carbide and Carbon Research Laboratories and the Electro Metallurgical Co., are truly a team and the lecture which is to follow is due largely to a few of the outstanding members of that team here cited and well known to you in their own right: John Lamont, whose work deserves special recognition, Walter Crafts, William Forgeng, William Binder, Robert Fowler, and David Swan. Thus it is a fitting pleasure to accept this honor for the research institution which I represent. Henry Marion Howe was truly inspirational and his breadth of view was such that he could use either the method of rigid proof or implied mechanism to further the broad understanding. But perhaps his outstanding characteristic was his sound judgment and his ability to weigh observations. The following is replete with observations and judgment as to their implication. We would like to feel that it is the type of dissertation which Howe, himself, would have enjoyed. Chromium carbides in stainless steel have, in a sense, been a key to the progress of the stainless steel industry. Brearley's first stainless steel, containing approximately 0.30 pct C, was successful because it differed from previous Fe-Cr alloys in that the carbides were present in such amount and form as to provide the stainless and cutting properties. Again, the 12 pct Cr die steels with a carbon content of 1.5 pct or more achieved their properties by virtue of the form and distribution of the chromium carbides. In the straight chromium steels with from 12 to 25 pct Cr, chromium carbide amount and distribution are once more the determinants with respect to physical properties and corrosion resistance. In each of the cases just cited the role of the carbide is in essence similar to that of iron carbide in carbon steel and this story is so well known that we need not dwell on it. In the austenitic stainless steels, the chromium carbide plays only a minor positive role with respect to mechanical properties, and for corrosion resistance the objective is to achieve maximum homogeneity. Thus the tendency has been to ever lower the carbon and correlatively the chromium carbide content of the austenitic steels. Today stainless steels with carbon less than 0.03 pct are beginning to find their true place, and this is being aided in major degree by our Company's introduction of a new type of ferrochromium having negligible carbon. The role of carbon in these austenitic steels has long been recognized, and a truly enormous amount of research has been devoted to the subject. Much has been learned about corrosion resistance, but the attack on the problem from the point of view of carbide precipitation and the nature thereof has left much to be understood. Because of its fundamental importance to the future of stainless steels, further study of the mechanism of carbide formation in austenitic steels is of interest. In the 18-8 Cr-Ni steels, which are typical of the austenitic class of steel, carbon is soluble in the aus-tenite to the extent of approximately 0.15 pct at 1832°F (1000°C). The solubility becomes less with decreasing temperature, so that at 1652°F (900°C) the solubility limit is approximately 0.06 pct, at 1472°F (800°C) it is about 0.03 pct, and at 932°F (500°C) it is approximately 0.01 pct. Obviously this is the ideal situation for the classical precipitation phenomenon. It is true that, in a sense, 18-8 austenite resulting from annealing is in a metastable state, and the phenomenon of ferritic precipitation from this austenite should be considered. However, in a so-called fully austenitic 18-8 steel this is secondary with respect to the carbides and can be neglected. A wide variety of chromium carbides is known. Only one chromium carbide, Cr,,C,,, frequently referred to as Cr,C, has been found in unmodified low carbon 18-8 steel. This, fortunately, simplifies the study of the precipitation phenomena. For this study an A.I.S.I. type 304 stainless steel of commercial quality was selected. Analysis showed C 0.07 pct, Mn 0.49 pct, Si 0.33 pct, Ni 9.34 pct, Cr 18.91 pct, and N 0.033 pct. As expected with this composition the steel was homogeneous after annealing except for a few pools of ferrite, which were disregarded in that their location was not coincident with subsequent carbide precipitation areas. On corrosion testing the steel was normal and typical of good commercial quality. Specimens 1/2xlx4 in. were annealed and subjected to precipitation heat treatment for 100 hr at 100" temperature intervals from 1000° to 1500°F and for ten selected intervals ranging from 5 min to 64 hr at 1300°F. These blocks were then cut into small pieces for metallographic studies. The first problem

Jan 1, 1953

By E. S. Machlin

The possibility that formation of voids under creep-rupture conditions may take place by the condensation of vacancies has been investigated theoretically. It has been concluded that nucleation of voids under creep-rupture conditions by vacancy condensation is highly improbable. However, growth of pre-existant voids by vacancy condensation is probable. A number of predictions made in this theory have been verified by the data. It has been predicted and checked that the product of rupture life and steady-state creep rate for preannealed metals and single phase alloys is an approximately invariant quantity, independent of stress, temperature, and atomic number for a given type structure. The direction of the effect of cold work on this product has been predicted and found in agreement with experiment. A number of experiments to evaluate the vacancy condensation mechanism further are described. SEVERAL papers have appeared recently which speculate on the origin of voids formed at grain boundaries under stress.' ' The object of this paper is to examine quantitatively the proposition that the voids produced in a creep test are a result of vacancy condensation. A result of this paper is a theory of creep-rupture. Void Nucleation Application of standard nucleation theory" to the problem of void nucleation leads to the following conclusions: 1—Homogeneous nucleation of voids requires a supersaturation ratio (concentration of vacancies in supersaturated to that in saturated solution) of 400 for a reasonable surface energy of 1000 erg per cm-and 1.4 for the improbably low surface energy of 10 erg per cm. 2—Heterogeneous nucleation of voids at plane interfaces between two phases requires a supersaturation ratio of 2.5 for a typical contact angle of 145 3-—Void nucleation about a solid particle may be accomplished at a supersaturation ratio of 1.17 for a typical value of work of adhesion? of 60 erg per The work of adhesion is the surface work 10 replace two solid-vauor surfaces by a solid-solid interface. enr ' between an oxide and a metal in the presence of a surface active element such as sulphur. Estimates of the supersaturation ratio at which voids are produced in diffusion experiments yield a maximum of 1.01. Inasmuch as the foregoing mechanisms of void nucleation probably will not operate at this level—too low a surface energy is required—the investigatol. is led to the conclusion that voids must already exist. That is, nucleation of voids probably does not occur. Rather, existing submicroscopic voids grow out to visible size. Already existing voids might be produced during solidification or working. Supercritical sized parlicles which contain cracks may act as heterogeneous void nuclei. Gas pockets may act as void nuclei. Experiments are desired to determine the nature of the heterogeneous void nuclei which grow out to voids in both diffusion and creep experiments. Void Growth Void growth might occur in at least two possible ways, depending upon whether the already existing void nuclei are at grain boundaries or within the grains. In the case of a spherical void far from a crystal boundary, vacancies are generated during creep as a consequence of the migration of suitable dislocation jogs' and are also annihilated at sinks. Under these conditions, a steady-state concentration of vacancies is built up in the crystal, defined by the condition that for any differential volume the rate of generation of vacancies in that volume equals the rate of annihilation of those vacancies." This equality would lead to the development of a gradient of vacancy concentration radially outward from the void surface up to a radius where the vacancy lifetime becomes equal for all directions of vacancy migration. The distance over which this vacancy concentration gradient extends equals about 2vD,T* where D, is the vacancy diffusivity and T:' the vacancy lifetime in a crystal outside the gradient in a zone of constant vacancy concentration. The vacancies generated in the region over which the gradient exists will annihilate more often at the void than elsewhere. Approximately a little over one-half the vacancies generated in the gradient zone will annihilate at the void. Hence, the growth rate of the void is given by on where R is the radius of void in centimeters, is the atomic volume, and R is the rate of generation of vacancies, number per centimeter" per second. R D and T* may be estimated in terms of other physical parameters." In particular, R = n.j e/b [3] where n is the average number of vacancy produc-

Jan 1, 1957

By T. I. Moore, L. A. Painter

THE electrolytic zinc plant of the American Zinc Co. of Illinois was described by Davidson' in 1944. Since then, improvements as well as expansion of the plant facilities have been made. In order to increase the production of high grade zinc which was needed for war purposes, an expansion program designed to double the slab zinc capacity was started in 1942 and completed in March 1943. This expansion was propagated by a contract between the American Zinc Co. of Ill. and the Defense Plant Corp. The contract included the facilities of the Fairmont City, Ill., property of the American Zinc Co., where a suspension-type roaster with contact acid plant, cadmium distillation furnace, Waelz oxide and densifying plant, and horizontal retort furnaces were installed. The expanded Monsanto, Ill., plant and the additional facilities of' the Fairmont plant were designed to integrate the metallurgical treatment of zinc concentrates for the production of special high grade zinc at Monsanto with the production of acid, cadmium, high grade zinc from furnace skimmings and the Waelz treatment of leach residue at Fairmont. In general, the original flowsheet was not changed, except for the addition of the filtering, drying and reclaiming of leach residue, and the treatment of purification cake for the recovery of copper, cadmium sponge, and zinc. Fig. 1 is a flow diagram of present operations. The original plant facilities, desi-gned for 50 tons daily production of slab zinc, had some units which were more than adequate. Therefore, in expanding the facilities to 100 tons per day, it was not necessary to double all operating components. Table I gives the comparison of the changes made in the unit operating components for the original facilities, 1941, the 1943 expansion, and the 1951 facilities. During the past 11 years a number of improvenients have been made resulting in: 1—an increase in slab production, 2—higher recoveries on the calcine treated, 3—better quality of slab zinc produced, 4—higher current efficiencies, and 5—less man hours Table I. Changes in Operating Facilities Operating Unit 1941 1943 1951 Calcine unloading (pneumatic), 10 tons per hr 12 calcine unloading track hpr. and elev., 60 tons per hr 1 Calcine storage, tons 1,000 2,000 2,000 Leach tanks, 35 vol. tons. No. 3 5 6 slurry mixing 6x6 ft stainless tank, No. 1 Ball mill. 4.5 rt x 16 in. conical. No. 1 CLassifier duplex, No. 1 1 Thickeners. 50 it diam. No. 2 9 2 Filter thickeners, sq ft '-- Moore filters, sq ft 5.760 11,520 Drum filters, 10 ft diam x 16 ft, No. 3 3 Rotary arlers, No. 1 2 1st stage Cu-As purificatlon tank. 90 vol. tons, No. 3 Solution heaters, No. 3 Filter press, 30x30 bronze, No. 4 Zinc dust purification tanks, 45 vol tons. No. 3 5 4 Filter press, 36x36 bronze, No. 3 5 3 Cadmium recovery plant: Process tank. No. 5 2 Cake roaster, 20 ft diam x 4 hearth. No. 1 Filter press, 24x24, No. 4 1 Sponge wash box, 4x6 ft, No. 1 Evaporative cooling unit (vacuum), No. 1 Purified storage tank, vol. tons 400 400 400 Cell acid storage, vol. tons 400 400 400 Electrolytic cells, No. 180 372 372 Cell room ventilation, cu ft per min 35,000 125,000 125,000 Cell cooling water. gal per min 1,500 2,300 2,300 Deep well 16 in. x 95 ft, 1500 gal per min, No. 2 3 3 Melting and casting furnace, 130 ton. No. I Furnace fume scrubber unit, No. 1 Dross drums, No. 2 Dross roaster, 8 ft diam x 8 hearth, No. 1 Electrolysis power conversion, kw 6,250 23,750 23,750 Power transformers, 13,800/440, kva 1,000 1,500 2,000 Steam boilers fire tube, 15 psi, lb per hr 12,000 18,000 18,000 Steam boilers water tube. 125 psi, Ib per hr 30,000 Air compressor, 2 stage. 300 cu ft per min. 100 psi, No. 1 2 Air compressor, 1 stage, 300 cu ft per min, 20 psi, No. 1 Vacuum pumps, 18x7, 720 cu ft per day, No. - - . Vacuum pumps, 24x11, 1.633 cu ft per day, No. 3 3 Building area, sq ft 60,854 113,568 115,000 per ton of metal produced. In the summer of 1944, the "reverse" leaching process was placed in operation and since it has been described,' no further description will be given. Other facilities and changes which have contributed to the process improvements were the scrubbing of fume from the melting and

Jan 1, 1953

By AIME AIME

THE first meeting" of the Iron and Steel Division was held Tuesday afternoon, Feb. 17, with nearly 100 men present and C. B. Murray as chairman. This was a round table discussion of iron ore beneficiation and of its relation to blast-furnace operations, and only one paper was presented, by Clyde E. Williams. Mr. Williams discussed the future of the Lake Superior district in its relation to the iron and steel industry. He stated that although estimates had shown that the total life of the present known iron ore reserves of the Lake Superior district was 20 to 30 years and that many believed that as a result of this depletion foreign ores would be brought to the Atlantic Seaboard, thus causing a displacement of the steel industry to the East, he believed that the Lake Superior district would continue to be the important producer of iron ore used in this country for decades beyond the predicted time before the exhaustion of the present known reserves. He described the three general types of ores that required concentration and pointed out the needs for research and the proper methods to be used in processes for the beneficiation of these low-grade materials. He discussed the effects of fine concentrates and the use of agglomerated ores in the blast furnace. This general summary led to interesting and valuable discussion on the various economic, ore-dressing, and blast-furnace problems in large-scale development of beneficiation. Perry Harrison described recent developments in sintering and the properties of iron ore sinter. T. B. Counselman and W. H. Coghill discussed the merits of classification of iron ore previous to tabling. Robert Faulkner and J. E. Little described the beneficiation plant of the Bethlehem Steel Corporation.

Jan 1, 1931

By E. J. Whittenberger

OVER the past two decades, the Cr-Ni stainless steels, popularly termed 18-8 steels, have been used in ever increasing amounts in the aircraft, automotive, chemical, transportation, and building industries.' The contribution of nickel to these steels is associated with its strong austenite forming tendencies which assure good hot and cold forming characteristics, and with its beneficial influence upon corrosion resistance, especially in oxidizing acids. The adjustment of the chromium and nickel contents has been used advantageously in controlling the rollability and the response of the mechanical properties of these steels to cold reduction. With the advent of the jet age and the government stockpiling of nickel, considerable emphasis has been placed upon the development of satisfactory low nickel or nickel-free substitute austenitic stainless steels, especially for nonmilitary applications.2 Fer-ritic 17 pct Cr (AISI Type 430) stainless steel has been effectively substituted for the austenitic grades in applications with mildly corrosive exposure. However, the limited formability and ferromagnetic behavior of this grade do not permit its satisfactory use for many applications. To achieve better formability in a nenmagnetic steel and conserve nickel, a number of low nickel substitute austenitic grades of stainless steel have been developed in recent years. In most of these steels, manganese and nitrogen have been substituted for a portion of the nickel content, for example, 17 pct Cr-4 pct Ni-6 pct Mn (AISI Type 201), 18 pct Cr-5 pct Ni-8 pct Mn (AISI Type 202), and 16 pct Cr-1 pct Ni-17 pct Mn, all of which contain approximately 0.15 pct N.1 The precipitation of ferrite when large ingots of these Cr-Ni-Mn-N steels are heated to temperatures high enough for conversion to slabs on conventional universal slabbing mills has seriously affected the hot workability of these steels. This has been especially true in the case of the 16 pct Cr-1 pct Ni-17 pct Mn (sometimes referred to as 15 pct Cr-1 pct Ni-15 pct Mn) composition." Because it is possible that little, if any, nickel will be available for nonmilitary stainless steel applications in the event of a national emergency, a need has existed for a nickel-free austenitic stainless steel containing hot and cold forming characteristics similar to 18-8 steels and corrosion resistance at least comparable to the 17 pct Cr ferritic steel. The aforementioned deleterious effect of ferrite upon the hot working characteristics of the 16 pct Cr-1 pct Ni-17 pct Mn, steel indicates the importance of establishing the austenite-austenite plus ferrite phase boundaries as the first step in the development of nickel-free or low nickel substitute steels. This paper presents the data collected in an elevated temperature constitutional study of the Cr-Ni-Mn-N system at the South Works of the United States Steel Corp. Experimental Procedure The 300 compositions used in the investigation were melted in a 26 Ib induction furnace and cast in 3 14 in. diam by 18 in tall hot-tcpped ingot mo1ds.

Jan 1, 1958

By E. Horton Jones

CONTENTS I Page Introduction ....:......................... 3 Chapter I. Unit Costs. . ...................... 4 Chapter II. Comparative Costs ..................... 20 Chapter III. Composite Costs..................... 32 Chapter IV. Wage Scale....................... 36 Chapter V. Raw Material Prices................. 37 Chapter VI. Description oF Costs ................... 64 Introduction We have endeavored in the following "sheets" to give the unit construction costs derived from 'the building of the Arizona Copper Co.'s new smelter, Clifton, Ariz., starting in February, 1912, and completing February, 1914.l In Chapter I—Unit Costs—are to be found the most elementary total unit costs which the accounts provide for. They are usable in the Clifton district. Here too are found the percentages to be added to an estimate for Engineering and General Expense. In Chapter II—Comparative Costs—these elementary costs have been classified, averaged and reported as labor and material unit costs. In such form the labor unit costs, when properly applied to similar conditions as these under which they were derived, are usable anywhere. The material unit costs are better disregarded for more accurate estimates and replaced by a newly priced bill of material. In Chapter III Composite Costs are given. They are unit costs built up from several elementary units, and likewise units of larger dimensions and simpler application, valuable for

Jan 1, 1915

Jan 1, 1967

By John A. Gann

THE Iron and Steel Division has shared in the general increased activity of the industry. The fall meeting at Chicago was not only well attended but particularly characterized by the virility and scope of the discussions at the various sessions. The program for the February meeting indicates a similar interest in technological developments. The Blast-Furnace Committee continues to have the hearty support of the industry with discussions of such a nature as to be of real help to the operator as well ass the technologist. The work of the Open-Hearth Committee continues to be outstanding, with ever increasing attendance and ever mounting interest in the meetings. Formation of a conference on the physical chemistry of steelmaking has given those interested in this phase of work an organization allowing the same

Jan 1, 1936

By RAYMOND M. LARSEN

The area covered by this paper is the same as that covered in 1944. Colorado, Montana, Utah, and Wyoming are included in the discussion and statistics, and brief mention is made of activities in the adjoining states, except Nebraska and New Mexico, which are considered separately (see pp. 187-189). The other adjoining states, North and South Dakota, Arizona, Nevada, and Idaho, have had virtually no oil or gas production, except in the gas field of Cedar Creek. Montana, which extends into North Dakota; the Pierre field in South Dakota, where gas is produced and separated in gas-water wells, and the abandoned Payette area in western Idaho. Seven gas wells were drilled in North Dakota, but no wildcatting was reported. Several unsuccessful wildcats were drilled in South Dakota, and some were continuing at the end of the year. Activity in Idaho and Nevada was negligible.

Jan 1, 1946

By B. Foley

IN THE first report1 disks of steel of known composition and history were exposed, under carefully prescribed conditions, to impacts of explosion products resulting from the explosion of 50-gm. charges of explosives having well-determined rates of detonation, or to impacts of a weight of known mass, falling through measured distances. There were available at the Bureau of Mines Experiment Station five explosives, having rates of detonation of 5716, 4470, 3190, 2296, and 1523 meters per second. It being desired to continue this exposure downward, and no explosives with a lesser rate of detonation being at hand, resort was had to the impact of a falling weight, the differences in the velocity of impact being attained by varying the height from which the weight was allowed to fall. It was expected by this means to demonstrate experimentally that the development of Neumann bands was determined by the suddenness of the deformation of the metal in which they were produced. But while progressive deformations were obtained with the explosive arges, the falling-weight method failed because the tup was ogressively deformed, also the disks were so flattened as to be unfit for metallographic examination. A return, therefore, has been made to the use of explosives as the source of energy, an air-gap being interposed between the charge of explosive and the disk, as the effect diminishes as the function of the distance between the charge and the object on which the products of its explosion, or the shock wave from its detonation, impinge. The details of the arrangement of the tests by the air-gap method were the same as those described by Foley and Howell,2 except that after the system had been set up and the desired apparent specific gravity for the explosive

Jan 2, 1926

By B. N. Das, H. E. LaBelle, G. A. Wolf

ZnxHg1-xTe and GUPxAS1-x crystals have been grown from solution by a traveling heater method (THM). In a floating zone type fashion a solution zone sandwich of liquidus composition is made to migrate through feed material of solidus composition by a slow-speed movement of <0.25 mm per hr of the total charge through a stationary radiation or rf heater. ZnxHg1-xTe of composition 0 < x < 0.3 and GUPxAS1-x of composition 0.5 < x < 0.7 have thus been grown from their solution in liquid tellurium and gallium, respectively. A simplified analysis of heat and material transport is given. In this method, Proper adjustment of the temperature distribution along the liquid solution zone axis eliminates instability of the crystal growth front and constitutional supercooling. 1) ZnxHg1—xTe: The solid-solution telluride was grown in sealed quartz tubes at a temperature of 500 to 650°C. For first Preparations mechanical mixtures SINGLE crystals of solid solutions of III-V and II-VI compounds have been grown by various techniques from the melt, from the gaseous phase, or from solutions in of ZnTe and HgTe were used. X-ray determinations showed the resulting crystals to be homogeneous throughout their entire length, except for the end sections. The homogeneity is greatly improved by reusing the ingots obtained from first Preparations as feed material in subsequent crystal growth runs. Ingots of up to 4.5 cm have thus been obtained. 2) GaPxAs1-x: Crystals were grown in a EN crucible at a temperature of 1050°C. For first Preparation a mechanical mixture of Gap and GaAs was cemented together by a minimum amount of liquid gallium through "wet" grinding of Gap and GaAs crystals in the liquid metal and subsequent exuding of excess gallium by compacting of the charge in the EN crucible. Electron microprobe measurements using P-Ka and As-Ka radiation showed uniformity over a length of 0.4 cm. Properties of the various Preparations are described. liquid metals.&apos;-&apos; Their growth from the melt by passing a liquid zone of pseudobinary liquidus composition through a solid rod of solidus composition is generally a convenient method of growth of single crystals of solid solutions of homogeneous composition. This method loses much of its usefulness, however, if the melting point or the vapor pressure of the solid solution is high. In the case of a compound having a high melting

Jan 1, 1969

By H. E. Brauer

PART I INTRODUCTION Among the zinc base alloys used for casting in metal moulds, particularly die casting, those alloys containing aluminum usually together with copper, are probably the most widely used. The reason lies in their desirable mechanical properties and the fluidity and absence of hot shortness which facilitate casting. Objection to the use of these alloys may bee raised on account of the rather frequent instances of failure of a type variously described as disintegration, exfoliation and intercrystalline corrosion. In all cases which have come to our knowledge where the conditions of service were known, failure has occurred upon exposure to moisture and heat varying from a hot moist atmosphere to high pressure steam. Moreover, the alloys described in all such cases contain impurities or alloying metals which the present authors have found to have enormous accelerating effects on intercrystalline oxidation in zinc-aluminum alloys. Thus Williams1 experimented with an alloy containing a large percentage of tin exposed in Cuba. Bauer2 reports failure of a boiler plug containing over 1 per cent. of impurities, doubtless lead for the most part, and Rawdon&apos;s3 examinations were made on an alloy containing 0.4 per cent. lead also exposed to the Tropics. No thorough investigation of the nature and cause of this type of failure has been made heretofore, so far as the authors are aware, although Rawdon correctly classified it as internal oxidation. We have, therefore, carried out this investigation in order to determine the exact nature of this phenomenon and the effect of other metals present as impurities or intentional additions in accelerating or-retarding it. In the appendix we include a considerable amount of data which may be valuable in selecting an alloy combining good physical properties with high resistance to oxidation.

Jan 8, 1922

While coal was mined in New Mexico in 1860, production was on a small scale during the period considered here, and Table 77 contains all the data found. [ ]

Jan 1, 1942

Aamot, Olav Crone, (Member &apos;29) Met , Det Norske Aktieselskab for Elektrokemisk, Industry, Inc , Oslo, Norway Aase, Glenn D , (Junior Member &apos;40) Met Engr , Western Research Dyv , American Smelt & Refin Co , Selby, Calif Abadye, Henry G., (Member &apos;43) Petr Engr , Sr Prod Engr , Long Beach Oil Development Co, 255 So Santa Clara Ave, Long Beach, Calif Abbott, Clarence E , (Member &apos;04) Civil Engr , &apos;P , Tennessee Coal, Iron & R R Co , 1542 Brown Marx Bldg Birmingham, Ala Abbott, Fred E , Jr , (Member &apos;44) Petr Engr , c/o Arabian American Oil Co , A P 0 816, c/o Postmaster, New York, N Y Abbott, John H , (Member &apos;46) Resident Mgr , Hickman, Williams & Co, 230 No Michigan Ave, Chicago 1, Ill Abbott, R W R., (Member &apos;47) Asst Power Supt, Cananea Consolidated Copper Co , Cananea, Son , Mexico Abbott, Robert H? Jr., (Junior Member &apos;47) Petr Engr, Producing Dept, The Texas Cc, P 0 Box 951, Houma, La Abel, Raymond L., (Member &apos;28) Chem Engr, Prof of Chem Engrg, Univ of Pittsburgh, Pitts¬burgh Pa Abel, Walter D , (Member &apos;32) Min Engr, Div of Corporations, 403 State Bldg , Los Angeles 12, Cal if

Jan 1, 1952

By E. H. Citron

CLEANING fine coal in jigs is not new in Europe, where the feldspar jig is being used almost exclusively for this purpose. A feldspar jig operates with an artificial bed made up generally of feldspar or other hard rock material. Dimensions and specific gravity of this material must be of suitable values for good bedding qualities. For good efficiency, the size should range between 1x2 in. and 2x2x1/2 in. thick.

Jan 4, 1958

By A. Hultgren, B. Herrlander

The original aim of the present investigation was to study the mechanism of cracking on hot-deforming red-short steels. During the microscopical examination of hot-deformed soft steels attention was directed to various patterns of so-called veining in the ferrite, as related to variations in the deformation procedure and the heat-treatment of the steels. Later, a hot-short steel of higher carbon content was also studied. Historical Storey1 suggested in 1914 that veining developed during the gamma-alpha transformation from the growing together of several alpha nuclei within the same gamma grain. Rawdon and Berglund2 found that in soft steel, forged somewhat below A3, the ferrite showed profuse veining, while usually there was very little veining after forging at very high temperatures. They emphasized the similarity between the pattern of veining formed by deformation about 600°C and that of slip lines formed by deformation at room temperature. They also suggested a relation between the gamma-alpha transformation and veining. They became convinced from etching and from the slip-line pattern in cold-deformed material, that the orientation within any ferrite grain showing veining was uniform. In recrystallized alpha grains veining was absent. Veining did not appear materially to affect the properties of ferrite. Independent of veining, continuous networks, attributed to preexisting delta and gamma grain boundaries were sometimes found, the former only in cast steel. Those networks were usually connected with small inclusions. Ammermann and Kornfeld3 confirmed that recrystallized ferrite grains showed no veining. In soft steel annealed between A1 and A3 there was veining only in the ferrite grains formed during cooling. Electrolytic iron deformed at 880&apos; was free from veining. Bannister and Jones&apos; considered veining in ferrite to be a manifestation of microscopical and sub-microscopica~ inclusions. NorthCOtt,5,6 in studies on veining in steel and other metals and alloys, attributed veining to oxide inclusions. It could generally be removed by annealing in hydrogen, a suggestion earlier made by Rawdon and Berg1und. TrittOn (discussion in ref. 5) emphasized the importance of perfect polishing for bringing out veining by etching, and thought the veins or subboundaries were produced during the gamma-alpha transformation, when it was rapid, as a result of the volume change and contraction during cooling. Slight distortion would not allow projecting parts of a growing crystal to join up with atomic symmetry. Hanemann and co-workers showed that veining in ferrite consisted of ridges

Jan 1, 1948

By B. Herrlander, A. Hultgren

The original aim of the present investigation was to study the mechanism of cracking on hot-deforming red-short steels. During the microscopical examination of hot-deformed soft steels attention was directed to various patterns of so-called veining in the ferrite, as related to variations in the deformation procedure and the heat-treatment of the steels. Later, a hot-short steel of higher carbon content was also studied. Historical Storey1 suggested in 1914 that veining developed during the gamma-alpha transformation from the growing together of several alpha nuclei within the same gamma grain. Rawdon and Berglund2 found that in soft steel, forged somewhat below A3, the ferrite showed profuse veining, while usually there was very little veining after forging at very high temperatures. They emphasized the similarity between the pattern of veining formed by deformation about 600°C and that of slip lines formed by deformation at room temperature. They also suggested a relation between the gamma-alpha transformation and veining. They became convinced from etching and from the slip-line pattern in cold-deformed material, that the orientation within any ferrite grain showing veining was uniform. In recrystallized alpha grains veining was absent. Veining did not appear materially to affect the properties of ferrite. Independent of veining, continuous networks, attributed to preexisting delta and gamma grain boundaries were sometimes found, the former only in cast steel. Those networks were usually connected with small inclusions. Ammermann and Kornfeld3 confirmed that recrystallized ferrite grains showed no veining. In soft steel annealed between A1 and A3 there was veining only in the ferrite grains formed during cooling. Electrolytic iron deformed at 880&apos; was free from veining. Bannister and Jones&apos; considered veining in ferrite to be a manifestation of microscopical and sub-microscopica~ inclusions. NorthCOtt,5,6 in studies on veining in steel and other metals and alloys, attributed veining to oxide inclusions. It could generally be removed by annealing in hydrogen, a suggestion earlier made by Rawdon and Berg1und. TrittOn (discussion in ref. 5) emphasized the importance of perfect polishing for bringing out veining by etching, and thought the veins or subboundaries were produced during the gamma-alpha transformation, when it was rapid, as a result of the volume change and contraction during cooling. Slight distortion would not allow projecting parts of a growing crystal to join up with atomic symmetry. Hanemann and co-workers showed that veining in ferrite consisted of ridges

Jan 1, 1948

By Robert A. Rapp, Ronald L. Pastorek

Solid-state electrochemical measurements by three alternative experimental procedures were made with the cell FeO, Fe3O4 |Zro.85Cao.15O1.85 |Cu| Zr0.85CaO.15O1.85 | FeO, Fe304 to establish the solubility and diffusivity of oxygen in solid copper in the temperature range 800" to 1030°C. The solubility of oxygen in solid copper and the diflusivity of oxygen in solid copper Dgu = 1.7 X 10-2 exp(-16,000/RT) Cm2/sec were determined and confirmed in alternative experiments. The enthalpy of solution of oxygen in solid copper equals —10,000 cal per mole; the partial excess entropy of the oxygen atoms in the Cu-O dilute solution is approximately the same as that found for interstitial atoms in other metals. The diffusivity of oxygen in solid copper is consistent with that expected for an interstitial atom. RELIABLE values for the saturation solubility N(s) and diffusivity DO of oxygen in solid copper have not been unambiguously established in the literature. Following three early determinations by others,1"3 Rhines and Mathewson4 reported that the solubility of oxygen in solid copper increased from 0.007 at. pct 0 at 600°C to about 0.015 pct at 1050°C. Phillips and skinner,, using essentiially the same analytical procedure, reported that the solid solubility increases from 0.0018 at. pct 0 at 550°C to about 0.0075 pct at 1050OC. The only previous value for the diffusivity of oxygen in solid copper was reported by Ransley.6 Ransley deoxidized Cu-Cu2O alloys in an atmosphere of carbon monoxide gas to yield a solubility-diffusivity product. He used the solubility data of Rhines and Mathewson to calculate the diffusivity values. Another method for obtaining the solubility-diffusivity product (N(s) DO) is by measuring the widths of internal-oxidation zones in copper alloys as reported by Verfurth and Rapp.7 However, the calculated N(S)Do products depend upon the alloy content of the specimen, so that the internal oxidation of copper alloys does not follow ideal internal oxidation kinetics. As a result, unequivocal values for the N(s) DO product were not obtained by this procedure. A solid-state coulometric titration technique similar to that employed in this work was introduced by C. Wagner8 to study the dependence on silver activity of the Ag/S ratio in silver sulfide in the temperature range of 160" to 300°C. Similar experiments have been carried out by C. Wagner and co-workers9-11 to study the stoichiometry range of silver and copper tellurides, cuprous sulfide, and cuprous selenide. Numerous authors have carried out electrochemical measurements with a solid oxygen-ion-conducting electrolyte to determine the solubility and/or diffusivity of dissolved oxygen in several liquid metals.12-l6 Rickert and Steiner17,18 have used solid-state electrochemical measurements to determine the diffusivity of oxygen in solid silver from 760" to 900°C. Two different cell geometries were used. In the cell of linear geometry Fe, FeO | ZrO2 + (CaO) | Ag + [0 (dissolved)] [1] oxygen diffused from the interior of the silver electrode to the silver/electrolyte interface where the oxygen activity had been lowered from a fixed initial value to practically zero by the application of voltage to the cell. The diffusivity of oxygen in solid silver was determined from the solution of the diffusion equation and the time dependence of the cell current. However, this determination of the diffusion coefficient depended upon a knowledge of the solubility of oxygen in solid silver. A cylindrical geometry was used for the cell Pt, O2(Po2 = 0.21 atm) | ZrO2 + (CaO) | Ag + [0 (dissolved)] [II] which also allowed the diffusivity of oxygen in solid silver to be determined. These values were in agreement with other available data.l9 Recently, Raleigh20,21 used a method involving the measurement of diffusion-limited currents in a cell involving the AgBr solid electrolyte to determine the diffusion coefficient of silver in Ag-Au alloys at 400°C. Diffusivity values on the order of l0-14 sq cm per sec were measured in the alloy composition range 10 to 60 at. pct Ag in a single experiment. From numerous electrical conductivity and galvanic cell measurements,9&apos;22"26 the solid solution Zr0.85 Ca0.15 O1.85 has been established as an electrolyte with predominant oxygen ion conduction over a wide range of intermediate and high oxygen activities. For interrelating the thermodynamics and the kinetics of the dissolution of oxygen in solid copper in this investigation, a galvanic cell was constructed with FeO-Fe3O4 as the reversible reference electrode, the Zr0.85Ca0.15 O1.85 electrolyte, and a pure copper specimen under-saturated in oxygen as the other electrode. THEORETICAL ANALYSIS Three variations of a high-temperature electrochemical technique were used in this study to provide two determinations each of the solubility and diffusivity

Jan 1, 1970

By R. W. Regehr, R. A. Burmeister

Epitaxial growth of GaAs 1-xPx on germanium substrates was achieved using an open tube vapor transport system. The compositional range of 0.3 < x < 0.4 was examined. The best results were obtained with (311) orientation of the germanium substrate. The physical and chemical properties of the resulting layers were investigated using several techniques. Spectrographic analyses of the layers indicate substantial incorporation of germanium into the GaAs t-X Px layer. Evidence is presented which indicates that this incorporation occurs via a vapor phase transport process rather than by solid phase dijfu-sion. Electrical measurements suggest that the germanium thus incorporated behaves predominantly as a deep donor in the compositional range of 0.33 < x * 0.40 and has a deleterious effect upon the luminescent properties of GaAs1-x Px. The increasing technological importance of GaAs1-xPx for use in light-emitting devices has led to an evaluation of several aspects of existing growth processes. The method most commonly used to prepare GaAs1-xPx for electroluminescent device applications is vapor phase epitaxial growth on GaAs substrates.&apos;-4 In a typical electroluminescent diode structure the active region of the diode is entirely within the epitaxial layer and thus the electrical properties of the substrate are relatively unimportant since it is effectively a simple series resistance (assuming hetero-junction effects to be negligible). The use of germanium rather than GaAs as the substrate material is of interest for several reasons. First, GaAs of reasonable structural quality has been epitaxially grown on germanium4-2 and it is reasonable to expect that GaAs1-xPx could subsequently be deposited on the GaAs layer. Second, germanium substrates are readily available with both lower dislocation densities and larger areas than GaAs. Finally, single crystals of germanium are more economical than GaAs single crystals. The principal objective of the present investigation was to test the feasibility of growing GaAs1-xPx epi-taxially on germanium substrates, and to evaluate the properties of such layers with regard to electroluminescent device requirements. The approach used was to a) demonstrate epitaxial growth of GaAs1-xPx on germanium, and b) characterize the relevant structural, electrical, and optical properties of the GaAs1-xPx layers. The possibility of germanium incorporation into the grown layers was of special interest since there was some indication of this in previous studies of GaAs growth on germanium.5&apos;11,12 Although a study of the electrical properties of germanium in GaAs1-xPx was not an intent of this investigation, several features of the electrical properties of the layers grown in the present study which appear to be due to germanium are described. EXPERIMENTAL PROCEDURE The open-tube vapor transport system used for the epitaxial growth of GaAs1-xPx is illustrated in Fig. 1. This system utilizes the GaC1-GaC13 transport reaction and is similar in most respects to the larger system described elsewhere.&apos; The germanium substrates were n-type, with a resistivity of 40 ohm-cm (Eagle-Picher Co.). These were cut to the orientations of {100), {111), and (3111, and were mechanically polished and chemically etched in CP-4 (5 min at 0°C) prior to growth. In some cases, a GaAs substrate was employed in addition to the germanium. The orientation of the latter was {loo}, and they were also mechanically polished and chemically etched prior to growth. The initial composition of the deposited layer was pure GaAs. After approximately 10 microns of GaAs was deposited on the germanium substrate, the phosphorus content of the layer was gradually increased over a distance of approximately 15 microns to the desired concentration and maintained at this value throughout the remainder of the growth. Typical operating parameters used during growth are given in Table I. Selenium was used as a n-type dopant in several runs to facilitate comparison of the electrical properties of the layers grown on germanium with those of layers grown on GaAs substrates, which are usually doped with selenium. The concentration of H2Se in the gas phase was adjusted to a value which would normally yield a carrier density of 1 to 5 x 101 7 at room temperature in layers grown on GaAs substrates. The terminal surfaces of the epitaxial layers were examined by optical microscopy for structural characteristics. Laue back-reflection photographs (Cu radi-ation) were also made on the terminal surface to verify the epitaxial nature of the deposit. After these steps

Jan 1, 1970

By A. M. Gsudin, F. W. Bloecher, C. S. Chan-s, P. L. De Bruyn

M ANY elements can now be obtained in radioactive form. The radioisotopes have the same chemical properties as the corresponding inactive forms, differing from them only by their nuclear instability. This combination of properties permits radioelements to act as atomic tracers, the tracer atoms revealing themselves at the time of their decay with emission of electrons and positrons (beta radiation), helium nuclei (alpha radiation), or electromagnetic radiation (gamma radiation). The tracer technique is particularly attractive for the study of surface chemistry problems in general and flotation problems in particular because of its extreme sensitivity. A large proportion of the ionizing events, initiated by the characteristic decay emissions in a radioactive sample, can be recorded by radiation probes of suitable design and geometry. An idea of this sensitivity can be gathered from the following example: assume that 0.001 mg of pure RaD (radiolead), a beta-ray emitter, is used in a flota¬tion tracer experiment. This is 1210 X 10-1 mole; since the fraction trans¬formed per second (the so-called decay constant), is 9.92 X 10-10, the number of ionizing events per second would be (6.023 X 1023 X y,,, X 10-1) X 9.92 X 10-10 = 2.845 X 106. (The decay constant is expressed in sec-&apos;. Its product with the half life, expressed in seconds, is the pure number, natural logarithm of 2 = 0.693. The half life is the length of time required for decay of one half of the total number of radioactive atoms of one particular type.) (See reference (1) at the end of this article.)

Jan 1, 1948