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By E. R. Funk, H. Udin, J. Wulff

The surface tension of solid silver is measured by a refinement of the Udin, Shaler, and Wulff technique. The tests were made in a purified helium atmosphere at four different temperatures and various times. The average of the results is y = 1140 + 90 dynes per cm for the range 875° to 932°C. A SUCCESSFUL measurement of the surface tension of solid copper was reported by Udin, Shaler, and Wulff.' They performed experiments on copper held in vacuum within 100°C of the melting point. At these temperatures, the creep rate in 5 and 3 mil copper wires was sufficient to permit determination of a balance between the contractile force of surface tension and the extending force of small copper weights joined to the end of the wire specimens. By suspending a series of variously weighted specimens and measuring the change in a gage length during the test, the balancing applied load could be found. This point of zero strain is a measure of the surface tension. The mathematical formulation of this gives: w = r y [1] where w is the weight for zero strain; r, the radius of the wire; and y, the surface tension. A subsequent note by Udin pointed out that eq 1 applies only to single-crystal specimens. Since the specimens used revealed a "bamboo" structure upon

Jan 1, 1952

By W. L. Fink, R. S. Archer

This paper describes results obtained on aluminum-beryllium alloys and aluminum-beryllium-copper alloys in the preparation of which aluminum of 99.95 per cent. purity was used. The constitution and structure of the high purity binary alloys are given. The aluminum-beryllium eutectic occurs at 0.87 per cent. beryllium and 645° C. The solid solubility of beryllium in aluminum at 631° C. is 0.05 per cent, beryllium, and at room temperature is less than 0.013 per cent. beryllium. The densities of the alloys are substantially the same as the values calculated on the assumption that the densities of the elements are not changed by alloying. All of the alloys age-harden at room temperature after a solution heat treatment, and the chill cast aluminum-beryllium-copper alloys age-harden at room temperature without a prior solution heat treatment. Hardness and tensile properties are given. Salt spray corrosion tests on sheet specimens are described. It is shown that beryllium docs not lower the corrosion resistance of aluminum to salt spray provided the alloy has received a solution heat treatment. Commercial Possibilities of Beryllium Alloys Beryllium has certain properties such as low density, high hardness, good corrosion resistance, low thermal expansion and fairly good electrical conductivity, which would indicate that it is worthy of consideration as an ingredient of light alloys. It was not until recently, however, that beryllium could be produced with sufficient ease to encourage research along this line. The production of beryllium by electrolysis from a fluoride bath has now developed to the stage where it could be used commercially if there were sufficient demand for the element. Claims made in certain newspapers and technical articles in regard to thc strength and corrosion resistance of thc aluminum-beryllium alloys are not justified by experimental results. Moreover, beryllium is an expensive metal ($200 a pound at the time this work was done) and, on account of the nature of its ore and the consequent difficulty of its extraction, it will in all probability continue to be a relatively expensive metal, although the cost would undoubtedly be a great deal less than $200 a

By L. R. Gremillion, G. H. McClellan

Chemical and mineralogical data on 25 apatitic phosphate ores of diverse origins from commerical deposits around the world were widely variable. From 14 to 18 chemical constituents of each sample were determined, and each sample contained four to ten mineral phases. Altogether there were 34 different mineral phases, many of which occurred in only one sample. Samples contained varieties of francolites, fluorapatites, and mixed hydroxyfluorapatites. The apatite compositions calculated from chemical data showed that certain types of substitutions were characteristic of different apatite varieties, but every variety had some degree of substitution of (CO3)2- for (PO4)3-, and the P205 contents of the apatites ranged from 35.1 %for a North Carolina apatite to 41.9% for a Missouri apatite. In a concentrate with a grade of 32% P2O5, the North Carolina sample should contain about 8.7% gangue and the Missouri sample about 23.7% gangue. To optimize the digestion process of a phosphate ore, it is necessary to consider the dissolution behaviors of different mineral phases during acidulation and their possible effects on chemical properties of a fertilizer intermediate. In the production of wet-process acids or superphosphates, acidulation of certain minerals may cause foaming and formation of sludge, filtration, corrosion, and handling problems.

Jan 1, 1982

By D. M. Dunbar, H. C. SCHLILTZ

HE Chile Exploration Co.'s mine and reduction plant are at Chuquicamata, Chile, on the eastern edge of the Atacama Desert, 163 miles northeast of Antofagasta, 80 miles from the Pacific Ocean, and 70 miles from the Bolivian border and at an altitude of 10,000 ft. Guggenheim Brothers started prospecting in -4pri1, 1912, and purchased the main property, in October of that year; the property is now controlled by the Anaconda Copper Mining Co., which acquired it in January, 1923. The orebody strikes north 10" east, and has a dip of about 64" west. The length, as far as developed, is approximately 1% miles; the width averages sightless than mile.

Jan 1, 1925

STANDING COMMITTEES. Executive. CHARLES F. RAND, Chairman. JAMES F. KEMP, JOSEPH W. RICHARDS, ALBERT R. LEDOUX, BENJAMIN B. THAYER. Membership. BENJAMIN B. THAYER, Chairman. WALTER R. INGALLS, JOHN H. JANEWAY, JR., JOHN D. IRVING, HERMAN A. PROSSER. Finance. EDWARD L. YOUNG, Chairman. KARL EILERS, JAMES GAYLEY. Library. JAMES F. KEMP, Chairman. JOHN HAYS HAMMOND,3 E. GYBBON SPILSBURY,2 ALEX. C. HUMPHREYS,4 BRADLEY STOUGHTON.

Jan 9, 1913

STANDING COMMITTEES. Executive. CHARLES F. BAND, Chairman. JAMES F. KEMP, JOSEPH W. RICHARDS, ALBERT R. LEDOUX, BENJAMIN B. THAYER. Membership. BENJAMIN B. THAYER, Chairman. WALTER R. INGALLS, JOHN H. JANEWAY, JR., JOHN D. IRVING, HERMAN A. PROSSER. Finance. EDWARD L. YOUNG, Chairman. KARL FILERS, JAMES GAYLEY. Library. JAM ES, F. KEMP, Chairman. JOHN HAYS HAMMOND,3 E. GYBBON SPILSBURY,2 ALEX. C. HUMPHREYS,4 BRADLEY STOUGHTON.

Jan 8, 1913

By P. R. Bradley

NO thought had been given to deep level mining at the Alaska Juneau mine prior to 1930, but in that year a prospect winze was started and continued for 1000 ft. vertically below the main haulage or adit level. The point at which this winze was started was selected entirely by the prevailing physical conditions, such as proximity to established workings, convenience, and the presence overhead of a substantial pillar of dimensions sufficient to lo¬cate the sheave wheel and loading pocket. It was, therefore, with some surprise that the unusually high-grade ore found just below the collar of the winze persisted, with certain vari¬ations, to a full depth of 1000 ft. vertically. Likewise this unusually high-grade ore was found to persist on the various crosscuts driven on levels at 200-ft. vertical intervals. To be brief, this deep level work found richer ore than any yet found in the mine over like areas, and inasmuch as the area was sufficient to plan for extensive deep level mining, plans for such work were given extended consideration.

Jan 1, 1936

By J. E. McConnell

No. 1 Electric Furnace and Billet Casting Shop, located in Plant No. 4 adjacent to our new 12 inch merchant bar mill, represents Inland Steel Company's first venture into both electric furnace steelmaking and continuous casting. Construction of the shop began in September, 1968 with excavation and foundation work, followed by awarding of general construction and major equipment contracts from December, 1968 to April, 1969. The first heat was tapped from No. 90 furnace on December 18, 1970, and the first cast attempted on December 29, 1970. The rated shop capacity is 500,000 tons per year of semi-finished billets and blooms, and was justified on the basis of approximately 85 percent of the product mix falling into the special bar quality category of steel grades. Auxiliary equipment has been designed into the shop to pursue free machining steels, including leaded and tellurized grades. The start-up and operation of the shop for the first year was essentially based on a two-part philosophy: 1. Initially operate with the "open stream" casting technique to accomplish machine shakedown and operator familiarization and training. 2. Operate with special bar quality, or "submerged casting", practices, make necessary refinements and gain customer acceptance

Jan 1, 1972

By S. L. Couling, D. J-P. Adenis

i2) twinning in poly crystalline magnesium does not usually occur homogeneously but is concentrated in localized regions or bands inclined at about 45 deg to the stress axis. These bands of twinning are in many ways similar to the Luders bands seen in mild steel alzd Al-Mg alloys. Within a hand the percentage of the volume twinned may approach 100 with no twinning outside the band. Bands are not produced in specimens of grain size a3 x 10-= in. or in samples deformed at temperatures 300°F. Annealing of heavily twinned regions leads to re crystallization textures markedly different from those found in the starting material. The exact preferred orientation appears to depend on the starting grain size, the amount of compres-sive strain and the manner in which it is introduced. MAGNESIUM, with its hexagonal crystal structure, deforms plastically at room temperature by a variety of deformation mechanisms, one of the most important of which is (1012) twinning. In single crystals such twinning occurs on compression parallel to the basal plane or tension perpendicular to it. In wrought polycrystalline magnesium, such as rolled sheet, the preferred orientation is such that almost all the grains have their basal planes nearly parallel to the sheet surface. Therefore, if sheet material is deformed by cold rolling or by tension in the plane of the sheet, deformation of the grains by (1012) twinning is not favored. On the other hand, twins form readily on compression parallel to the plane of the sheet as in a compression test, on the compression side of a bent sample, or at the outer rim of a drawn cup. The first part of this paper deals with the morphology of (1012) twinning in polycrystalline magnesium. Some aspects of this subject were covered by Barrett and ~aller.' Of direct bearing on the work to be reported here are the comments of Jetter' in discussion to Ref. 1. He shows that Mg + 1.5 pct Mn alloy sheet when bent about a transverse axis develops a pattern of transverse bands on the surface of compression. A longitudinal section through the sheet shows that the bands are wedge-shaped and inclined to the rolling direction. The bands were identified as "outcrops of regions in which deformation is localized as the result of the reorientation brought about by (1012) twinning". The second part of the paper treats the recrys-tallization texture of the specimens twinned by compression. Magnesium deformed and then annealed generally recrystallizes with no change in orientation.1'3'4 This statement is true provided that the amount of (1012) twinning occurring during deformation is small. In an isolated observation, Anse15 has shown that, when deformation produces much (1072) twinning, annealing causes new maxima to appear in the pole figure. The second part of the paper elaborates on this point. I) MATERIALS AND PROCEDURES All of the experimental work to be reported was done using the commercial sheet and plate alloy AZ31B (nominal composition: 3 pct Al, 1 pct Zn, 0.4 pct Mn, balance magnesium). This alloy is available commercially in two tempers, -0 (fully annealed) and -H24 (partially annealed), and specimens of both tempers were used in various experiments. The average grain. diameter (AGD) for -0 temper 0.125-in. sheet is about 0.5 x 10- in.; where material of larger grain size was desired, anneals at elevated temperature were used to grow the grains. compression tests were made on a screw-driven "hard" machine. The specimens were clamped in a

Jan 1, 1964

By R. A. Vandermeer, J. C. Ogle

The preferred crystallographic orientations (texture) developed in randomly oriented, poly crystalline niobium during rolling were studied by means of X-ray diflraction techniques. The evolution of texture at both the surface and center regions of the rolled strip was carefully examined as a function of increasing defamation throughout the range 43 to 99.5 pct reduction in thickness. Certain aspects of the center texture development in niobium are in agreement with the predictions of a theory by Dillamore and Roberts, but others cannot be explained by the theory in its present form. Above 87 pct reduction by rolling, a distinctly different texture appeared in the surface layers which was unlike the center texture. The present results are compared with previous results obtained from other bcc metals and alloys. RANDOMLY oriented, poly crystalline metal aggregates when plastically deformed to a sufficiently large extent develop preferred orientations or textures. In a recent review article, Dillamore and Roberts1 pointed out that the nature of the developed texture may be influenced by a large number of variables. These include both material variables such as crystal structure and composition and treatment variables such as stress system, amount of deformation, deformation temperature, strain rate, prior thermal-mechanical history, and so forth. From a practical point of view, the control of preferred orientation may often be important for the successful fabrication of metals into usable components. During the past few decades many experiments have been devoted to the study of deformation textures. This work, however, has been confined in large part to metals and alloys that have an fcc crystal lattice. By comparison, bcc metals and alloys have received much less attention, and consequently our understanding of preferred orientations in these materials is only shallow. This state of affairs worsens when it is realized that almost all of our present howledge about this class of materials derives from studies on irons and steels.' The bcc refractory metals, which are relative newcomers to the industrial world, have, on the other hand, been given at best only passing glances in the area of texture development. Our understanding of the evolution of preferred orientations in bcc metals can only remain fairly limited until systematic studies of metals and alloys other than the irons and steels have been carried out and the influence of the many variables has been determined. To that end a program was initiated to investigate in detail texture development in niobium. The present paper reports some of the results of this study. Textures were determined at both the center and surface of strips rolled variously to as much as 99.5 pct reduction in thickness at subzero temperatures. Emphasis in this paper is on texture description and on texture evolution during rolling to progressively heavier deformation. EXPERIMENTAL PROCEDURE The niobium was purchased from the Wah Chang Corp. as a 3-in.-diam electron-beam-melted billet. Chemical analysis indicated the impurities to be less than 300 ppm Ta, 40 ppm C, 10 ppm H, 170 ppm 0, and 110 ppm N. All other impurities were below the limits of detection by spectrochemical analysis. This large-grained billet was fabricated into specimen stock so that a fine-grained randomly oriented grain structure resulted. This was accomplished in three deformation steps alternated with recrystalli-zation anneals of 1 hr at 1200°C in a vacuum of low 10"6 Torr range after each deformation step. The first step was to alternately compress the billet 10 to 20 pct in each of three orthogonal directions. The second step was to compress in only two directions 90 deg apart to produce a 2-in.-sq bar. The final step was to roll this bar 50 pct to give a 1-in. by 2-in. cross section. After the final anneal, metallo-graphic examination showed the material to have an average grain size equivalent to ASTM No. 5 at 100 times (i.e., 0.065 in. diam). Specimens cut from the center and edges of this bar gave no indication of detectable preferred orientation when examined by X-ray diffraction. Samples 1.5 in. long, either 0.625 or 0.750 in. wide, and approximately 0.400 in. thick were machined from this fabricated ingot. The surfaces corresponding to the rolling planes were ground so as to be parallel. The samples were chemically polished in a solution of 60 pct nitric acid and 40 pct hydrofluoric acid (48 pct solution) prior to rolling to remove any cold work introduced in the machining operations. Rolling was accomplished with a 2-high hand-operated laboratory rolling mill that had 2.72-in.-diam rolls. Prior to operation, the rolls were polished with 600 grit paper, cleaned with acetone, and then soaked in a container of liquid nitrogen for several hours. The samples were also soaked in liquid nitrogen prior to rolling and were recooled between each pass. While some slight heating of the samples occurred during rolling, this procedure maintained the sample temperature well below 0°C at all times. The samples were rolled unidirectionally, and the rolling plane surfaces were not inverted during any phase of the operation. The draft per pass averaged between 0.010 to 0.012 in. After 96 or 97 pct reduction the draft was reduced to 0.001 to 0.002 in. per pass. Samples were rolled to various reductions in thickness between 43 and 99.5 pct.

Jan 1, 1969

By V. A. Phillips

Transmission-electron micrographs of electro-thinned samples of bulk-aged Cu-3.1 pet Co alloy show an aging sequence, supersaturated solid solution — coherent particles — quasi -coherent particles — noncoherent particles. Hardening is due to precipitation of coherent spherical fee coball-rich particles showing coherency strain fields, which are resolved at between 15 and 30A diameter. Loss of- full coherency did not occur until well into the overaged region, even with the assistance of deformation after aging. Different average particle diameters of 123, 92, and 149 ± 10Å were observed in samples aged to peak yield strength at 600°, 650°, and 700°C, respectively, indicating that there is no critical size for peak hardening. Noncoherent particles tended to develop (111) faces and became octahedral in shape. Dislocations tended to nucleate spherical coherent particles which eventually grew together forming large elongated particles. The surface energy of a noncoherent (low-angle) inter-phase boundary is estimated to he about 50 ergs per sq cm. A number of particle lining-up phenomena were observed. Overaging is principally attributed to increase in particle spacing, progressive loss of coherency, and increase in amount of discontinzdous precipitation. COPPER dissolves about 5.6 at. pet (5.2 wt pet) of cobalt at 1110oC1 and the solubility decreases to 0.75 at. petl (0.54 at. pet)2 at 650°C and to 0.1 at. pet or less at lower temperature.' It has been known for many years3-5 that Cu-Co alloys are capable of age hardening. Since cobalt is fee above 417°C and its atom size is only about 2 pet smaller than that of copper, precipitation of coherent particles would be expected. The equilibrium phase precipitated at 700°C and below contains about 10 pet Cu in solution which tends to stabilize the fee structure, lowering the transformation temperature to 340oc.l The alloy is known to undergo discontinuous precipitation in addition to general precipitation; while the former can be seen with an optical microscope, the latter precipitates are not visible except in the grosly overaged condition.5, 6 Extensive use has therefore been made of the ferromagnetic properties of the precipitate in order to follow the course of aging, and it has proved possible to measure the average particle size, spacing, approximate shape, and volume fraction and to determine that the particles are coherent without ever seeing a particle (see for example Refs. 2, 7, and 8). The magnetic measurements of particle size are limited to diameters below about 120Å.7 The present study was undertaken using the techniques of transmission-electron microscopy in order to check the above conclusions, to extend the previous magnetic work to larger particle sizes, and to attempt a more detailed correlation of properties and structure. A portion of this work has already been published.9-11 The present paper is concerned with the metallographic features of precipitation in relation to aging curves. Bonar and Kelly12'13 have published preliminary results of a similar study on single crystals of Cu-2 at. pet Co. EXPERIMENTAL Preparation of Alloy. A Cu-Co alloy, containing 3.12 wt pet (3.36 at. pet) Co by analysis, was prepared from 99.999 pet purity oxygen-free copper and electrolytic-grade cobalt. The alloy was melted and cast in vacuo in a high-frequency furnace using a graphite crucible and mold: Analysis showed chat 0.004 pet C was picked up during melting. The 1-1/2-lb ingot was homogenized in hydrogen for 24 hr at 1000°C. Slices were cold-rolled to 0.005 or 0.003 in. thickness, with an intermediate 650°C anneal in hydrogen at 0.080 in. thickness. Batches of six to ten strips were solution-treated in sealed-off quartz tubes in high vacuum in a vertical furnace and quenched by dropping into iced brine containing a device which snapped off the nose of the tube. Solution treatment consisted of 1 hr at 990°C or 2 hr at 965°C. The latter was employed for all mechanical-property studies, since a tendency was noted for the higher temperature to give porous material. Strips were usually aged individually in a horizontal vacuum furnace, inserting into the hot zone and withdrawing into a cold zone without breaking the vacuum. This method gave a rapid heating rate, permitting the use of short aging times. In some cases, particularly for the longer aging times at the higher temperatures, samples were sealed individually in quartz tubes in high

Jan 1, 1964

By Rex B. McLellan, W. A. Oates, William W. Dunn

Vapor transport experiments have been carried out in order to determine the saturation solubility of carbon in cobalt and nickel with respect to graphite over a large temperature range. Some of the uncertainty existing in the thermodynamics of these systems has been removed and accurate estimates of the heat and entropy of solution of carbon in nickel and cobalt have been deduced. The relatitle partial entropies of carbon are considered to be predominantly vibrational in origin. ALTHOUGH the thermodynamic properties of solid solutions of carbon in cobalt and in nickel have been investigated previously, the results of these investigations show a marked lack of agreement amongst themselves. In this investigation simple vapor transport experiments have been undertaken in order to measure the temperature variation of the saturation solubility of carbon in nickel and cobalt with the object of removing to some degree the uncertainties in the thermodynamics of these systems. In the case of the C-Co system, early measurements of the saturation solubility of graphite in Co show results at the same temperature differing by almost an order of magnitude.1-5 Recently Schenck, Frohberg, and Jaspert6 measured the equilibrium between CH4-H2 mixtures and C-Co solid solutions at five temperatures ranging from 850° to 1050°C. Their values of the activity of carbon in the solid solution, deduced from a knowledge of the chemical potential of carbon atoms in the gas, varies linearly with composition at each temperature and the authors concluded that the solutions were Henryan. The saturation solubilities deduced from Schenck, Frohberg, and Jaspert's measurements are in good agreement with the solubility determined at 1000°C by Rao and Nicholson7 and the value at 1000°C deduced from the measurements of smith8 on the equilibrium between CO/CO2 mixtures and C-Co solutions. The object of this investigation was to check the results of Schenck, Frohberg, and Jaspert and also to extend the range of solubility data to higher temperatures. In the case of the C-Ni system, there is fairly good agreement between the values given by schaefer9 for the saturation solubility of carbon C, in nickel, the value given at 1000°C by Rao and Nicholson,7 and that deduced from Smith's8 measurements of the equilibrium between C-Ni solutions and CO/CO2 mixtures at 1000°C. The values of C, deduced from Schenck, Frohberg, and Jaspert's6 measurements of the carbon activity in nickel are also in fairly good agreement. However, there is a large discrepancy between the above set of measurements and those of Lander, Kern, and Beach," who measured the carbon solubility of nickel in equilibrium with graphite by depositing an excess of graphite on the surface of Ni-foils from a gaseous mixture and subsequently diffusing in the carbon. These results are considerably higher than those mentioned above. The determinations6 of carbon activity a, in nickel from the CH4-H2 equilibrium showed that at all the temperatures investigated, the ac-values varied linearly with concentration up to about a, = 0.6 and then began to increase with concentration much more rapidly. On the basis of auxiliary experiments in which C-Ni solutions were equilibrated with graphite, Schenck, Frohberg, and Jaspert concluded that their CH4-H2 solid equilibrations at high values of a, were spurious and a, behaved linearly with composition up to the saturation limit (a, = 1). The short linear extrapolations of Schenck, Frohberg, and Jaspert's a, data yield values of C, in good agreement with those of Lander, Kern, and Beach." The solubility data for nickel discussed here is presented in Fig. 1 in the form of plots of ln[Cc/(l - 2cc)] vs 1/T. The discrepancies mentioned are clearly seen. In this investigation a vapor transport method is used to measure C, over as wide a temperature range

Jan 1, 1969

By Peter N. Blakey, Thiann R. Yu, Douglas O. Tansey

A flexible, efficient, and relatively low-cost drilling and blasting operation has evolved at Texasgulf Canada, Ltd.'s Kidd Creek underground mine. Located 32 km (20 miles) north of historic Timmins, Ontario, the underground operation utilizes a sublevel open stoping method developed with large-diameter blastholes. Over the years, Kidd Creek engineers have experimented with various drill-and-blast techniques, leading to the present-day practice. Continuing work is being conducted to further mechanize the blasting operation.

Jan 6, 1976

By J. E. Burke

Recent investigations have elucidated many of the phenomena of the grain growth process, but have also revealed some conflicting and unexplained results. Beck and his co-workers 1,2,3 have shown that grain growth continues with continued heating at one temperature until the average grain size approaches the thickness of the specimen. Plotting the logarithm of the average grain diameter, D, against the logarithm of the annealing time, t, they find the straight line relationship given by the equation: D = Ktm [,] with possible small deviations at short annealing times. K and n are constants at constant temperature. On the other hand, Wa1ker4 finds that, in alpha brass) the rate of growth falls off more rapidly than indicated by Eq i after long annealing, so that an equilibrium grain size is approached. Burke5 showed, using Walker's data, that time and temperature for grain growth are apparently related through an activation energy, H, according to the relationship: ] where h and tz are the annealing times necessary for the average grain size to increase from Do to D at temperatures T1 and T2 and R is the gas laws constant. In the case quoted, H had the value 60,000 cal per mol. The apparent validity of this relationship, at least up to 700°C, and the approximate value of H were later confirmed by Beck3 The results obtained by Beck Kremer, Demer and Holzworth2 for aluminum and aluminum-magnesium alloys could not be satisfied by such a relationship, since the slope of the isothermal curves increased with increasing temperature. These data indicate and Beck, Holzworth and HuV ater demonstrated expcrimentally that the only apparent factor controlling the rate of grain growth at constant temperature is the grain size. While many of the phenomena have been elucidated, a reasonably quantitative explanation of the process is lacking. There exists uncertainty even as to the primary driving force for grain growth, although there is a considerable body of indirect evidence that the interfacial energy of the grain boundary is responsible.7,8,io Making implicit use of this assumption, several authors14,16 have pointed out that grain boundaries should migrate toward their center of curvature. Harker and Parkerg have advanced this concept furthest, and show that one can expect grain boundary migration to occur only when the grain faces are curved, or when they meet at non-equilibrium angles. When the face is curved toward a grain A, for example, the atoms in the surface of grain B, on the other side of the boundary7 are on the average more surromded by the atoms of their lattice than the atoms on the surface of grain A. Because of thermal motion, the atoms at the interface will sometimes be attached to grain A and

Jan 1, 1949

By Young C. Kim, Francis Martino, Ish Chopra

For the past two and one-half years, various emission control strategies were jointly developed by the University of Arizona and the Homer City Owners to com- ply with the clean air requirements at its Homer City Power Plant. Some of the available options for the SO2 emission control by the Homer City Owners were:l)cleaning of run-of-mine coal,2) scheduling of mining operations at its captive mines, 3) blending the captive mine coal with externally purchased low sulfur coa1,and 4) placing a partial FGD system. This paper describes two specific emission control strategies developed and implemented. These are: 1) production scheduling at the mining face to minimize sulfur variability, and 2) external purchasing of low sulfur coal for further blending. The main objective of this paper is to share our experiences on implementing the developed strategies. This paper emphasizesthe need for proper mine planning on a short-term basis, if the power plant is to meet the desired objective.

Jan 1, 1983

By Benjamin Miller

FROM earliest recorded times, limestone has been employed in the industrial life of peoples of all sections of the world where it exists. It is widely distributed and therefore has been available in almost every country and has been employed for a variety of purposes. Its earliest use was for buildings, since it is more easily quarried and dressed than most other kinds of rocks and so could be utilized by early inhabitants, poorly equipped, with tools. The next use was in the manufacture of lime, which was employed by the ancients in mortar. These two uses of limestone are still all important, but they have been overshadowed in scores of regions by literally hundreds of new applications for raw lime-stone or products made from it. Anyone familiar with modern manu-facturing processes can quickly call to mind many industries in which limestones are essential, and new uses are continually coming to notice. Any enumeration would probably be incomplete and soon out of date.

Jan 1, 1935

By M. E. Nicholson

IN order to understand how alloying elements influence hardenability through their effect on the rate of pearlite nucleation, it is advantageous to use a model to describe the mechanism of pearlite nucleation. The model which currently is most widely accepted has been described, among others, by Mehl et a1.1-3 and by Hultgren.4 Pearlite nucleation involves both the nucleation of ferrite and cementite.4,5 Cementite nucleates first, followed rapidly by the formation of ferrite at the cementite-austenite interface. The conclusion that cementite is the first phase to form during the process of pearlite nucleation is based, in part, on the observation that proeutectoid cementite is continuous with pearlitic cementite and that the orientation relationship between proeutectoid ferrite and austenite is not the same as that between pearlitic: ferrite and austenite.6,7 Recently, Smith" has suggested that the dissimilarity in orientation relations, instead of indicating the pearlite nucleation sequence, may be evidence that pearlite is nucleated in a grain adjacent to the one in which it is growing. Also, Modin9 as shown that ferrite in pearlite may be continuous with proeutectoid ferrite, and he has also shown that proeutectoid cementite is not always continuous with pearlitic cementite. Thus, it appears that the conclusion that cementite always initiates the pearlite nucleation process is open to serious question. The author believes that important evidence on the question of how pearlite is nucleated exists in the rates of pearlite nucleation of different steels. The following is a review and analysis of this evidence. If the nucleation of pearlite is considered as requiring the nucleation of both cementite and ferrite, then the time to nucleate pearlite consists of the time to nucleate the first phase to form, plus the time to nucleate the second phase at the advancing interface of the first. In the absence of a visible proeutectoid constituent, the second phase must be nucleated very soon after the first. If cementite is the first phase to form, the time to nucleate pearlite will be approximately equal to the time to nucleate cementite. Since the nucleation rate of a phase precipitating from solid solution increases with the degree of supersaturation, it would be expected that the rate of pearlite nucleation should increase with increasing carbon content. Digges10 has shown that the reverse is true. He determined the quenching velocity necessary to suppress pearlite formation in plain carbon steels as a function of carbon content. His results showed that pearlite formation could be suppressed with lower quenching velocities as the carbon content of the steel increased, at least up to 1.25 pct C, the highest carbon content he studied. Thus, it would appear that in plain carbon steels, in which pearlite is formed at the nose of the C-curve, the hypothesis that cementite nucleates first does not predict the observed change in nucleation rate. The validity of the current model of pearlite nucleation may be tested also by determining whether or not it predicts the relative influence of alloying elements on the nucleation of ferrite and pearlite. This can be done best by a consideration of the changes in a TTT diagram produced by a change in alloy content. Hultgren4 suggests that the general shape of a TTT diagram, as well as what constituents will form from austenite, can be predicted from the relative nucleation rates of ferrite and cementite. This is illustrated in the schematic diagram for a hypoeutectoid low alloy or plain carbon steel shown in Fig. 1. The C-curves for ferrite nucleation and for cementite nucleation are represented by AA' and CC', respectively, and the start of ferrite formation and pearlite formation are represented by AX and BC respectively. Adopting the current theory, Hultgren suggests that between X and C pearlite forms directly from austenite, since cementite nucleates more rapidly than ferrite in this temperature range. Conversely, at temperatures above Tx, ferrite is the first to nucleate and proeutectoid ferrite is formed. As proeutectoid ferrite grows, austenite becomes enriched in carbon, so that the cementite nucleation rate is increased. As a result, pearlite nucleation above temperature Tx is accelerated, beginning along BX and not along C'X. Temperature T, represents the temperature at which the nucleation rate of ferrite equals that of cementite. By extending this type of reasoning, using the principle that the nucleation rates of ferrite and cementite increase with increasing supersaturation of austenite with respect to these phases, it should be possible to predict the influence of carbon content on the shape of the isothermal transformation diagram. For a slightly higher carbon content than that of Fig. 1, the curve A A' representing the beginning of ferrite formation should be moved to the right, whereas the curve CC' representing the beginning of cementite formation should be moved to the left. For hypoeutectoid plain carbon and low alloy steels, this prediction does not agree with existing data. Instead, as the carbon content is increased, the pearlite curve moves to the right along with the ferrite curve and always appears to join it tangen-tially. This behavior is demonstrated by the TTT-curves of a series of manganese steels" shown in Fig. 2. In this figure, TTT-curves for four alloys containing 0.20, 0.40, 0.60, and 1.20 pct C are superimposed. The curves are identified in the figure legend. The letters F, P, and C are used to indicate the start of ferrite, pearlite, and cementite formation, respectively. In these steels, it appears that the nucleation of ferrite and of pearlite are related processes. As a result of the foregoing evidence, it is proposed that the formation of pearlite be considered as follows: Either ferrite or cementite may nucleate from austenite, then grow until the remaining phase is nucleated at the interface between the growing phase and austenite. (The growth of the first phase to form is assumed to be accompanied by a composition change in the adjoining austenite.) According to this model there are two modes of nucleating pearlite: 1—where cementite initiates the succession

Jan 1, 1955

By W. L. Fink, L. A. Willey

DURING the war there was introduced a new higher strength aircraft alloy designated 75S.1,2.3 The properties of this alloy assure extensive applications in both military and commercial aircraft. It is therefore important to understand its limitations as well as advantages. The properties of 75S alloy in the heat-treated tempers, like those of many other aluminum alloys, are affected by the rate of quenching. In general, low quenching rates tend to cause low strengths and to adversely affect the resistance to corrosion. Commercial quenching procedures are usually sufficiently rapid to permit the development of maximum properties. However, in the center of heavy sections, such as large rolled and extruded shapes, the quenching rates may not be rapid enough to develop maximum strength or resistance to corrosion. The investigations of 75S alloy described herein reveal the most critical range of temperature and the effects of quenching sheet and extrusions at different rates through that temperature range, on the tensile properties and resistance to corrosion. PROCEDURE Preparation of Specimens Sheet specimens (3 in. by 9 in.) were prepared from a commercial lot of 75S "as rolled" sheet (0.064 in. thick). The composition° was as follows: Weight Percentage Cu Mn Mg Zn Cr Al with Normal Impurities 1.64 0.16 2.58 5.66 0.27 Remainder The specimens employed for the determination of the time-temperature relations during quenching of sheet were prepared by the following procedure: the hot junction of a 28-ga. iron-constantan thermoelement (asbestos insulated) was brazed to a thin strip of aluminum which was in turn brazed into the end of a 3/16 in. i.d. by ¼ -in. o.d. by 30 in. long aluminum protection tube. Before brazing, the tube was flattened to 0.064 in. thickness over the junction and adjacent lead wires for a distance of 1 ½ in. The flattened end of the tube was fitted into a slot cut from one edge to the center of the 3 in. by 9 in. by 0.064 in. sheet and brazed into place. The brazed joint was dressed to the thickness of the sheet and the entire specimen buffed to approximately the brightness of "as rolled" material. Because of the brazing operations required, these specimens were prepared using 3S alloy. For the investigation of quenching rates in heavy sections, there were used 75S extrusions which had compositions similar to that of the sheet. The specimens consisted of 14 in. lengths of ¾ in., 2 in., 3 in. and 4 in. diam cylindrical rods and a roughly rectangular extruded section which was approximately 4 in. by 7 ½ in. Thermocouples were peened into the bottoms of 1/8 in. diam wells drilled into one end and parallel to the axis to a depth of

Jan 1, 1947

By C. F. Weinaug, P. M. Blair

Many difference equations used to approximate reservoir flow problems treat the phase pressures implicitly but not the mobility-density coefficients. Such difference equations are neither wholly explicit nor implicit, but might be described as mixed. Mixed equations are relatively easy to apply. But the associated time truncation error is relatively large, and when used to solve problems characterized by high flow rates, these equations may be unstable for practical size time steps. This paper outlines the development of a completely implicit difference analogue for reservoir simulation, along with a Newtonian iterative method for solving the resulting nonlinear set of algebraic equations that arise at each time step. While this implicit equation requires two to three times more work than does the mixed equation, it is shown to markedly decrease the time truncation error and to yield a stable solution for much larger time steps than does the mixed equation. INTRODUCTION Calculation of multiphase, multidimensional flow in porous media is generally accomplished by numerical methods that involve approximating systems of partial differential equations by systems of partial difference equations. The early development of difference equation techniques was directed toward the solution of linear differential systems. However, the equations of multiphase fluid flow through porous media are highly nonlinear in that mobility and density often are strong functions of pressure and saturation. Thus, solution of the flow problems by difference equations involves solving sets of algebraic equations whose coefficients change from step to step of the calculations. Current practice is to evaluate most coefficients at the beginning of a time step and then to apply difference methods well suited for solving linear problems. At least two major difficulties arise from the use of this practice. First, evaluation of coefficients at the old time level and evaluation of pressures at the new time level results in larger time truncation errors near displacement fronts than if all quantities in the distance derivative were evaluated at the same time level. Second, evaluation of mobility coefficients at old time levels results in an unstable difference equation in regions of high flow rate. The first of these difficulties often results in optimistic recovery behind the displacement front, and the second requires the use of small time steps or special techniques in solving coning problems1 or problems of gas percolation.2 Evaluation of all quantities in the distance differences at the new time level results in a completely implicit difference system. Such a difference system has a lower time truncation error than equations in which mobilities are evaluated at the old level and pressures evaluated at the new level. The algebraic equations resulting from the fully implicit difference equation are, in this case, nonlinear and require some iterative method for their solution. This paper develops fully implicit difference equations for two-phase flow in porous media, describes their solution by Newtonian iteration, and gives examples of problems more easily solved by the new method than by previous methods. EVOLUTION OF DIFFERENCE EQUATIONS FOR NONLINEAR PROCESSES The completely implicit difference equation used in these calculations results from evolution beginning with explicit difference equations, proceeding through what we will call "mixed" equations, and arriving now at a completely implicit equation. Let us illustrate the evolution by considering different analogs of the nonlinear equation

Jan 1, 1970

By Henry D. Hibbard

FoR the purpose of this paper all steels will be divided into two divisions: effervescing and non-effervescing. This classification must be borne in mind as many statements true of one class are not true of the other. Effervescing steels are those that effervesce in the molds, evolving large volumes of gases, which, in escaping, throw up a spray of molten particles of steel and give the molten metal in the mold a rapid stirring or churning motion until it freezes, or, at least, becomes pasty. This effervescing is intended and proper. The non-effervescing steels, which are more or less completely killed, should not evolve any gas in the molds. Many years ago the writer proposed to call the two classes of steels evolution and solution steels, depending on whether the gases in the molds are evolved or kept in solution, but the present designations seem to be better. In America, at least, effervescing steels are in a class by themselves and include steels made for boiler, tank, and ship plates, steel pipe, soft steel wire, wire nails, soft machincry, and structural steels. These are all low-carbon steels with less than 0.40 per cent. carbon. Low-carbon steels that do not effervesce but are killed are also made; these are put into ingots and castings but they are outside the limits of this paper. Effervescing steel is cast into square or slab molds, either top or bottom poured. Most of the effervescing steel is made in the basic open-hearth furnace, though much is made by the acid Bessemer process and a little is made in the acid open-hearth. It is not made by the crucible or electric-furnace process, in which the volume of gas required cannot be produced in the steel and evolved from it in the mold to give the effervescence desired. Any steel ingot, when teeming is finished, is of a certain size or volume. Leaving out of consideration for the moment the lessening of the over-all dimensions of the ingot through contraction due to loss of heat, the ingot may be placed in one of three classes: It may settle, losing in volume; it may stand, that is, remain constant in volume; it may rise, that is, increase in volume. Into which class an ingot will fall depends on the gases of the steel,

Jan 1, 1920