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By A. Steiner, E. Miller, K. L. Komarek

Equations have been derived to calculate the chemical potentials of the components of liquid binary alloys from liquidus and enthalpy data. The equations are applicable to systems with intermetal-lic compounds of limited solid solubility. The liquidus curve of the Mg-Sn system was accurately redetermined, and the melting point of the compound Mg2Sn was found to be 770.5° * 0.3°C. The thermo-dynamic properties of the liquid alloys were calcu -lated from the liquidus data. The activity of tin shows both positive and negative deviations from Raoult&apos;s law and the relative integral entropy is positive. THE phase boundaries in an equilibrium phase diagram are functions of the thermodynamic properties of the coexisting phases. Specifically, the shape of the liquidus curve near an intermediate compound is determined by the properties of the compound and the liquid phase. Hauffe and wagnerl have derived an equation for the chemical potential differences (ui) of the liquid alloys near the congruent melting point when the heat of fusion of the compound and the liquidus curve are known. In their derivation the temperature dependence of the ui value was neglected, and liquid of stoichio-metric composition was chosen as the reference state. To obtain chemical potentials over the ntire composition range of the phase diagram with the pure components as the reference state, equations have been derived which incorporate the temperature dependence of the chemical potentials. Such a thermodynamic analysis can be carried out when the liquidus curve and the enthalpy values are known. The equations have been applied to the redetermined liquidus curve of the Mg-Sn system to obtain the thermodynamic properties of liquid Mg-Sn alloys. The Mg-Sn phase diagram has been studied by various investigators and the results have been compiled and critically assessed by Hansen.2 Preliminary thermal analyses showed that the basic features of the diagram were correct. The system has one congruent melting compound, Mg2Sn, two eutectics, and some solid solubility of tin in magnesium. However, the published liquidus temperatures show considerable scatter. In order to perform an accurate thermodynamic analysis, the liquidus curve between the magnesium-rich and tin-rich eutectic compositions was therefore redetermined by careful thermal analysis using highest-purity starting materials. EXPERIMENTAL PROCEDURE The magnesium metal (Dominion Magnesium Ltd, Toronto, Canada) had a purity of 99.99+ pct with the following impurities (in ppm): 20 Al, 30 Zn, 10 Si, <1 Ni, <1 Cu, <10 Fe; tin (Cominco Products, Spokane, Wash.) contained 99.999+ pct Sn. The graphite crucibles were machined from graphite rods (United Carbon Products Corp., Bay City, Mich.) which had an average total ash content of less than 30 ppm. All the graphite parts were baked out in vacuum at 950" to 1000°C for a minimum of 24 hr to remove volatile organic materials. The graphite crucibles (3 in. long, 1-5/8 in. OD, 1-3/8 in. ID) were tightly closed with a threaded cover (1/2 in. thick) which had a central thermocouple well (2-1/4 in. long, 5/16 in. OD, 3/16 in. ID). Cover and thermocouple well were machined from one piece of graphite. A thin sheath of quartz was inserted in the well to protect the thermocouple

Jan 1, 1964

By I. S. Servi, N. W. Marr

TIN markedly increases the softening temperature of pure copper with only a moderate effect on conductivity. Smart and smith&apos; indicated a substantial increase in softening temperature upon addition of different levels of tin up to 0.05 pct in oxygen-free copper. However, after finding no effect of tin in copper containing about 0.025 pct O, they concluded that "the presence of oxygen completely nullifies these changes in properties by precipitation of SnO2.&apos;&apos; Since preliminary tests on Cu-Sn-0 alloys indicated that this conclusion is not valid outside the extremely limited tin range investigated by Smart and Smith, research was primarily directed toward the partition of tin between the copper-base matrix and the oxide phases. Experimental. Material was prepared by sealing an 80 to 100-g charge consisting of oxygen-free copper, tin, and copper oxide into a y-in.-ID quartz capsule which was evacuated and back-filled with tank-grade argon to 3 atm. The sealed vials were heated to 1125ºC, kept at temperature for 30 min, occasionally agitated, and then allowed to air-cool outside the furnace ("as-cast"). The castings were cold-swaged to 0.370 in., annealed 1 hr at 600°C ("as-annealed"), and cold-swaged 16.5 pct ("as-swaged"). Cross sectional wafers were annealed for 1 hr at various temperatures and Rockwell F hardness measured. The "softening temperature" was defined as the interpolated temperature at which the hardness was 50 Rockwell F. Electrical resistance was measured by an ac poten-tiometric method on 0.040-in.-diam wires after a 1 hr anneal at 600°C. Relative accuracy was of the order of 1 pct. Chemical analyses for tin were performed by two methods, identified as "Acid flux" and "HN03 dissolution". The first consisted of fusing the sample with potassium pyrosulfate and leaching the resulting melt with 10 pct sulfuric acid. The spectrophotometric procedure which followed is described in Ref. 2. Oxygen analyses were performed by the vacuum-fusion method with a bath temperature of 1650°C. Tin content of oxides and copper phases was measured with the X-ray microprobe, using a calibration curve based on a series of oxygen-free Cu-Sn alloys. Results. In Fig. 1, the nominal tin composition for alloys containing nominally 400 ppm O was plotted vs the analyzed or the computed tin content—the latter derived from one of the following: 1) "acid flux" chemical analysis; 2) "HN03 dissolution" chemical analysis; 3) X-ray microprobe analysis of the copper matrix; 4) equivalent tin content, based on the measured increase in electrical resistivity (an increase in resistivity of 1.6 microhm-cm is approximately equivalent to 1 wt pct of sn1,3). It is evident from Fig. 1 that the results of the acid flux analysis are consistent with the nominal composition while the HNO3 dissolution method generally yields lower values, consistent with those obtained from electrical resistivity measurements and somewhat higher than those obtained by microprobe analysis. Microprobe and resistivity values are affected only by the tin in solid solution; hence the HN03 method

Jan 1, 1970

By K. R. Evans

The yield and fracture characteristics of a Ti-8 pct A1 alloy have been examined at room temperature as a function of exposure temperatures to 1700°F. Em-brittlenzent of the alloy is observed to occur following exposures below 1280°F, the critical temperature for formation of ordered Ti3Al particles. Presence of Ti3Al particles results in an increase in the yield and flout stresses which, for a given exposure time, are a maximum at the critical temperature. Embrittlement is attributed to the occurrence of plastically induced fracture controlled by the increase inflow stress. This interpretation is compatible with the decreasing ductility and increasing crack popagation rates observed with increasing exposure time below the critical temperature. Slip line observations and measurements of the gain size dependence of the yield stress indicate that the slip character of the alloy is not influenced by exposure treatments below the critical temperature. ThE potential usefulness of Ti-A1 alloys in structural applications has prompted considerable interest in studies of their microstructure and mechanical characteristics. The relatively recent surge of interest in the system and at the same time its complexities are indicated by the fact that despite numerous recent efforts14 there has not even been general agreement upon a phase diagram for the system. However, most of the Ti-A1 alloys of commercial interest range in aluminum content from 5.0 to 8.5 wt pct where most investigators achowledge the existence of the ordered Ti41 phase. Goldak and parrs confirmed the existence of Ti41 from X-ray diffraction analysis of the stoichiometric composition while Blackburn6 has determined its presence in commercial Ti-8 pct A1 alloys by electron diffraction analysis and dark field transmission microscopy. The embrittlement of Ti-8 pct A1 alloys was initially reported by Crossley and carew7 who showed that aging in the temperature range of 750" to 1110°F resulted in extreme brittleness. This behavior was later confirmed and correlated by Soltiss to the presence of long-range order in the commercial Ti-8 pct Al-1 pct Mo-1 pct V alloy (Ti-8-1-1) which exhibited anomalous creep and yield strength behavior at 950°F. This investigation was designed to study the mechanical characteristics of a binary Ti-8 pct A1 alloy in order to further examine the nature of the embrittlement behavior reported for the alloy system. EXPERIMENTAL PROCEDURE A Ti-8 pct A1 binary alloy was obtained from the Wah Chang Co. in the form of 0.050 in. sheet. The chemical composition of the alloy is given in Table I. Heat-treatment of the alloy was conducted in evacuated vycor capsules for the tensile sDecimens (1.25 bv 0.25 in. gage section) and in evacuated stainless steel jackets for the crack growth-rate panels. All specimens were annealed at 1700°F for 24 hr and water quenched before an exposure at elevated temperatures. Tensile testing was conducted at room temperature in an Instron testing machine equipped with a pushbutton control for instantaneous control of the cross-head motion. In this manner the change in stress, Aa, accompanying an increase in strain-rate by a factor of 10 was measured immediately after yielding. The crack growth-rate panels used were 3 by 18 in. with a precrack 0.33 in. long cut into the center of the panel. The panels were cycled at a maximum gross stress of 40,000 psi at a maximum to minimum tension stress ratio of 20. Cycling of the specimens was continued until the crack propagated the width of the panels. The crack length was intermittently measured during slow crack growth. EXPERIMENTAL RESULTS Correlation of Embrittlement to the Critical Temperature for Ordering. Blackburn6 has found that the critical temperature, T,, for the formation of Ti41 in a Ti-8 pct A1 alloy with an interstitial content on the order of that given in Table I is approximately 1245°F. The relationship between the ductility of the binary alloy used in this investigation and T, was estimated from tensile tests on specimens which were exposed for 24 hr at various temperatures followed

Jan 1, 1970

By Paul Torrey

IN his recent article, Stanley C. Herold described some of the principles involved in oil-field rejuvenation. The purpose of this paper is to present certain facts, based upon observations of systematic water flooding of oil sands in the Bradford and Allegany fields of northwestern Pennsylvania and southwestern New York, which, it is believed, will partially substantiate Mr. Herold&apos;s opinions. The writer gratefully acknowledges his obligation to the Forest Oil Corpn., the Kendall Refining Co., the Petroleum Reclamation Co., W. J. Healey, of Bradford, Pa., and the Producers Oil Corpn., of Olean, N. Y., for information included in this article; and he is especially indebted to Jerry B. Newby, of the Petroleum Reclamation Co., for helpful suggestions and criticism in the preparation of the paper. It is obvious that, in the calculation of mathematical formulas for the prediction of oil recovery, uniform conditions must be assumed; otherwise, the applicability of the results of such work will be, by necessity, restricted to the area or individual property for which they were designed. Geologic and physical conditions are so variable throughout the fields where water flooding is practiced that it is difficult to devise methods that can be universally applicable. However, it is believed that there are certain fundamental principles which, when applied with specific modifications designed to meet the conditions existing in local areas or individual properties, may yield considerably greater recoveries of oil.

Jan 1, 1927

By F. C. Blickensderfer

During 1914 and 1915, extensive experiments were conducted at the concentrator of The Detroit Copper Mining Co. of Arizona, at Morenci, Ariz., in order to test the relative grinding efficiencies of the Marathon, the Chilean and the Hardinge Mills. The Marathon mill used in the experiments was the first one ever built by the manufacturers; and since this was the first time that it was given a thorough tryout, with such remarkable results, the data have been assembled in the following paper and submitted in the interests of the milling profession. Description of the Flow Sheet The concentrator flow sheet, as it existed at the time of these tests, is shown in Fig. 1. The total mill feed, as delivered from the crushing plant, had as its maximum size 1-in. material. This material was screened dry on Cole Zig-Zag screens fitted with 4-mesh openings; the oversize going to one pair of coarse rolls, the undersize to five Wilfley tables fitted with Butchart National riffles. The coarse-roll product was elevated and screened wet on Zig-Zag screens fitted with 4-mesh openings, the undersize going to the five tables mentioned above and the oversize to one pair of fine rolls. The fine-roll product was elevated and screened together with the coarse-roll product; so that, before any concentrating operations took place, the total mill feed was reduced to pass a 4-mesh screen. After concentration on Butchart-riffled tables, the tailings from these tables passed to de-sliming hoppers, where two products were formed, viz., an overflow and a discharge. The discharge constituted the feed for the Chilean mills. It contained 16 per cent. slime (—200 mesh) and 3 parts of water to 1 of solids. The concentrator was equipped regularly with four Chilean mills (Monadnock type) to the unit; but, in order to

Jan 1, 1917

By E. R. Stover, J. Wulff

A tentative 870°C isothermal section, the solidus equilibria, and the solubility of Tic and graphite in the nickel solid solution have been determined with arc-cast specimens. Each of the Ni-Ti intermetallics forms two-phase fields with Tic; only Ti,Ni has a measurable solubility for carbon. A quasi-binary eutectic exists between Tic: and a titanium-rich nickel solid solution; ternary eutectics occur on either side, with TiNis or with graphite. Graphite is in equilibrium with a nickel solid solution containing a ratio of Ti/C which increases below the solidus. DURING the past decade, titanium-carbide cermets containing nickel or cobalt alloy binders have been widely studied for high-temperature and wear-resistant applications. Recent work has emphasized control of the carbide grain size and distribution2&apos;s and the thermal exparision match between the phases/ The present study was initiated to determine the phase compositions during the sintering or infiltration process, and during subsequent heating. Although the effects of additives and impurities in commercial compositions are important, it was necessary to consider first the pure ternary, Ni-Ti-C. Emphasis was placed on the nickel corner, since the nickel solid solution is the continuous phase in compositions having highest strength. Early metallographic work by Zarubin and Molkov- showed that solid solubility at either end of the section Ni-Tic is small. Edwards and Raine reported that nickel dissolves between 3 and 5 wt pct Tic at 1250° C, on the basis of microstructures of annealed, vacuum-melted powder mixtures. Although the solubility of nickel in Tic has not been measured, Tic formed by the menstruum process in liuid-iron solution contains as little as 0.06 wt pct Fe. BINARY SYSTEMS The most recent determinations of the three binary systems are summarized in Fig. 1. Titanium carbide (6) is predominant among the binary phases, and it melts above the boiling point of nickel. The Ni-C binary is based on separate liquid solubility&apos; ) and solid solubilitylo measurements. The eutectic temperature, 1328OC, is taken from the solidus determination described below. However thermal analysis data by Morrogh and williams also show eutectic arrests in this vicinity in the absence of supercooling. The Ni-Ti binary is taken from the work of Poole

Jan 1, 1960

By F. H. Wilson, R. M. Brick

WITH the papers of Palmer and Smith1 and of Burghoff and Bohlen,2 published in 1942, understanding of the problem of the development of ears on deep-drawn brass cups was brought to the point where, from a practical standpoint, producers of brass sheet could keep earing under control. However, these very practical results raised questions concerning the reasons for the observed behavior. Since plastic deformation is fundamentally a function of crystallographic slip, these reasons must be considered in the light of the textures to be found in the annealed sheet. Various attempts have been made to interpret the mechanism of ear formation by correlation with the preferred orientation and the anisotropy of the sheet. However, none of these has satisfactorily explained why ears are usually found in the rolling and transverse directions for copper but at 45° to these directions for brass. This paper offers additional data on these two related subjects and has been divided into two sections. The first section provides a survey of the recrystallization textures of brass as affected by annealing temperatures and prior reductions, and an account of the earing behavior which accompanies these textures. The second section reports new data on the anisotropy that accompanies a characteristic texture for brass and endeavors with partial success to establish the factors that account for the difference in earing behavior between copper and brass. EARING BEHAVIOR OF BRASS Review of Textures The recrystallization texture of severely rolled brass was first reported by Bass and Glocker.3 The preferred orientation to which this texture, obtained from metal annealed at 500°C., was most akin was (113)[112]. This texture has been confirmed by Bauer, von Göler and Sachs;4 Brick;* Brick, Martin and Angier;6 and Cook and Richards7 who, in general, used annealing temperatures under 500°C. Bass and Glocker3 also reported the recrystallization texture at 850°C. to be the same as at 500°C. They analyzed their textures by means of single photograms obtained with white radiation. The photograms for both temperatures are reproduced in their paper, and there is a strong similarity between the two. However, for the latter temperature, the number of spots registered on the film was remarkably large for the grain size to be expected, and the authors made no mention of using a technique that would integrate a large enough area of the specimen to produce so many spots. The only other attempt to determine textures resulting from high temperature

Jan 1, 1945

By C. G. Dunn, J. L. Walter

Either (100) (001] or (110) (0011 oriented secondaries, or both, depending on annealing atmosphere and material composition, grew in a matrix of 2-dimensional grains. The growth dependence of secondaries of(hkl) orientation in the iron and 0.6 pct Si-Fe is explained in terms of surface energy driving forces and the dependence of these driving forces on both (hkl) orientation and amount of oxygen at the surace. (100) [001] grains grow when suflicient oxygen is present at the gas-metal interface to make y100 the lowest (hkl) surface energy. When little or no oxygen is present at the gas -metal interface, making Y110 the lowest surface energy, (110) (001 1 grains grow. When Y100 and y110 are approximately equal but lowest among (hkl) surface energies, both (100) (0011 and (110) (0011 grains grow simultaneously. An excess of surface oxygen produces complete grain growth inhibition in high-purity iron; i.e., a 3-dimensional grain structure is retained over long annealing periods. However, a small amount of silicon in the iron reduces the growth inhibiting tendency of oxygen without altering the (hkl) surface energy relationships. Some information on annealing twins is presented. SECONDARY recrystallization in high-purity a iron, cold rolled to strip a few mils thick, was reported recently.&apos; The present paper provides additional information on this subject, also furnishes similar information on both zone-refined iron and a high-purity 0.6 pct Si-Fe alloy, and treats the more important surface phenomena that determine to a large extent whether one or more secondary recrystallization textures will appear during annealing. Previous studies on high-purity 3 pct Si-Fe2,3 have shown the importance of (hkl) and impurity-atom dependent gas-metal surface energies on the formation of secondary recrystallization textures in high-purity material. For example, either (100) or (110) oriented grains may have the lowest surface energy depending on the amount of oxygen at the gas-metal interface, and hence the highest driving force for growth if other growth factors are the same. Another important phenomenon in high-purity material is stabilization of the matrix structure; neither impurities nor a strong matrix texture seem to be necessary when recrystallization and normal grain growth occur, developing a stable two-dimensional type grain structure. The average grain size in such a structure is usually between one and two times the specimen thickness. Consequently, most of the grains have grain boundaries terminating at exposed surfaces, and this results in gas-metal interfaces (or hkl surfaces) as well as

Jan 1, 1962

By William H. Kobbé

The word recovery as used in this paper is applied in its broader sense and not limited to wells producing from horizons of unconsolidated sands. Certain problems connected with the winning of petroleum from such horizons present themselves while drilling is in progress and should be considered before drilling is commenced wherever unconsolidated sands are expected. Maximum recovery depends not alone. upon efficient pumping methods but also upon the selection of a proper drilling system and the completion of a well especially adapted to the extraction of these oil-bearing sands. 11. DRILLING Determining Factors in Selection of Method The selection of a drilling method for the development of a region in which the oil reservoir consists of unconsolidated sands is based on the following considerations: (1) Character of the formations overlying the reservoir; (2) thickness of the oil-bearing stratum or strata; (3) thickness and character of water or gas sands; (4) gas pressure to be encountered or expected; (5) total depth of well. 1. Character of the Formations Overlying the Reservoir The strata overlying the reservoir may consist of beds of shales, clays, sands and other soft materials, or they may be slates, limestones and sandstones consolidated to all degrees of hardness. Very often the overlying strata are predominantly of soft materials but interbedded with harder formations. Likewise, soft strata are encountered at different horizons in an overburden consisting almost entirely of the harder sedimentaries. The physical character of this overburden is frequently

Jan 1, 1917

By A. Gemperle

Electron diffraction and transmission electron microscopy on foils at room temperature were used to investigate the ordering of Fe-Si alloys containing 10 to 23 at. pct Si. A certain degree of DO3 order was found in all alloys. With the exception of the lowest silicon concentration for which the antiphase domains could not be clearly resolved, the alloys have a domain structure of two-domain type with boundaries having 1/4a01<111> displacement vectors for less than 12.3 at. pct Si and with boundaries having 1/2 a0<100> displacement vectors for more than 12.3 at. pct Si. The alloys with 12.3 at. pct Si have a domain structure consisting of fine domains with 1/2a&apos;o<100> boundaries within much larger domains with 1/4a&apos;o<111> boundaries. The development of these structures can be explained by transition of the alloy from the disordered state into the B2-type order and then into the D03-type order by the mechanism proposed previously for the FeSAl alloys. The existence of the B2 structure in the lower part of the investigated concentration range reported in some articles can be explained by fine domains with 1/2a&apos;o<100> boundaries formed by several disordered planes within large domains with 1/4a&apos;o<111> boundaries. The ordered structure predicted by the theory —with practically no domain boundaries —is found in the alloys having 12.3 at. pct Si where it develops in the B2 structure region. ORDERING in Fe-Si alloys was first studied by phragmenl who found that beginning with 13 at. pct Si the DO3 (Fe3Si) superlattice reflections appear in the diffraction patterns. The equilibrium diagrams constructed later by corson2 and Haughton3 from various measurements proposed the existence of a homogeneous solid solution (a phase) in the range from 0 to 25 at. pct Si. Jette and Greiner4 and Farquhar et al.5 measured the relation between lattice parameter and composition and they considered the break in the curve at 9 to 10 at. pct Si to be caused by the ordered solution a"(Fe3Si). Glaser and Ivanick6 Determined critical ordering temperatures of the alloys containing from 10.9 to 27.9 at. pct Si from the measurement of the electric resistivity of quenched samples. In all cases the critical temperature was lower than the melting point and it was highest for 25 at. pct Si. Lihl and Ebel7 measured the lattice parameter curves at various temperatures up to 1000°C. The region between two breaks on these curves, corresponding to 10 to 12.5 at. pct Si at room temperature, was considered by them to be two-phase (a + a"). They concluded by extrapolation of the measured values that a" in the alloy having 25 at. pct Si is stable up to the melting point. Davies8 studied superlattice reflections in the X-ray diffraction patterns of an alloy containing 8.7 at. pct Si. He found the B2 structure and short-range order in the slowly cooled samples and the DO3 A. GEMPERLE is Research Scientist, Institute of PhysicS, Czechoslovak Academy of Sciences, Prague, Czechoslovakia. __Manuscript submitted January 2, 1968. IMD structure in the quenched and annealed samples. This investigation first reports the presence of the B2 structure, phase a&apos;: in Fe-Si alloys. Meinhardt and krisement9,10 also found its existence in Fe-Si alloys by neutron diffraction. No order was detected by them in the alloy containing 8 at. pct Si. The onset of B2 order was observed at a composition of 9.2 at. pct Si. They found almost perfect B2-type order with partial DO3-type order at room temperature in the 10 to 12.5 at. pct Si range and almost perfect DO3-type order in the 12.5 to 25 at. pct Si range. They established the critical temperatures Tc of both the structures through measurement at higher temperatures. The critical temperature for the B2 structure was found to be always higher than the critical temperature for the DO3 structure of the same alloy. They extrapolated the curves of the critical temperatures and concluded that the alloys with more than 17 at. pct Si have the B2 structure up to the melting point and the alloys with more than 23 at. pct Si have the DO3 structure up to the melting point. The results of Meinhardt and krisement9,10 were confirmed by Dokken&apos;s measurement" of the temperature dependence of the electrical resistivity in the 10.8 to 15 at. pct Si range. On the other hand chessin12 detected by X-ray diffraction a considerably lower degree of order in the 12.7 at. pct Si alloy. ANTIPHASE DOMAIN STRUCTURE IN ALLOYS WITH B2- AND DO$-TYPE ORDER The ordered structures B2 and DO3 can be described in terms of a subdivision of the bcc lattice into four fcc sublattices with a parameter double that of the bcc lattice. Following Marcinkowski13 we will label them I. 11, 111, IV. The B2 structure in the AB alloy is formed by placing A atoms on sublattices I and II and B atoms on sublattices III and IV. In a non-stoichiometric perfectly ordered alloy having concentration (A) > (B), A atoms occupy sublattices I and 11. and A and B atoms distributed at random occupy sublattices III and IV. The DO3 structure in an A3B alloy is formed by placing A atoms on sublattices I, 11, and 111, and B atoms on sublattice IV. In a nonstoichiometric, perfectly ordered alloy having concentration (A)/3 > (B), A atoms occupy sublattices I, 11, and 111, and A and B atoms distributed at random occupy sublattice IV. As further shown in Ref. 13, two types of domains are possible in the B2 superlattice and the antiphase domain structure has associated with it boundaries with displacement vectors 1/4 a&apos;o<ll1> only. Four types of domains are possible in the DO3 superlattice and the antiphase domain structure has associated with it boundaries with displacement vectors 1/4a&apos;o<111> and 1/2a&apos;o<l00>. Bethe14 suggested on the basis of theoretical considerations that in a structure with two sublattices at low temperatures only one domain should be present in the whole crystal at equilibrium. Similarly Bragg15 concluded that at low temperatures the domain struc-

Jan 1, 1969

By R. F. Campbell, R. F. Miller

SIXTEEN steels containing different combinations of chromium and molybdenum, in amounts up to 5 per cent of each element, were examined for microstructure and hardness after air cooling and after furnace cooling from suitable austenitizing temper- [ ] atures. Diagrams summarizing the structures and hardness resulting from these two cooling rates are presented (Figs. 3, 4, 7).† Data were also obtained on the influence of these two elements, individually and in combination, on the depth 0f decarburization observed after furnace cooling (Fig. 8). COMPOSITION OF MATERIAL The chemical composition of the steels investigated is given in Table I. The plain carbon steel was made in a 25-ton Heroult electric-arc furnace, while the alloy steels were made in a 4.5-ton Ajax-Northrup elec- [ ] tric induction furnace. The cast ingots were hot-rolled to I.75-in. bars. AIR-COOLED SERIES The I.75-in. bars were forged to flat plates, 0.25 in. thick, from which specimens 0.5 in. wide by 2 in. long were cut. The specimens were pretreated to improve their homogeneity by heating to I830°F. (1000°C.) for 20 min. and oil quenching. They were then reheated to I740°F. (950°C.) for one hour and air-cooled, after

Jan 1, 1941

By M. C. Flemings, T. S. Piwonka

Pore formation is examined analytically in cellu-larly solidified single crystals, in cylindrical castings of pure metals and alloys, and in unidirectionally solidified castings. Effects are considered of fluid flow. dissolved gas, and surface tension. Experimental results show that flow through a partially solid "mushy" aluminum alloy obeys relations similar to those applicable for flow through other types of porous beds. Flow rate is proportional to the square of the fraction liquid (fL) for fL less than 0.3. Results indicate models chosen for mathematical description of pore formation in alloy solidification are valid. ONE of the solidification variables which is influential in determining mechanical properties of castings and of wrought material produced from cast ingots is porosity. This porosity may result from shrinkage, dissolved gases, or a combination of both, and may take the form of either macroscopic pores in localized areas of the casting or microporosity evenly distributed throughout large areas of the casting. In alloys which freeze over a range of temperatures some microporosity is generally found after solidification. There has been much discussion as to whether the micropores are caused primarily by shrinkage, by gas, or by a combination of both.&apos;-3 There has, however, been little experimental or analytical work performed to delineate the relative contributions of shrinkage and gas in causing microporosity or to quantitatively describe the influence of processing variables on microporosity. Porosity other than microporosity (e.g., center line shrinkage) has also received only limited quantitative treatment.495 The aim of this work is to examine analytically the general problem of pore formation in solidification and to perform limited confirmatory experiments. Consideration is first given to shrinkage porosity in pure and nearly pure metals (shrinkage in cellular growth, centerline shrinkage in cylinders). Next, a discussion is given of formation of shrinkage-caused microporosity in unidirectional solidification and in "mushy" solidification. Finally, influences of dissolved gas and surface tension on microporosity are examined. Numerical results are given for aluminum and an aluminum alloy. Typical examples of microporosity are shown in Figs. 1 to 3, for 1) cellularly solidified aluminum, 2) sand-cast aluminum alloy, and 3) low alloy steel. SHRINKAGE POROSITY IN PURE AND NEARLY PURE METALS 1) Cellular Growth. In cellular freezing of an impure liquid, liquid grooves may exist over a considerable length of the growing crystal.6 The cross-sectional area of the grooves decreases with increasing distance from the cell tips; the root of the groove is blunt if the liquid at this point reaches eutectic composition. As a result of volume contraction during solidification, flow occurs down the groove, parallel to the growth direction; hence, there must be a pressure drop from the groove entrance to points within the groove. The geometry of the groove is simplified to that of a circular cylinder of radius r and length L, Fig. 4, L » r, and a simple and useful limiting case is obtained. This is that, in the circular cylinder freezing at constant velocity, the pressure at the root of the groove is lower than that pressure at radius r in grooves of other geometries where r decreases continuously to a lower limit at L. As discussed below, a pore resulting from a given pressure difference

Jan 1, 1967

HAVING produced one ton of concentrate that contains the copper that can profitably be saved or recovered from anywhere between 3 and 60 tons of mine ore, the ore-dressing engineer passes the result of his handiwork over to the metallurgist for smelting. Essentially, the concentrate is a damp mass of finely ground material consisting of copper sulphide minerals; pyrite, the sulphide of iron; and gangue, usually silica and silicates of aluminum. The purpose of smelting is to produce metallic copper and a slag containing the iron and siliceous constituents by pyrometallurgy-the kind of metallurgy that employs fire and heat to produce the chemical reactions through which it attains its object. Though blast furnaces have been used to smelt Porphyry-mine concentrates, notably at Braden, more than 90 per cent of [ ]

Jan 1, 1933

By Eduardo Martinez Baca

Although Mexico has always been justly considered by the whole world essentially a mining country, this reputation is due principally to the richness of the deposits which have been worked. In reality, the majority of the States have been little

Jan 1, 1902

By E. N. Aqua, C. M. Owens

The microstructure in cold-rolled Zircaloy-4 sheet has been characterized by transmission electron microscopy and X-ray diffraction line-broadening techniques. The rate of work hardening is measured in terms of microstrains, particle sizes, micro hardness, the tensile yield strength. The tensile yield strength of Prestrained (less than 10 pet) Zircaloy-4 sheet is found to be a linear function of the square root of the dislocation density, us predicted by work-hardening theories. ZIRCALOY is one of the most important structural alloys for nuclear reactor applications.&apos;&apos;2 In order to understand the effects of the nuclear environment on Zircaloy structural members, the preirradiation microstructure must be characterized. This report will describe the substructure existing in cold-rolled Zircaloy-4 sheet as observed by electron microscopy and by X-ray diffractometry. I) EXPERIMENTAL PROCEDURE Commercial Zircaloy-4 strip was cold-rolled and annealed in argon at 760°C for 1 hr. The annealed sheet was again cold-rolled to a maximum of 50 pct reduction in thickness with segments of the strip being retained after 2.5, 5, 10, and 20 pct reductions. One-inch-square portions were chemically polished (50 parts HNO3, 50 parts H2O, 10 parts HF) and mounted onto lucite blocks using double-sided tape to prevent bending during handling. Thin foils suitable for transmission electron microscopy were prepared from each of the sheet samples. The sheets were mechanically polished through 600-grit paper to a 3 to 4 mil thickness and were chemically polished by immersion for 1 to 2 min in a solution containing: 50 parts HNO3, 50 parts H2O, and 10 parts HF, and rinsed in distilled water. The final metallographic procedure was to electropolish for less than 10 sec with an electrolyte of 10 parts perchloric acid and 90 parts glacial acetic acid, at 20 v, at room temperature using stainless-steel point cathodes and a vigorous stirring action. These foils were examined in a Philips EM 200 electron microscope using a 100-kv accelerating voltage.

Jan 1, 1968

By William L. Fink, Dana W. Smith

Many investigators have attempted to determine the true nature of the internal changes taking place during aging. Merica, Waltenberg and Scottlt were the first to propose a theory of age-hardening. They attributed the age-hardening of duralumin to the precipitation of minute particles of the compound CuA12. Later Zay Jeffries and R. S. Archer2 augmented the above theory by suggesting that the small CuAl2 particles acted as "keys" causing "slip interference." Many other investigators3 have devoted much time to the mechanism of age-hardening, some corroborating the "precipitation" and "slip interference" theory, others coming to different conclusions because of effects that to them seemed unexplainable by the early theories; e.g., the increase of electrical resistance upon aging, the fact that no change in lattice parameter could be detected during the early stages of aging even at elevated temperatures, and the fact that the density did not change as would be predicted. In 1932, Merica revised4 his earlier precipitation theory to take care of these presumably anomalous effects. This revised theory, the "knot" theory, assumes that substantial age-hardening occurs before actual precipitation of solute as discrete particles. The assumed mechanism is one of diffusion of solute atoms into groups forming distorted regions which act as "hard" spots and cause slip interference. These "knots" are assumed to have approximately the same lattice arrangement as the solvent. The anomalous behavior of lattice parameter, density, and electrical conductivity, require a modification of the simple precipitation-hardening theory, only if it is assumed that the precipitation of minute particles during aging has the same effects on these properties as the formation of large particles under equilibrium conditions. Data obtained by the authors in the study of the aluminum-magnesium system5 indicated that

Jan 1, 1936

By Maurice Cook

A procedure of obtaining specimens of lead, tin, and cadmium with a moderately equiaxed structure and a smooth surface suitable for etching without grinding and polishing is described; the advantages of general oblique illumination (as opposed to vertical illumination) in the photomicrography of pure metals .are stated. Using these methods, a statistical study of the changes brought about in deformed lead, on annealing, has been made. The common mode of structural change was found to be recrystallization. Cases of growth of one original grain into another were rare; but if a totally recrystallized specimen was annealed above 200° C., a few large grains appeared and spread over the whole area, except for a certain number of small grains which remain unabsorbed. The temperature needed to produce structural change in lead was found to be lower, the greater the degree of deformation. With tin, an extension of 28 per cent. caused practically no recrystallization at ordinary temperatures, but recrystallization occurred if the specimen was heated to 100° C. In cadmium at low temperatures, the grain size decreased in marked manner as the deformation increased; but at high annealing temperatures, the grain size was comparatively independent of the degree of deformation. THE structural changes in metals brought about by annealing following a deformation at a low temperature has been the subject of many investigations. No less than eleven metals and alloys have been studied in some detail; namely, aluminum,1 tungsten,2 iron,3 copper,4 gold,5

Jan 11, 1924

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.&apos; 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&apos;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 J. E. Smith, P. B. Crawford, B. L. Laudrum

A method is presented by which the displacement efficiency of a steam-injection process may be determined. The method consists of a determination of the position of the steam front at a particular time and the application of the Buckley-Leverett theory to the hot and cold regions of the reservoir with the appropriate hot or cold fluid properties. A uniform, linear flow system is assumed for this study. The method is applied to a number of reservoir conditions. The primary properties which affect the results of a steam-injection process are the oil viscosities at steam and reservoir temperatures, injection rate, pay thickness and oil saturation. The efficiency of steam injection is compared with that of water injection for a 10-acre, linearly flooded reservoir. In reservoirs where the oil viscosity is about 35 times the water viscosity, the steam process yields recoveries about 20 per cent higher than the water flood. Where the oil viscosity is near 500 cp, the steam process may recover from three to four times as much oil as a water flood. INTRODUCTION In a previous paper,&apos; a method was suggested for evaluating a steam-injection process. Emphasis was placed upon the determination of the steam-front position in the reservoir at various times. In the present paper, the displacement process is studied in more detail. The approach is essentially that employed by Jordan, Rayne and Marshall3 in their study of hot-water injection. The steam-injection process is quite complex, and an exact description of the process has not been found; however, a qualitative description can be made. When a quantity of steam is injected into a relatively cold reservoir, the leading edge of the steam front continaally flows into cool sand, and some steam condenses. As more and more steam condenses, a bank of condensed water will form which is augmented by displaced connate water. The reservoir may be divided in three regions—(1) a region containing displaced fluids at reservoir temperatures, (2) a region contain- ing water at reservoir temperature and (3) a region occupied by steam. The reasons for the lack of a region of hot water or oil of any appreciable size have been discussed.&apos; A method has been presented by which the location of the steam front can be determined even when heat losses to the regions surrounding the reservoir are sizable compared with the amount of heat injected. It remains to show the effectiveness of steam injection in increasing the ultimate oil recovery. To illustrate the technique for determining the increase in displacement efficiency and, hence, the increase in ultimate production, two sets of hypothetical reservoir conditions will be studied in detail.

By J. W. Eckerd, P. S. Lewis, N. H. Coates

USBM designed, constructed, and operated an 18-in.-diam fluidized-bed combustor for highly caking coals to evaluate the method for possible application to power generation. In initial tests, combustion was carried on successfully, with carbon losses as low as 0.8%+ Coal ash was not satisfactory as a bed material, but mullite proved to be acceptable. Overall heat-transfer coeffients from the bed (-20 t 50 mesh mullite) to a water-cooled tube were about 75 Btuper hr-sq ft. Fluidized-bed combustion is being investigated by the U. S. Bureau of Mines (USBM) as a potentially cheaper method of firing power plant boilers. Lower combustion temperatures in fluidized beds—only 1400" to 1800°F-—should reduce fireside deposition and corrosion and nitrogen oxides production. Fluidized-bed combustion also offers possible advantages in the use of additives to control SO, emissions and amenability to processing of high-ash, high-moisture coals without special preparation. This paper describes development of a fluidized-bed combustor and the results of preliminary tests. Equipment and Procedure Fig. 1 shows the 8-ft-high combustor. A cone-shaped plate perforated by I/8-in.-diam holes supports the fuel bed and admits the fluidizing air; water passing through the 3/4-in. pipe extracts heat from the combustion bed. Fig. 2 is a flow diagram of the system. Coal is fed pneumatically to the base of the fuel bed. Combustion products are passed through two centrifugal separators to remove most of the entrained solids, then to a water

Jan 1, 1971