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By J. Gurland

The strength variation ofWC-Co alloys with composition and particle spacing falls into two ranges. 1) Above a critical value of the mean free path, the strength follows a dispersion hardening relation; it is proportional to the volume fraction of WC and in-versely proportional to the particle size. 2) Below a critical value of the mean free path, the strength variation is based upon the Griffith criterion of fracture and is controlled by the case of crack propagation. The maximum strength is attained when the structure and load just satisfy the critical conditions for spontaneous crack propagation. THE relation between the strength and the structure of cemented carbides is of interest not only for its own sake but also for its contribution to the understanding of composite materials in general. The cemented carbides are essentially aggregates of a brittle carbide constituent and a ductile binder phase and, at certain structures of intermediate composition, the composite alloys attain strengths far higher than those of the separate constituents in bulk. It is this variation of strength with structure which will be discussed. The experimental procedures of sample preparation, mechanical testing, and quantitative metallo-graphic analysis have been described in Ref. 1. The reported strengths are the maximum fiber-stresses in bending calculated on the basis of completely elastic behavior. The plastic contribution to the over-all deformation is small and was neglected in the calculations (the elongation at fracture of the most ductile alloy, 50 pct WC-50 pct Co, is less than 0.5 pct).2 The pertinent structural parameters are: f, the volume fraction of tungsten carbide; p, the mean free path through the binder phase; the contiguity C, defined as the average fraction of surface area shared by a particle of wc with all neighboring particles of the same phase; and d, the average equivalent spherical diameter of the WC particles. The strength variation of WC-Co alloys with composition and particle spacing is shown in Fig. 1 and Table I, which are based partly on previously published data' and partly on new experimental work which extends the range to higher values of the mean free path. Each data point represents the average of five samples. For a given composition, the transverse rupture strength reaches a maximum

Jan 1, 1963

By PDF

C. 0. FAIRCHILD, Washington, D. C. (written discussion*).-Referring to the paragraph entitled "Effect of change in temperature of absorbing glass on its transmission," Dr. Foote and the writer have been using, since

Jan 11, 1919

By Clarence Zener

The heat-treatment of steels will not pass from the stage of an art into that of a science until the mechanism of the phase transformations associated therewith is thoroughly understood. Such an understanding will involve two distinct types of inquiry: (I) the direction of the transformation, (2) the speed of the transformation. The first type of inquiry, which involves only the equilibrium relations between the various phases in steel, is the subject of the present and of a succeeding article; the second type, which involves the kinetics of transformations, will be the subject of a later article. In order that a theory of equilibrium relations may give useful information, certain experimental data must be put into the theory. It is customary for these data to be of a thermal nature. In the system discussed herein—steel—such thermal data are very scarce. In medium-alloy steels another type of data might possibly be used. In these steels the concentration of solute is so low that it may be possible to use the approximation of dilute solutions, an approximation that leads to linear relationships. The unknown functions, or constants, may then be determined in some cases from the binary iron-alloy system and in other cases from the ternary iron-carbon-alloy system. From the functions, or constants, so determined may then be computed the equilibrium relations in more complex medium-alloy steels. In the present article this second approach is used, the necessary experimental data 'being obtained from the simple binary and ternary equilibrium diagrams. Fundamental Equations In this article are derived the equilibrium relations between the gamma, alpha and cementite phases. In these three phases all important alloying elements except carbon form substitutional solutions, each alloying atom occupying a potential site for an iron atom. Carbon, on the other hand, forms an interstitial solution. In the face-centered cubic gamma phase, the interstitial positions are at the center of the unit cube and at the centers of the cube edges. There are therefore four interstitial positions per unit cell, or one per iron atom. In the body-centered cubic alpha phase, the interstitial positions are probably at the face centers, and also at the centers of the edges, which are crystallographically equivalent to the face centers. There are therefore six interstitial positions per unit cube, or three per iron atom. In cementite, (Fe, Alloy)3C, the carbon atoms may. be regarded as filling all the lattice positions of a given type. The introduction of the fundamental the of equilibrium has presented the author with a quandary. These equations are formally so similar to other equations that have been derived1,2 for related systems that they may appear to some readers

Jan 1, 1947

By Clarence Zener

The heat-treatment of steels will not pass from the stage of an art into that of a science until the mechanism of the phase transformations associated therewith is thoroughly understood. Such an understanding will involve two distinct types of inquiry: (I) the direction of the transformation, (2) the speed of the transformation. The first type of inquiry, which involves only the equilibrium relations between the various phases in steel, is the subject of the present and of a succeeding article; the second type, which involves the kinetics of transformations, will be the subject of a later article. In order that a theory of equilibrium relations may give useful information, certain experimental data must be put into the theory. It is customary for these data to be of a thermal nature. In the system discussed herein—steel—such thermal data are very scarce. In medium-alloy steels another type of data might possibly be used. In these steels the concentration of solute is so low that it may be possible to use the approximation of dilute solutions, an approximation that leads to linear relationships. The unknown functions, or constants, may then be determined in some cases from the binary iron-alloy system and in other cases from the ternary iron-carbon-alloy system. From the functions, or constants, so determined may then be computed the equilibrium relations in more complex medium-alloy steels. In the present article this second approach is used, the necessary experimental data 'being obtained from the simple binary and ternary equilibrium diagrams. Fundamental Equations In this article are derived the equilibrium relations between the gamma, alpha and cementite phases. In these three phases all important alloying elements except carbon form substitutional solutions, each alloying atom occupying a potential site for an iron atom. Carbon, on the other hand, forms an interstitial solution. In the face-centered cubic gamma phase, the interstitial positions are at the center of the unit cube and at the centers of the cube edges. There are therefore four interstitial positions per unit cell, or one per iron atom. In the body-centered cubic alpha phase, the interstitial positions are probably at the face centers, and also at the centers of the edges, which are crystallographically equivalent to the face centers. There are therefore six interstitial positions per unit cube, or three per iron atom. In cementite, (Fe, Alloy)3C, the carbon atoms may. be regarded as filling all the lattice positions of a given type. The introduction of the fundamental the of equilibrium has presented the author with a quandary. These equations are formally so similar to other equations that have been derived1,2 for related systems that they may appear to some readers

Jan 1, 1947

By C. H. Li, J. Washburn, E. R. Parker

Yield stress in single crystals of zinc was shown to be dependent on prior annealing temperature and rate of cooling after annealing. Rate of strain hardening beyond the yield was not sensitive to annealing procedure. A tentative mechanism for the effect was discussed. COMPARATIVE mechanical tests of metal single crystals generally show a large scatter of results. Many factors may contribute to this variation. The present investigation was undertaken to evaluate the influence of two factors not ordinarily considered of major importance—annealing temperature and cooling rate after annealing. Since plastic properties of crystals are thought to depend on the presence of lattice imperfections, it follows that a variation in properties should result from differences in the distribution and density of these imperfections. Furthermore, in the light of current theories of interaction between edge dislocations and lattice vacancies,' it seems possible that annealing temperature and subsequent rate of cooling might play a part in this distribution. Analysis of the effect of these variables on plastic properties may lead to a better understanding of the annealed state. Single crystals of zinc which have been strained in simple shear undergo complete recovery of mechanical properties upon subsequent annealing. This behavior is useful because a single specimen can be used repeatedly, thus eliminating the uncontrollable variables introduced by the use of many crystals. The test section of the crystals used for these experiments was a cylinder 1/8 in. thick with a cross-sectional area of about 1/3 sq in. The specimens were acid-machined from spherical single crystals 1 in. in diameter in such a way that the slip plane (0001) was at right angles to the axis of the cylinder. In testing, a shear stress was applied along one of the slip directions, [2110] in the (0001) plane.' All crystals, unless otherwise noted, were grown in a helium atmosphere from molten 99.99 pct pure Horse Head Special Zinc. Fig. 1 is a family of curves obtained from one crystal. Each curve was obtained by deforming the crystal to 10 pct strain at a rate of 3 pct per min, following an anneal of 1 hr at 260°C. The strain direction was reversed every other time so that no progressive change in shape of the specimen took place during the series of tests. Since the stress-strain curve was reproducible within a fairly narrow range under the same conditions of annealing, it was then possible to determine the effect of changes in annealing conditions. A crystal was tested repeatedly as previously described but the annealing temperature and rate of cooling preceding each test were varied (Fig. 2). Two cooling rates were used, approximately 30°C per min and 3°C per min. Annealing was conducted at three temperatures, 200°, 300°, and 400°C. For a given annealing time and rate of cooling the room temperature flow stress was higher the higher the annealing temperature. The increase in flow stress was small as the annealing temperature was increased from 200° to 300°C but rose rapidly when the annealing temperature was raised from 300" to 400°C. Faster cooling from the high temperature produced an even greater strengthening of the crystal. The curves were obtained in the order indicated by the numbers. The stress level associated with a certain annealing treatment was, within the limits of scatter, always the same regardless of the

Jan 1, 1954

By John F. Elliott, David L. Schroeder

SURFACE depletion can be a significant source of error in the use of the effusion method for determining the thermodynamic properties of alloys, particularly in the case of solid alloys.' However, as yet this problem has not been treated quantitatively by experimentalists in the field.1-3 The method necessitates diffusive transport of material out of the source alloy to obtain a relevant measurement. The general character of the system is shown in Fig. 1. The diffusive flux (nD) at the alloy surface is equal to the effusive flux (nE): where D is the diffusion coefficient of the diffusing substance, C is its concentration, X is the distance from the surface into the source alloy, p is the density of the alloy, and AS is the actual available surface of the alloy. The concentration of the diffusing substance at the surface is CS, and Here fE is the Clausing factor, AE is the area (sq cm) of the orifice, y is the activity coefficient of the substance in the alloy, PO is the vapor pressure (Torr) of the pure substance, M is the molecular weight of the effusing species, and T is the absolute temperature. The differential equation for diffusive transport in the source alloy is Fick's second law. Its solution is dependent on the geometry of the source alloy. Two boundary conditions must be established to obtain the proper solution, the first being Eq. [1], or its equivalent. The solutions for several important forms of the source alloy are considered below.

Jan 1, 1967

[Murray and Renard: chondres on ocean floor, XLI, [155]. Musconetcong Iron Works, Stanhope, N. J., blnst furnace, XL, 467. MUSSEN, H. W.: Discussion on The Bogoslovsk Mining Estate, XXXIX, 897. death, XLIII, [Ixxvi]. Muzo, CoIombia: emerald deposits, LV, 910. emerald mine, XXXIX, 317. minerals, LV, 925.]

Jan 1, 1918

By Robert Sisselman

Dramatic changes have been effected since 1955 in the mode of iron ore shipments within the U.S. In 1955, less than one percent of ore consumed in domestic blast furnaces-approximately 125 million long tons-consisted of pellets. Five years later, in 1960, some 94.8 million tons of ore were consumed, 14% in the form of pellets. Projections for 1980 reinforce confidence in this most economic of blast furnace feeds, and set U.S. pellet consumption at 151.2 million tons. In the last decade, pellets caught fire in other countries. By 1969, world production reached 96 million tons. This compares strikingly with present world output of 154 million tons.

Jan 9, 1973

By E. L. Jr. Spain

THE deposits described in this paper occur over much of Wayne County and in the southeast portion of Hardin County, Tennessee, and in the northeast and northwest portions of. Mississippi and Alabama respec-tively. The presence of tripoli in this region has been known for some time and descriptions of various occurrences in Tennessee have been made by Jewell1 and Miser2, while Hilgard3, Lowe4 and Logan5 have treated the deposits that occur in the adjacent portions of Alabama and Mississippi, but there has not been a general survey of the entire field summarizing completely the geology of the deposits with respect to their origin and stratigraphic relationships. During 1934, while engaged in surveys of mineral resources through this region for the Tennessee Valley Authority, the writer collected data with this in mind, and has assembled them in this report. He is indebted to H. S. Rankin and N. A. Rose, of the Tennessee Valley Authority Geology Division, for field assistance.

Jan 1, 1936

THE developments in construction and operation of the basic open-hearth furnace that have taken place in recent years have placed new emphasis on the vital influence played by regulation and control of the combustion system in both furnace production and furnace life. The several factors influencing flame development, such as regulation of fuel supply, air supply, air preheat, furnace draft, and furnace reversal, must be controlled closely if successful furnace performance is to be accomplished. Rugged instrumentation of distinctive design is required to supply dependable information concerning operating conditions, which is requisite to proper regulation of combustion. This chapter attempts to give a practical discussion of the various features of open-hearth practice and instrumentation related to the efficient heating of the furnace, control of temperature, and protection of the furnace structure from accidental overheating. The appended references will be helpful to readers wanting a more detailed treatment of individual topics. OPEN-HEARTH FUELS The most important open-hearth fuels are listed in Table 4-1 together with their compositions and other data useful for combustion calculations. A detailed discussion of open-hearth fuels is not attempted here, but certain factors peculiar to their application to this type of process are briefly indicated on the following pages. Very often an open-hearth shop is forced by economic factors to use a fuel to which its furnaces are not very well adapted, because the general fuel-allocation scheme for the works as a whole seems to demand it. Often a variable marginal supply of coke-oven gas or tar must be consumed by the open hearth, together with another fuel such as producer gas, natural gas, or oil, to fill out total requirements. Most furnaces employ more than one kind of fuel during their usefu1 life; however, the partial redesign

Jan 1, 1951

YEAR OF ELECTION 1913. FRANK DAWSON ADAMS Montreal, Canada. 1921. WILLIAM CUTHBERT BLACKETT Sacriston, Durham, England. 1923. GELASIO CAETANI Rome, Italy. 1920. HENRY STURGIS DRINKER Merion Station, Pa. 1921. FERDINAND FOCH Paris, France. 1922. FEDERICO GIOLITTI Torino, Italy. 1906. SIR ROBERT A. HADFIELD London, England. 1921. FRANK WILLIAM HARBORD London, England. 1917. HERBERT HOOVER Washington, D. C. 1905. HENRI LECHATELIER Paris, France. 1921. C. MCDERMID London, England. 1913. EZEQUIEL ORDONEZ Mexico City, -Mexico. 1909. ALEXANDRE POURCEL Paris, France. 1911. ROBERT H. RICHARDS Cambridge, Mass. 1891. DR. ING. H. C. EMIL SCHROEDTER Mehlem a Rhein, Germany.

Jan 1, 1923

[YEAR of ELECTION 1913. FRANK DAWSON ADAMS Montreal, Canada. 1921. WILLIAM CUTHBERT BLACKETT Sacriston, Durham, England. 1923. GELABIO CAETANI Rome, Italy. 1920. HENRY STURGIS DRINKER Merion Station, Pa. 1921. FERDINAND FOCH Paris, France. 1922. FEDERICO GIOLITTI Torino, Italy. 1906. SIR ROBERT A. HADFIELD London, England. 1921. FRANK WILLIAM HARBORD London, England. 1917. HERBERT HOOVER Washington, D, C. 1905. HENRI LECHATELIER Paris, France. 1921. C. MCDERMm London, England. 1913. EZEQUIEL ORDONEZ Mexico City, Mexico, 1909. ALEXANDRE POURCEL Paris, France. 1911. ROBERT H. RICHARDS Cambridge, Mass. 1891. DR. ING. H. C. EMIL SCHROEDTER Mehlem a Rhein, Germany.]

Jan 1, 1928

YEAR OF ELECTION 1913 FRANK DAWSON ADAMS Montreal, Canada. 1921. WILLIAM CUTHBERT BLACKETT Sacriston, Durham, England. 1923. GELASIO CAETANI Rome, Italy 1920. HENRY STURGIS DRINKER Merlon Station, Pa 1921. FERDINAND FOCH Paris, France 1922. FEDERICO GIOLITTI Torino, Italy 1906. SIR ROBERT A. HADFIELD London, England. 1921. FRANK WILLIAM HARBORD London, England 1917. HERBERT HOOVER Washington, D. C 1905 HENRI LECHATELIER Paris, France 1921. C. MCDERMID London, England 1913. EZEQUIEL ORDONEZ Mexico City, Mexico 1909. ALEXANDRE POURCEL Paris, France. 1911. ROBERT H. RICHARDS Cambridge, Mass 1891. DR. ING. H. C. EMIL SCHROEDTER Mehlem a Rhein, Germany.

Jan 1, 1923

By B. J. Bayles, J. A. Ford, M. J. Salkind

Unidirecltonally sulidijied AL-A13Ni was found to exliibit excellent microstructural stability to 508°C. Above this temperature the transverse micvostructure coarsened in a manner analogous to "Ostwald Ripening", but the A13Ni whiskers did not spheroidize or foreshorlen. The coarsened whiskers reinforced the matrix as effectirely as the finer whiskevs and no reduction in room-temperature tensile strength mas noled after exposure for 90 hr at 96 pel 0-1 the absollite melling temperature of the eulectic. The slability was auribucted to the presence of low-index (and presurrlably lor-energy) cvystallographic interjaces between phases. The driving force for microstructural coursening was presumed to be the reduction of interfacial energy between the phases. Finer micvoslructures were found to coarsen more rapidly than those which weve imitally coarser. FIBER-reinforced metals have been shown to exhibit good short-time tensile properties at elevated temperatures.1-4 However, many elevated-temperature applications require long-time property retention which dictates the necessity for microstructural stability. Two types of instability are common to fiber composities. The first, chemical instability, is due to chemical reaction between the matrix and the relnforcing phase. Petrasek and weeton5 have shown that chemical reaction between tungsten wire and several copper alloy matrices resulted in degradation of the reinforcing phase and reduction in strength of the composite materials. Kreider and Leverant4 have noted similar behavior for boron-reinforced aluminum above 400°C. The second type of instability arises in systems in which the phases are essentially chemically stable with respect to each other and is characterized by spheroidization and/or agglomeration of the reinforcing phase. parratt6 has reported this type of degradation to occur relatively rapidly at moderate temperatures for nickel and cobalt alloys containing whiskers of silicon nitride, aluminum oxide, and silicon carbide and has identified it as "physicochemical" instability. Although both types of instability are the result of the differences in the chemical potential between the fibrous phase and the matrix, they may be differentiated by the fact that the second type, "physicochemi-call' instability, results in little or no net change in the composition or amount of each phase present. It does, however, result in a morphological change which profoundly affects the strengthening mechanism. Chemical interaction. on the other hand, results in a change in the composition and amount of one or more of the phases present, or the appearance of a new phase. In the latter case, the appearance of the mi-crostructure, that of a fibrous phase aligned in a matrix, may remain essentially unchanged, but the chemical degradation of the fiber results in markedly reduced composite strength. This investigation was initiated in an attempt to evaluate the microstructural and related mechanical property changes of a whisker-reinforced eutectic composite after elevated-temperature exposure. The composite system chosen for this study was the uni-directionally solidified A1-A13Ni eutectic which Lemkey. Hertzberg, and Ford7,8 demonstrated to behave as a whisker-reinforced composite. The microstructure of the unidirectionally solidified material consists of aligned needles of A13Ni in a matrix of aluminum. EXPERIMENTAL PROCEDURE Eutectic ingots having a nominal composition of 6.2 wt pct Ni were produced by melting 99.99+ pct A1 and 99.99 pct Ni in recrystallized alumina crucibles in an argon atmosphere. The analyses of the starting materials are listed in Table I. These master heats were cut up? melted by induction heating, and unidirectionally solidified vertically in 1/2 -in.-ID. 6-in.-long graphite crucibles, in a dynamic argon atmosphere. Therma1 gradients in the liquid were determined by measuring the temperature in the melt as a function of apparatus travel distance, as discussed by Kraft and Albright.9 The rate of solidification was assumed to be equal to the uniform rate at which the crucible assembly was withdrawn from the furnace. Ingots representative of nominal growth rates of 2, 5, and 11 cm per hr were produced and the solidification parameters are summarized in Table 11. Ingots used for met allographic analysis were pre-

Jan 1, 1968

By A. N. Diehl

IT is the object of this paper (1) to deal with the elements in blast-furnace gas from the standpoint of their importance, and the part they are to play in future consumption, and (2) to give detailed information about the construction and operation of a plant designed among the first in this country for washing gas on a large scale.

Jan 5, 1914

By R. W. Borio, T. S. Spicer

Inflammability and crossing point apparatus were used to study the influence of increasing amounts oi such inerts as ash, fly-ash, pyrite, limestone and Portland cement with several pulverized bituminous coals. The results indicate that it requires large proportions of inerts to substantially suppress the ease of ignition. The ignition phase of the combustion of solid fuels has been under study periodically since 1900. With the advent of pulverized coal firing some 60 years ago, a technique now dominating the coal utilization field, the ignitibility characteristics have become increasingly important. Much of the early work was directed toward the prevention of coal dust explosions and improved mine safety rather than combustion. Many factors affecting ignitibility have been studied and numerous methods devised in an effort to measure the ignition temperature and these have been reviewed by other authors. A recent survey was made by Finney and Spicerl in 1959, so a discussion of previous work will be omitted. Despite these efforts, the initial stage of combustion, termed ignitibility, is still a little understood phenomenon. In fact, no ASTM method has been adopted for the calibration of this illusive property. However, a D-5 Subcommittee XVI on Ignitibility of Coal and Coke of ASTM (now inactive) defined ignitibility:' "The ignitibility of a fuel is that char-

Jan 1, 1963

By E. L. Corp, J. C. Johnson, W. G. Pariseau, M. Poad, M. E. Fowler

This paper describes a comprehensive geomechanics case study of a full¬scale test stope at the Carr Fork Mine. The mine is owned by Anaconda Miner¬als Company and is located near Tooele, Utah. Large diameter blast holes were used in a new method known as blast hole post-fill (BHPF). The new design proved highly successful in contrast with earlier experience with vertical crater retreat (VCR) stopes. The main geomechanics objective was the estab¬lishment of a design procedure for optimizing stope and pillar dimensions with respect to productivity and recovery. An integrated three-part approach consisting of mine measurements, laboratory testing and finite element calcu¬lations was taken. Independent correlations of 0.88 and 0.89 between calculated and measured displacements and between estimated and observed yield :ones substantiate the finite element design approach using UTAH-II. Subse¬quent stope and pillar design analyses indicate that doubling of the original stope width at 50% extraction is quite possible without sacrifice of safety or stability.

Jan 1, 1984

By R. B. Yerxa, C. F. Green, H. O. Hofman

MASSACHUSETTS INSTITUTE OF TECHNOLOGY, In refining copper, the metal is melted down in a reverbera tory furnace in a more or less oxidizing atmosphere and then further subjected to an oxidizing smelting in order to eliminate the common impurities, most of which have a stronger affinity for oxygen than has copper. In these operations some of the copper is oxidized to cuprous oxide and dissolved by the metal bath. When the quantity of dissolved cuprous oxide has reached about 6 per cent, the metal is said to have been brought to " set-copper." A button-sample will show a depressed surface and, when broken, a single bubble at the apex of the depression; the fracture will be brick-red and dull. It is essential to carry the oxidation to this point in order to know that the impurities have been oxidized as far as it is possible under the given working-conditions. Nearly all the cuprous oxide of the set-copper is now reduced to the metallic state by poling, when "tough-pitch" copper will be obtained. A button-sample will show a flat surface. Upon breaking, it will be found that the former bubble has disappeared and that the fracture has become rose-colored and shows a silky luster. The quantity of cuprous oxide allowed to remain in the copper will vary with the impurities still present in the metal and with the degree of pitch that it is desired to reach. It is essential for the general physical and the mechanical properties of the resulting copper that such impurities as arsenic, antimony, bismuth, lead shall be present in the oxidized state, as they are

Jan 1, 1904

By B. F. Oliver, C. W. Haworth

Pure tin and Sn-0.5pct Pb ingots have been frozen unidirectionally from the base. For quiescent melts that were initially undercooled, a transition from lower eqlciaxed structure to an upper columnar structure is found in the alloy ingots. Columnar to equi-axed back to columnar transitions are observed in superheated alloy ingots, but no such equiaxed band is observed impure tin. The reproducible equiaxed band is associated with a thermal undercooling followed by a recalescence. This undercooling is <5"C, whereas the critical (maximum obtainable) under-cooling for both the pure tin and the alloys used is -20°C. A similar undercooling is observed at the same position in the pure tin ingots, although in this case no clear transition in structure can be seen. The structure of the pure tin ingots is either entirely columnar or mixed columnar-equiaxed. A consideration of the detailed thermal history of the ingots indicates that the ingot macrostructures are determined by the occurrence of a local therlnal undercooling in conjunction with nuclei multiplication and transport mechanisrris. GENERALLY it is found that a pure metal ingot solidifies so as to produce an entirely columnar structure. Frequently an alloy ingot is found to have a columnar outer zone and an equiaxed central portion. Early systematic work to examine the factors controlling the formation of the equiaxed structure was reported by Northcott&apos; who showed that, for copper alloys frozen unidirectionally with a given ingot practice, the alloying element influenced the length of columnar crystals and the extent of the equiaxed structure. Northcott showed that alloys with a wider freezing range more readily produced the equiaxed structure. The nucleation process can be important in producing equiaxed structures; frequently an alloy which readily solidifies with an entirely columnar structure will produce an entirely equiaxed structure when a nucleating agent is added to the melt.&apos; The formation of the equiaxed structure was attributed by Winegard and chalmers3 to the presence of constitutional supercooling; that is, a region of liquid in front of the growing solid could have a temperature below its equilibrium liquidus temperature. Thus, with a small enough temperature gradient in the liquid, it was suggested that the presence of constitutional supercooling may be sufficient to bring about the nuclea-tion necessary for the formation of an equiaxed structure. Although this explanation is plausible, and may be relevant in many ingots, Walker has described an experiment&apos; for which constitutional supercooling seems to be an unlikely cause of nucleation. A Ni-20 pct Cu alloy, repeatedly undercooled more than 50"C, was crystallized and found to show the typical colum-nar-equiaxed structure. The separation between the liquidus and the solidus for the alloy is 40°C. Thus, in this experiment the nucleation required for the formation of the equiaxed structure must have come about in some other way than by the nucleation catalysis constitutional supercooling hypothesis. Chalmers has suggested more recently5 that nuclei (in a typical ingot) are present immediately after pouring and are prevented from redissolving by the constitutional supercooling effect. More recently Uhlman, Seward, Jackson, and ~unt&apos; have shown direct evidence using ice and organic materials that freeze dendritically that the "remelt mechanism" may be an extremely effective crystal multiplication process during the freezing of ingots under conditions involving dendritic growth. JSlia" experimentally demonstrated the detachment of dendrite arms. chernov14 has analyzed the dendrite arm detachment process as a coarsening phenomena driven by the minimization of interphase area. Katta-mis and ~lemings" working with undercooled steel melts give evidence supporting this mechanism. Mechanisms of dendrite arm detachment such as those assisted by convection are believed to be the origin of the macrostructures obtained in this study. This study makes no attempt to distinguish the relative contributions of these mechanisms. The object of the present work was to obtain accurate temperature measurements during the solidification of an ingot and to correlate these measurements with the formation of equiaxed grains in the resulting ingot structures. Similar previous work is very limited. The measurements carried out by Northcott are neither sufficiently extensive nor sufficiently accurate for any interpretation. Plaskett and winegard7 carried out experiments on A1-Mg alloys in which they observed values of the temperature gradient, G, in the liquid and rate of freezing, R (for a given alloy solute content Co), at the transition from a columnar to an equiaxed structure. They reported that equiaxed crystals were produced at values of G/G approximately proportional to the solidus composition. Similar experiments using Pb-Sn alloys carried out by £111011" showed a linear relation between G/R and the solidus composition. However, the thermocouples were in the mold wall rather than in the melt and, in one case, ingot surfaces were examined. There is ambiguity in the meaning of the values of G and R measured in all these experiments. APPARATUS AND EXPERIMENTAL PROCEDURE Alloys were prepared by induction melting 99.999 pct Sn and 99.999 pct Pb to form a Sn-0.5 wt pct Pb alloy in air in a graphite crucible and casting into a cylindrical graphite mold 6 in. long, 1 in. in diarn , and with a & in. wall thickness. This mold was mounted on a copper base through which cooling water could be

Jan 1, 1968

By N. A. D. Parlee

The diffusion rate of hydrogen in liquid iron has been measured by a gas-liquid metal diffusion cell technique. The diffusion cell was formed by immersing an alumina tube containing hydrogen gas at 1 atm in a bath of stagnant liquid iron. The change in the composition of the melt in the cell was determined by measuring the rate of absorption of the gas in the cell. The appropriate solution to Fick&apos;s second law was used to examine the data and calculate diffusivi-ties. The absorption of hydrogen in stagnant pure liquid iron has been found to be diffusion-controlled. The results show that the chemical diffusion coefficient, D, of hydrogen in pure iron in the range of 1547" to 1726°C can be represented by the following Arrhenius relation: D(sq cnz per sec) = 3.2 x X exp(- 3300 i 1800/RT) where the uncertainty in the activation energy corresponds to the YO pct confidence level. Oxygen in the melt (above 0.015 pct 2) increased the apparent rate of absorption of hydrogen. The importance of diffusion data on liquid metals for predicting the rates of certain metallurgical processes has been recognized for a long time. Moreover, these data are much needed to test and develop theory for diffusion in liquid metals. Despite this practical and theoretical interest, however, relatively little reliable information about diffusion in liquid metals is available in the literature. This is particularly true for gas components such as hydrogen, oxygen, and nitrogen in liquid metals, where almost no data on chemical diffusion coefficients are to be found. This is probably due to a multitude of experimental difficulties particularly associated with high-temperature melts. In an effort to fill this gap in information, a research program was undertaken to study the diffusivities and rates of solution of gases in liquid metals. This paper presents the results of a study of the diffusion of hydrogen in liquid iron. EXPERIMENTAL METHOD Two methods for the study of the kinetics of dissolution of gases in liquid metals are being employed in this laboratory. Both involve the measurement of the volume of gas absorbed by the melt as a function of time and as such both avoid the uncertainties involved in chemical analyses of quenched samples for relatively small amounts of gas. In the first method, the gas dissolves in an inductively stirred melt and, in the absence of a slow surface reaction, the results are often interpreted in terms of mass transport across a liquid "boundary layer" between the homogeneous gas phase and well-stirred part of the melt. Other interpretations of the results of such experiments have also been described in the literature.1&apos;5 In the second method a gas-liquid metal diffusion cell is used.&apos; The gas dissolves in a cylindrical column of stagnant liquid metal and, in the absence of a slow surface reaction, the results are interpreted in terms of a non-steady-state diffusion solution to Fick&apos;s second law. The weakness of the first method is that while it gives information on the mechanism of absorption by stirred melts it yields an overall rate constant which even in the simplest cases depends on the nature and the thickness of the "mass transport layer" or "boundary layer". It yields no values of diffusion coefficients. The second method was used in this research because in many cases it is possible to determine the diffusion coefficient of the gas component in the liquid metal. In this research it has been utilized to measure diffusion coefficients of hydrogen in liquid iron. The apparatus used was essentially the same as that described by Mizikar, Grace, and par lee but certain modifications have been introduced to meet the elevated temperatures and special conditions of this research. Fig. 1 is a schematic drawing of the apparatus and Table I gives the identification of various parts in this figure. The diffusion cell, shown in detail in Fig. 2, was formed by immersing an impervious alumina tube (hereafter called absorption tube) in a bath of pure liquid iron contained in an alumina crucible. Two types of tubes were used, Morganite triangle RR and McDanel AP35. The crucible was contained in a vertical impervious alumina combustion tube (32 mm ID by 914 mm long) which was closed at both ends by water-cooled brass heads employing O-ring compression seals, Fig. 1. A protection tube enclosing a Pt, 5 pct Rh-Pt, 20 pct Rh thermocouple was introduced through the lower end of the combustion tube

Jan 1, 1968