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By Robert Anderson

ROUGHLY, a crack in a casting may be considered, for the moment, to be due to fracture of the alloy resulting from the stress set up by the contraction in volume on passing from the liquid to the solid state. Cracks may be external or surface, internal or deep-seated, or incipient or so-called strain cracks. Draws may appear as depressions, unevenness, and "pinching in" of the surface of a casting. Roughly, draws are due to contraction in volume and are governed largely by the same factors as cracks. A casting may be cracked, however, without being drawn. Shrinkage is a foundry term; it connotes the contraction in volume of a metal or alloy on passing from the liquid to the solid state. Pattern-makers' shrinkage is the total contraction in volume (including liquid, solidification, and solid shrinkage) that occurs when an alloy is poured into a mold. The contraction in volume of aluminum is one of the principal factors affecting the occurrence of cracks, as will be shown later, and the question of shrinkage and its control requires careful study in any treatment of cracking in castings. In the case of, sand and die castings, cracks have been attributed to a number of causes, but in the case of most light aluminum alloys, they are due to the large contraction in volume on freezing and the lack of its regulation or control. The most important factors are: (1) Contraction in volume; (2) composition of alloy; (3) quality of melting charge; (4) design of casting; (5) method of molding; (6) hardness of ramming; (7) method of gating; (8) hardness and characteristics of cores; (9) chills; (10) risers; (11) melting temperatures; (12) furnace used for melting; (13) pouring tempera-tures; (14) inclusions in alloys; (15) hot shortness of alloys; and (16) physical properties of alloys at high temperatures. Of these, the first, second, third, and eleventh to sixteenth involve ascertainable metallurgical facts; the remainder are largely matters of opinion, judgment and technical knowledge of men experienced in the production of castings.

Jan 10, 1921

By Ogden B. Malin

In conducting the manufacture of extruded brass rods and brass forgings it has been noticed that there is considerable variation in the physical properties, particularly the machinability of different rods and forgings, even though the composition, working temperatures and method of manufacture remain the same. It is the purpose of this paper to attempt to establish the cause and extent of this variation, and also to determine whether variations in composition within certain defined limits influence hardness and machinability to any marked degree. The writer believes that this variation probably is due to the variation in grain size of the metal, therefore the microstructures of several compositions under various conditions of working were examined in an effort to establish this as a fact. Survey of Previous Work in Studying Extrusion Defects Very little has been published concerning defects in extruded and forged brass. In 1921, R. Genders1 $ delivered a paper before the British Institute of Metals, relative to a defect inherent in the extrusion process. This defect, commonly found in extruded brass rod, shows up in the form of a circle in a transverse section consisting of oxides and impurities, and is generally present in the last 25 to 30 per cent of the rod. Genders concluded that the defect is caused by the turning inward of the skin of the billet at the edge of the plunger because the inside of the billet is more plastic than the outside, which is in contact with the cooler container; the outside skin will lag behind to form the rear end of the rod. Genders suggested a movable die as a remedy, and while this met with fair success, it is not used in present practice. About the same time as Genders' experiments, Tucker2 and Rosen-hain3 conducted experiments using colored plasticene in the form of a billet, and found much the same defect as Genders. Dr. Rosenhain State College, Pa.

Jan 1, 1932

Charles Catlett, Staunton, Va. (communication to the Secretary*):—Mr. Chance's suggestions that the brown hematite-ores of the Potsdam formation are due to the alteration in place of iron sulphide, and that their extension may reasonably he expected to contain valuable deposits of unaltered sulphide, are interesting; hut the facts marshaled by him are too limited

Jan 1, 1909

By R.J. Anderson

Roughly, a crack in a casting may be considered, for the moment, to be due to fracture of the alloy resulting from the stress set up by the contraction in volume on passing from the liquid to the solid state. Cracks may be external or surface, internal or deep-seated, or incipient or so-called strain cracks. Draws may appear as depressions, unevenness, and "pinching in" of the surface of a casting. Roughly, draws are due to contraction in volume and are governed largely by the same factors as cracks. A casting may be cracked, however, without being drawn. Shrinkage is a foundry term and refers to the contractio~l in volume of a metal or alloy on passing from the liquid to the solid state. Patternmaker's shrinkage is the total contraction in volume (including liquid, solidification, and solid shrinkage) that occurs when an alloy is poured into a mold. The contraction in volume of aluminum is one of the principal factors affecting the occurrence of cracks, as will be shown later, and the question of shrinkage and its control requires careful study in any treatment of cracking in castings. In the case of sand and die castings, cracks have been attributed to a number of causes, but in the case of most light aluminum alloys, they are due to the large contraction in volume on freezing and the lack of its regulation or control. The most important factors are: (1) Contraction in volume; (2) composition of alloy; (3) quality of melting charge; (4) design of casting; (5) method of molding; (6) hardness of ramming; (7) method of gating; (8) hardness and characteristics of cores; (9) chills; (10) risers; (11) melting temperatures; (12) furnace used for melting; (13) pouring temperatures; (14) inclusions in alloys; (15) hot shortness of alloys; and (16) physical properties of alloys at high temperatures. Of these, the first, second, third, and eleventh to sixteenth involve ascertainable metallurgical facts; the remainder are largely matters of opinion, judgment, and technical knowledge of men experienced in the production of castings.

Jan 1, 1923

By R. J. Anderson

Roughly, a crack in a casting may be considered, for the moment, to be due to fracture of the alloy resulting from the stress set up by the contraction in volume on passing from the liquid to the solid state. Cracks may be external or surface, internal or deep-seated, or incipient or so-called strain cracks. Draws may appear as depressions, unevenness, and "pinching in" of the surface of a casting. Roughly, draws are due to contraction in volume and are governed largely by the same factors as cracks. A casting may be cracked, however, without being drawn. Shrinkage is a foundry term and refers to the contractio~l in volume of a metal or alloy on passing from the liquid to the solid state. Patternmaker's shrinkage is the total contraction in volume (including liquid, solidification, and solid shrinkage) that occurs when an alloy is poured into a mold. The contraction in volume of aluminum is one of the principal factors affecting the occurrence of cracks, as will be shown later, and the question of shrinkage and its control requires careful study in any treatment of cracking in castings. In the case of sand and die castings, cracks have been attributed to a number of causes, but in the case of most light aluminum alloys, they are due to the large contraction in volume on freezing and the lack of its regulation or control. The most important factors are: (1) Contraction in volume; (2) composition of alloy; (3) quality of melting charge; (4) design of casting; (5) method of molding; (6) hardness of ramming; (7) method of gating; (8) hardness and characteristics of cores; (9) chills; (10) risers; (11) melting temperatures; (12) furnace used for melting; (13) pouring temperatures; (14) inclusions in alloys; (15) hot shortness of alloys; and (16) physical properties of alloys at high temperatures. Of these, the first, second, third, and eleventh to sixteenth involve ascertainable metallurgical facts; the remainder are largely matters of opinion, judgment, and technical knowledge of men experienced in the production of castings.

Jan 1, 1923

By Robin O. Williams

Robin 0. Williams (Oak Ridge National Laboratory)— The authors have questioned the degree to which the coherency strains between the iron-rich and chromium-rich phases are isotropic as proposed in Ref. 5 on the basis of the difference between the elastic properties of the two phases. The relative magnitude of the stresses is determined by the moduli as shown by Eqs. [2], [3], and [4] of Ref. 34. However, the moduli of the two phases have no direct bearing on the uniformity of either the stress or strain within either phase. The idea that the strains are isotropic within each phase (but normally of different magnitude and always of different sign) is based entirely upon the experimental observation that X-ray line broadening has not been detected even when the particles become rather large. It has not proven possible to grow the particles sufficiently large that they lose coherency. Based upon this lack of line broadening one can estimate an upper limit for the nonuniformity of the strains within each phase as follows. It is considered possible to detect line broadening if it is as great as 10 pct of the separation of the K, doublet for the (211) line using chromium radiation. The doublet separation would correspond to a total strain of 0.0017 such that the total variation of lattice parameter relative to the average lattice is now k0.05x0.0017 or something less than ± * For the present case the strain in each phase is roughly 0.002 such that the variation of strain within a phase will not exceed 5 pct. It is stated that the expression derived for strengthening for the hydrostatic straining as observed in this system would substantially overestimate the magnitude due to dislocation flexure. This is contrary to the conclusion reached in the original paper34 for the present range of particle sizes. What is the lowest temperature at which a has been observed to form in this alloy? M. J. Marcinkowski, R. M. Fisher, and A. Szirmae (nutlzors' reply)— -Williams' arguments based on X-ray findings for a chromium-rich precipitate and an iron-rich matrix strained to a common lattice parameter are certainly convincing. This being the case, there are no shear components of strain associated with the precipitate-matrix aggregate to interact with the shear components of the dislocation stress fields, contrary to the opinion expressed by the present authors. On the other hand, the present authors, in spite of this error, did not expect the shear interactions to be significant. The chief objection to Williams' model in the present case is that the various segments of the dislocation line are assumed to pass from one potential valley to the next independently of neighboring segments. This is only true for a highly flexible dislocation line, i.e., one whose radius of curvature is something less than the center to center distance between precipitate particles which amounts to about 90A in the present alloy. In order to maintain this curvature, an externally applied shear stress of at least 230,000 lb per sq in. would be required or about four times the observed stress. It is therefore concluded that the dislocation lines move rather rigidly through the lattice. This being the case, the forces on the dislocation resulting from the hydrostatic interaction between the stress fields of the edge-dislocation components and the precipitate particles should average out to zero; that is particles above the below the slip plane produce forces on the dislocation of opposite sign and therefore will cancel when averaged over the entire length of the dislocation. On the other hand, since the dislocation is not perfectly rigid, Williams' model may lead to some strengthening, but far less than that predicted. A second and equally serious objective to using Williams' strengthening model for the present alloys is that profuse wavy slip due to the motion of screw dislocations played a predominant role not only in the unaged alloys but in the fully aged ones as well. Since the screw dislocation has associated with it only shear components of stress the hydrostatic strengthening model no longer applies. In view of these arguments the present authors must reject Williams' model of strengthening as being pertinent to the present alloy system. The present authors have made no detailed study of the lowest temperature at which a forms in the quenched ferritic alloys. None was ever observed n the alloys aged at 500°C so that forma-tion must occur at temperatures higher than this and was therefore not a factor in the present study.

Jan 1, 1965

By M. B. Bever, A. B. Michael

The solidification of aluminum-rich aluminum-copper alloys was investigated for different solidification rates. The measured amounts of nonequilibrium eutectic were compared with the amounts calculated on the assumption of no diffusion in the solid. The morphology of the eutectic was studied and dendritic spacings were measured. Compositions of the cored primary solid solutions were determined by quantitative autoradiography after activation of the solute copper by neutron irradiation. CONSIDERABLE progress has been made in recent years toward an understanding of the solidification of metals. This is especially true for the factors involved in nucleation, the controlled growth of metal crystals from melts, and various aspects of the freezing of ingots and castings. Many features of the solidification process, however, are not understood adequately and much further work is required. In the research reported here the solidification of a series of aluminum-rich aluminum-copper alloys was investigated. Different degrees of deviation from the idealized case of equilibrium solidification were attained by varying the rate of solidification. Thermal analysis and microscopic examination were supplemented by lineal analysis and autoradiography. The latter required the use of radioactive tracers and was adapted to the quantitative determination on a microscale of concentrations of copper dissolved in aluminum. The tracers were produced in situ by activating the copper with neutrons. These techniques permitted a correlated interpretation of the temperature-time history, the amounts of nonequilibrium eutectic and the general morphology of the alloys, and yielded quantitative data on microsegregation. Nonequilibrium Solidification In the general case of the solidification of a solid solution in accordance with the phase diagram, the composition of the coexisting liquid and solid must change continuously. This change requires diffusion, which at best remains incomplete in the solid during solidification. The resulting inhomogeneity of the solid solution crystals is well-known as coring, dendritic segregation, or microsegregation. The generally accepted analysis of coring assumes that, 1—no diffusion occurs in the solid, 2—diffusion is complete in the liquid, and 3—the composition of the solid formed on cooling through any infinitesimal temperature interval is given by the solidus of the phase diagram. On this basis, In this equation m designates mass, x designates the concentration of solute and the subscripts L and S refer to the liquid and solid phases, respectively; the subscript o indicates the initial state, in which the liquid is identical with the entire system. Eq. 1 or its equivalent has been derived repeatedly'-7 and an analogous equation, known as Rayleigh's equation, has been used for distillation. The assumptions on which Eq. 1 is based do not include an infinitely fast cooling rate as sometimes stated. In fact, Olsen and Hultgren %ave shown by experiment that very high solidification rates suppress coring, presumably by preventing diffusion in the liquid and by the effect undercooling has on the composition of the nucleus. Coring thus occurs at rates of cooling which are too fast for a close approach to equilibrium in the solid and too slow to suppress diffusion in the liquid and to cause marked undercooling. If the liquidus and solidus approximate straight lines, Eq. 1 simplifies to The fraction of liquid present at any temperature during nonequilibrium solidification can be calculated from Eqs. 1 or 2. By the lever relation it is

Jan 1, 1955

By J. L. McK. Yardley

There are two general classes of problems in connection with electric furnaces. First, those relating to the best utilization of the electrical power inside the furnace; second, those connected with the bringing of the electrical power to the point where it is to be utilized. Referring particularly to the latter class of problems, many calculations have been made, covering a wide variety of furnaces, to check the electrical capacity and power factor found from test. The following analysis, made for the purpose of determining the maximum capacity and approximate performance of a new furnace, to operate at 160 volts on a 60-cycle circuit, may be of interest. The first characteristic necessary to determine is the inductance, since this is the factor that limits the amount of power which can be put into the furnace. Inductance is calculated from the dimensions of the furnace, according to formulas of the U. S. Bureau of Standardsll as follows. Formula 99, for mutual inductance, is: M = 2l [log 2l/d-1 +d/l] Formula 103, for self-inductance, is: Both of these formulas are approximations which are nearly correct when 1 is great compared to d or R. 1 = length of circuit, in centimeters, d = separation between the conductors, in centimeters, R = geometric mean radius of conductors. For a circular tube, R = the radius of the circle. For a solid circular conductor, R = 0.7788 X the radius of the conductor. For a rectangular conductor, R = 0.2235 X (a + b), a being the length, and b being the width of the rectangle. To calculate the inductance of a three-phase circuit, it is customary to figure the inductance of one leg. In circuits where the three conductors

Jan 1, 1920

By Frank Harwood Riddle

Products made from the ores of the sillimanite group, and synthetic substitutes for them, have unique properties, and service tests prove that they are playing, and will continue to play, a major part in improvements in metallurgical processes. Also, they are examples of how comparatively useless, rare, or "museum" minerals have been made available in commercial quantities. Until research developed uses for them, the minerals had been observed only in small, scattered masses, or crystals. Improvements in ceramic processes and compositions have come about very gradually. In the last 10 or 15 years, several outstanding developments have been made. Of greatest interest to the mining and metallurgical engineers are those affecting the composition, necessitating the mining of new minerals, and improvements in ceramic products used in metallurgical processes. Sillimanite and Mullite Sillimanite (Al2O3.SiO2), then known as Fibrolite, according to Dana,' was first reported in 1792. In 1796, cyanite (Al2O3.SiO2) was first described.2 In 1873, Behrens published results of his microscopic study8 of porcelain. He showed that it contained what appeared to be crystals, which dissolved less readily than quartz and glass in hydrofluoric acid. Hussak4 described the presence of needlelike crystals as occurring in some instances in small amounts in certain porcelains. In 1908, Plenske and Zoellner both studied porcelain structure.5 Plenske observed that "sillimanite" might be present in masses of exceedingly minute grains which apparently were amorphous. Zoellner advanced the theory that the needlelike crystals were formed by a molecular change in the clay,

Jan 1, 1932

By S. Luszcz, R. W. Vook, Fred Witt

In situ X-ray investigations were made on polycrys-talline silver films deposited by vacuum evaporation on (111) CaF2 and (100) NaCl single-crystal substrates at 80°K. The films were evaporated and annealed in an X-ray diffractometer attachment having a residual gas pressure of 2 x lo-' Torr. All measurements were made without exposing the films to the atmosphere. Measurements were made on the films in the as-deposited state and after various annealing treatments. The intrinsic stacking and twin fault densities, the magnitudes of the uniform and nonuniform strains, and the crystallite sizes were determined. In addition the textures in the films were measured qualitatively. The results obtained for the as-deposited films on single-crystal substrates are in substantial agreement with previously reported results for silver films deposited on glass. Intrinsic stacking and twin faults, as well as uniform and nonuniform strains, were present in these films. During the various annealing treatments (up to 350°C) the faults and nonuniform strains annealed out. Considerable grain growth and texture changes occurred also. The effects were much greater for the NaCl substrate than for the CaFz substrate. The relative magnitudes of the grain growth in the variously oriented grains could be explained qualitatively in terms of the thermal strains and strain energies introduced into the differently oriented grains during the initial, irreversible anneal. These strains were due to the different thermal expansion coefficients of the film and substrate. X-RAY diffraction measurements on evaporated films deposited on substrates at low temperature have the advantage that many of the imperfections introduced into the film during deposition are "frozen in". Thus, the influence of a very important experimental variable, substrate temperature, on the imperfection structure of evaporated metal films may be studied. Moreover, the effects of annealing such films makes possible the study of thermally activated recovery processes in these films. The present study was designed to determine the influence of single-crystal substrates on the resultant film structure relative to the previous results obtained using glass substrates.' To this end great care was taken to keep the experimental variables the same in the two cases. Different experimental conditions would, of course, result in films having different physical properties. Again the initial substrate temperature was in the neighborhood of 80°K and the films were subsequently annealed to 350°C. The pure metal silver was chosen for evaporation, primarily because of its relatively low stacking fault energy and consequent high fault density in the as-deposited state. The silver films were formed by evaporation onto air-cleaved {ill} CaF, and (100) NaCl surfaces cooled to 80°K in an X-ray diffractometer attachment2 having a base residual gas pressure of 2 X l0-' Torr. The films were not exposed to the atmosphere until all of the X-ray data had been recorded. In this way one of the most important experimental variables, environment, could be well-controlled and reproduced. X-ray measurements were made at the temperature of deposition and included determinations of the diffraction line peak positions, line shapes, and integrated intensities. The peak position measurements were used to determine the intrinsic stacking fault densities and the average uniform strain in the film. The shapes of the diffraction lines provided information on the twin fault density, true crystallite size, and average nonuniform strain. The preferred orientation in the film was determined qualitatively from the integrated intensities. I) EXPERIMENTAL PROCEDURE The evaporator attachmentZ was charged with 99.999 pct Ag pellets positioned in a tantalum filament which had been outgassed previously at l0-8 Torr. The CaFz and NaCl single crystals were cleaved in air and then placed in position in the chamber so that their cleavage surfaces were on the diffractometer axis. The chamber was prepumped using a sorption pump, sealed off, and then baked at 150°C for 24 hr. The ion pump operated during the bakeout cycle. The substrate was then heated to 500° C by means of an auxiliary heater and kept hot until the rest of the chamber was cooled slowly to room temperature. This bakeout procedure consistently resulted in an ultimate pressure in the low lo-' Torr range. The substrate was then cooled down on 80°K. Its temperature was monitored by a thermocouple wedged into the rear of the copper substrate holder. The diffracted intensity and peak position of the 111, 222, and 333 CaF, lines were measured prior to evaporation. Nickel-filtered, pulse-height-discriminated copper radiation was used. Similar measurements were made for the 200 and 400 lines from NaC1. These measurements were used as a lattice parameter check and to determine the thickness of the evaporated silver films from the attenuation of the substrate lines. The evaporation rates were approximately 3A per sec for both films while the maximum pressures during evaporation were 3 x lo- ' and 7 x 10"8 Torr for the CaF, and NaCl cases, respectively. The film thickness was measured by the attenuation of the CaF, and NaCl substrate lines and by at optical interference method. Values of 1700 and 1500A, respectively, were obtained for the silver

Jan 1, 1969

By H. E. Cline, D. F. Stein

A fine lamellar structure with interlamellar spacings ranging from 0.I to 3.7 was Woduced by high-speed directional solidification and used to study the interaction of dislocations with an array of bimetallic bound-am.es. The tensile flow properties were measured at elevated temperatures and the results interpreted in terms of an image force model. As the width of the plates decreases the strength increases according to a Hall-Petch relation, suggesting pile -ups. Considerable strengthening was also observed at elevated temperatures. In recent years composites containing a strong but brittle phase and a ductile but weak phase have been studied extensively in the search for a high-strength material with adequate ductility.' However, the strength of these composites has been limited to a volumetric average of the strengths of the component phases. In this paper we will show that a composite made by directional solidification of a Cu-Ag eutectic has strength properties that substantially exceed the strength of either of the metals in the composite, and an explanation for this increased strength will be presented. The Cu-Ag eutectic is a near ideal material to study for several reasons. The eutectic structure is made up of alternate plates of silver-rich and copper-rich phases with the same crystal structure.2'3 In addition, the mechanical and elastic properties of silver4 and copper5 are known over a wide range of temperatures and the eutectic microstructure is stable to nearly the melting point.6'7 Therefore, the properties of the eu-tectic are readily comparable to various theories involving the mechanical properties of copper and silver. The experiments consisted of producing directionally solidified Ag-Cu eutectic material, determining the structure and performing tensile tests at different temperatures and strain rates. EXPERIMENTAL A) Solidification and structure. A cylindrical ingot 2.2 cm in diam and 15 cm long was prepared from 99.999 pct pure materials and cast in argon. The ingot was placed in an alumina crucible and was direc-tionally solidified at 7 x 10"4 cm per sec in a Bridg-man apparatus in argon. Metallographic examination at both ends and along the length of the ingot showed a uniform structure that consisted of fine alternate plates of silver- and copper-rich phases. The plates were aligned parallel to the growth direction, as shown in the longitudinal section, Fig. 1, and were in small randomly oriented domains of parallel plates as viewed in a transverse section, Fig. 2. Four smaller ingots 0.175 in. in diam were solidified at speeds of 2Xlo-', 4.4 xlo-', 4.4 x 10-3, and4.4x10-4 cm per sec to produce a range of lamellar spacings. Replica electron micrographs were taken to resolve the individual plates as shown in Fig. 3. The copper-and silver-rich plates in the larger ingot were about

Jan 1, 1970

By Perry Olcott .

THE Taka oil field is in the northern part of Titus and Franklin Counties, in north- eastern Texas. The small town of Talco, from which the field derives its name, is on the P. & Mt. P. R. R., about 15 mi. north of Mt. Pleasant, the county seat of Titus County. The oil field extends east and west from the town of Talco, the approximate center. In May, 1924, R. L. Cannon, while doing surface geology for the Humble. Oil & Refining Co., mapped the surface trace of this fault structure as evidenced by Upper Midway faulted down against Lower Midway. In the interval of sixteen years between the original surface work and the recent productive well completion, approximately eight wells were drilled into the Woodbine sand without success. The structure was well defined by those tests and all of the wells en-

Jan 1, 1936

By Donald H. Beilstein

The St. Joe Minerals Corporation Lead Smelter and Refinery have been in continuous operation at Herculaneum, Missouri since 1892 as the Herculaneum Lead Smelting Division. Over the years, the plant has undergone many changes to increase production, improve the technology, reduce the cost of operations, and produce the highest grade of primary lead. The latest alterations were begun in 1950 and completed in 1970 and have resulted in the Herculaneum Smelter capacity being increased to in excess of 200,000 tons per year (180,000 metric tons). This increase has made the Herculaneum Smelter the largest in the United States, and second largest in the world. This paper will describe the Herculaneum Smelter equipment, processes, and material balances. Particular emphasis will be given to the McDowell-Wellman 10' x 100' (3 x 30 meter) up-draft sinter machine and ancillaries; to the new Heavy Smelting Department (blast furnaces and drossing section); to the new Lead Refinery and its automatic casting equipment and to the gas handling system including the 340 tons per day (305 metric) Sulfuric Acid Plant utilizing a -four pass converter. The description will include photographs, flow sheets, tonnages and assays where applicable.

Jan 1, 1970

By S. W. Ihle, L. J. Warren

A sample of garnet from the scheelite ore body on King Island, Tasmania, Australia, was upgraded and ground to produce an ultrafine portion suitable for micro-electrophoresis. X-ray and chemical analysis identified the garnet as andradite with some admixed grossular. At neutral pH the garnet particles were negatively charged, their zeta potential becoming rapidly more negative with increasing pH. An iso-electric point was found at pH 3.5. The effects on the zeta-potential of various electrolytes, hydrolyzable metal ions, and flotation collectors were also studied.

Jan 1, 1976

By Henry M. Otte, Ralph P. I. Adler

The changes of line position and integral line breadth in the X-ray diffraction pattern of a polycvys-talline Cu-30Zn tensile test piece, incrementally loaded (and unloaded) up to fracture, have been an-alyzed in detail. The stacking-fault probahility, cv, increased linearly with increasing strain, E, wheveas the effective domain size, De(hkl), decreased with decreasing E-1 Over the greater part of the stress-strain curve the rate of work hardening was essentially constant (about 86 kg per sq mwz), and could be correlated with the slope of stage II of the single-crystal stress-strain curve. Consequently the theories of work hal-dening (particularly those parts relating to stage 11) as developed by Mott and Hivsch and others could be applied to the observations made on the polycrystalline brass. A relationship of the form Aa = Aao - MhklEhkl between the change, Aa, in the extrapolated lattice pararneter and the rvns strain, Ehkl, was derived and found to fit the results acceptably well. From this and other relationships developed in the papev it was estimated that the equilibrium stacking-fault energy of Cu-30Zn was between 8.4 and 12.5 ergs per sq cm, in fuirly close agreement with the (corvected) value obtained by Howie and Swann (1961)43 using transmission electron microscopy. The theory of work hardening in the jorm developed and recently presented by Hirsch (1964)3 successfully described all the pvesent observations. In order to test certain aspects of the theories of work hardening, as developed by Mott,1 Hirsch,2,3 Seeger et el,4-7 and others (for review see Nabarro, Basinski, and Holt8), several recent investigations have been concerned with relating the dislocation density, p, with the shear stress, 7 (and strain, y), applied to the specimen. The results of these investigations have shown that the square root of the dislocation density appears to be linearly related to the applied shear (or flow) stress for fcc as well as bcc metals and alloys. Furthermore, the relationship appeared to apply not only to the deformation of poly-crystalline specimens, but also to stages I and I1 of the deformation of single crystals. An expression of the form has thus come into wide use. Here b is the Burgers vector for a total dislocation, G is the shear modulus, and 70 and q are constants. A review9 of available values of q shows it to have values (at room temperature) in general between 0.3 and 0.6. Forms of Eq. [1] can be deduced from, or predicted by, the current theories, and the various constants adjusted so that they are compatible with the experimentally found value of q . No unique relationship has yet been found between the dislocation density and the applied shear strain. There are several serious objections to the use of Eq. 11]. In the first place, it relates the shear stress to the density of the dislocations without regard to their arrangement, type, or distribution; the significance of the relation may therefore be justly questioned.5 In the second place, the values of the experimental quantities usually substituted into Eq. [11 are those of the applied shear stress and the total dislocation density measured after unloaditzg. The dislocation density value that should in fact be used is that for the mobile dislocations present in the specimen when under the applied load.* Finally, in cases where the values used for p, the dislocation density, are those obtained by electron microscopy, p is subject to considerable error,' both systematic and random. The corrections to be applied are still controversial. Dislocation densities can also be measured by etch-pit and other techniques,'' each having their specific limitations. An objective of the present investigation has been to obtain information about the dislocation configyration and distribution by analyzing the changes in the position and shape of X-ray diffraction profiles as a function of deformation. The X-ray techniques employed, also open to criticism, have certain advantages, however. Thus, although the X-rays diffract only from the surface layers to an effective depth of about 20 p, the measurements can be made while the specimen is under load. The value of the dislocation density obtained by the X-ray method is also subject to errors, which are different from those of the electron microscope. Though a considerably larger volume of material is sampled by the X-rays, thereby reducing some of the statistical errors inherent in the electron microscope data, the information obtained is less detailed and is dependent on the method of analysis used to obtain a value for the dislocation density. Nevertheless, important observations can be made because the aforementioned advantages outweigh some of the limitations. In the present paper the X-ray method is briefly described and applied to a brass specimens deformed in tension. The results are then discussed in terms of some of the current concepts of work hardening. 1) EXPERIMENTAL PROCEDURE Details have already been extensively published elsewhere11-14 and therefore will only be dealt with briefly here. 1.1) Materials and Specimen Preparation. Commer-

Jan 1, 1968

By M. W. Bartley, Hugh M. Roberst

ERRATA In the paper (Mining Technology, January 1943) Page 19, second column. Line 5 should read: of Lower Huronian and Early Middle Huronian. Page 20, second column, line 17. Read magmatic solutions instead of magnetic solutions. In the discussion (Mining Technology, September 1943) On page 2 of this discussion pamphlet there is a typographic error. In the chemical analysis, the first reaction should read: CaCO3 + H2O = CaO2H2 + CO2 AMERICAN INSTITUTE OF MINING AND METALLURGICAL ENGINEERS

Jan 1, 1943

By A. L. Field, P. H. Royster

The first report on the slag viscosity work of the Bureau of Mines was made by one of the authors1 in 1916. It was concerned chiefly with the method of measurement. A paper2 on this phase of the work was presented to the Institute in February, 1917, and a similar paper3 was communicated to the Faraday Society. PuRpose and Scope of the Slag Viscosity Tables Following the development of a suitable method, the next step in the investigation was the measurement of slag viscosity over a wide range of temperature and composition, and the arrangement of this experimental data in the form of tables for the use of the blast-furnace operator. It is believed that this information, if used intelligently, should help the operator to reduce losses in production caused by off-grade pig iron, to improve fuel economy, to promote operating efficiency, and to extend present-day practice to meet the increasing need of smelting lean and complex ores. In the experimental part of the work it was carly, observed that the temperature-viscosity relations of a slag were determined, for practical purposes, by the relative proportions of the three major constituents, lime (CaO), alumina (Al2O3), and silica (SiO2), provided the percentage

Jan 1, 1918

By Don M. Jackson, Robert W. Howard

Techniques for the epilaxial growth of single -crystal silicon carbide films on silicon were developed. The vapor-phase decomposition and bydrogen reduction of silicon tetrachloride (SiC14) and Propane (C3H8) resulted in clear films of silicon carbide, lip to seveval microns in thickness. The growth took place in a horizontal . silicon epilaxial reactor at 1100°C (pyrometer) at a rate of- 3000Å per minute. Electron diffraction and X-ray diffraction studies demonstrated that the films were single-cyrstal, ß -phase, or cubic silicon carbide. SiO2 film were used to mask areas of the silicon sur-lace in order that the silicon carbide might be grown in controlled geometries. Both n- and p-type films were grown on p-type silicon waters.. Heavily doped silicon films of the same conductivity type as the silicon carbide films were deposited over the silicon carbide in order to affect better probe contact to the structures. when n-type silicon carbide mesas were grown on p-type silicon substrates the de vollage-current relationships between films and substrates were that of junction diodes. These diodes showed a sensitivity to while light ill that the incident light increased forward- and reverse-satro,ation currents, P-type silicon carbide mesas grown on p-type silicon were ohmic rather than rectifying in their voltage -current relationship. No conclusions could he reached concerning heterojunc-tiou rectification in the structure. SILICON carbide is a semiconductor with many interesting properties. It decomposes at temperatures above 2200°C.1 It occurs in two general crys-tallographic forms—hexagonal (a Sic) and cubic (ß Sic)—with the cubic form having a forbidden-gap energy of 2.32 ev and the hexagonal form (specifically the 6H polytype) a gap energy of 2.86 ev.3 It behaves as an extrinsic semiconductor at temperatures approaching 5003C. It has been shown to have a high resistance to radiation damage4 and p-n junctions formed in Sic have been shown to radiate visible light under forward- or reverse-bias conditions. Epitaxial silicon carbide on silicon carbide has been successfully grown through the use of a variety of techniques, such as gaseous cracking of SiCL4 and CC4, nearly all of which require a deposition temperature above 1500°C.6 This paper will cover very recent work on the gas-phase deposition of highly ordered films of silicon carbide on high-quality silicon single-crystal substrates. The films have been shown to ex- hibit junction-rectification properties when geometrically isolated regions are electrically biased with reference to the silicon substrate. There will be no discussion of the mechanism of heterojunction rectification, but the methods of film fabrication, geometry control, and structural evaluations will be covered in detail. Electron diffraction, X-ray diffraction, and diode electrical properties were used to characterize the films and the junctions. GAS-PHASE DEPOSITION OF Sic The techniques for the deposition of silicon carbide films were a logical outgrowth of the standard silicon epitaxial process. The major premise followed was that, for any film to nucleate in an ordered fashion where there is considerable mismatch in lattice parameters (in this case 22 pct), an extremely clean, damage-free substrate surface must be presented to the gas stream. Thus a standard gas-phase HCl etching step was used to prepare the substrates for growth. A minimum of 5 µ of substrate-surface material was removed prior to the deposition of Sic overgrowth films. The techniques used for growing silicon carbide films were those of growing silicon alone, with the added injection of a hydrocarbon gas into the hydrogen and silicon tetrachloride gas stream. The hydrocarbon gases used thus far have been research-grade (99.99 pct) methane (CH4) and propane (C3H8). Propane ultimately gave the best results. The gas flows were controlled through a panel shown schematically in Fig. 1. A hydrogen main stream of 30 liters per min passed through the horizontal quartz-tube epitaxial reactor, while SiC14, C3H8, HC1, and doping gases were injected as side streams. The

Jan 1, 1965

By John Birkinbine

This piece of Carpenter crucible steel has exhibited a tensile strength of 116,000 pounds per square inch, an elastic limit of 89,170 pounds, an elongation of 25 per cent., and a reduction in area of 42 per cent. The test-specimen had a length of 2 inches between fillets, and was an inch bar reduced to an area of 0.504 square inch. This unusual tensile strength was determined before the bar had been hardened so as to show the remarkably velvety fracture here exhibited; but a piece of the same, tested after hardening, is reported to have stood a stress of 214,000 pounds per square inch. Another test-piece is said to have shown 136,000 pounds tensile strength, 97,300 pounds elastic limit, 20 per cent. elongation in 2 inches, and 40 per cent. reduction of area. Mr. Carpenter claims that this steel, containing 2.27 per cent. chromium and 1.04 per cent. carbon, can be welded like iron, and has proved most desirable,for the highest class of die-steel.

Jan 1, 1895

By E. S. Machlin, A. A. Hendrickson

Use was made of a recently developed surface-accumulation diffusion technique to measure the self-diffusivity of edge-type dislocation singular lines (Burgers vector along <110>) in a bent and polygonized single crystal of silver. Two factors are involved in obtaining a number for the diffusivity. One factor, the dislocation-line density, was measured by a microscopic technique recently developed, and the value of the other parameter, the effective "diameter" of a dislocation pipe, was estimated. With these parameters equal to 7.3 X 106 dislocations per sq cm and 10-7 cm, respectively, it was found that the self-diffusivity along edge-type dislocation lines was a factor of about two times the average grain-boundary self-diffusivity. The dislocation-pipe diffusivity is about 7x10-7 sq cm per sec at 450°C. THE possibility of short-circuit diffusion through imperfections in crystals has been recognized for some time. It was only recently, however, that concerted effort has been expended to obtain values of diffusivity for grain-boundary short circuits. The work, for example, of Hoffman and Turnbull1 on silver, Fisher&apos;s&apos; analysis of the grain-boundary dif-

Jan 1, 1955