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By Selma Hermann

For the past century, the so-called aluminum bronzes have been assuming a role of ever-increasing importance in the metallurgical field. The last quarter of that century has marked many efforts to find alloying elements that would alleviate some of the undesirable qualities of the aluminum bronzes-for example, the heterogeneity of the alpha solid solution and the extreme hardness of the beta. Much of this work, however, has been rather haphazard-compositions and heat treatment are usually so much at variance as to make comparison impossible. It is the object of this research to study the effect of certain percentages of common alloying elements on aluminum bronzes of analogous compositions and heat treatments. It is hoped that the material gathered here will at least serve as a basis for further systematic study.

Jan 1, 1930

By C. BARING HORIVOOD

(Chattanooga Meeting, October, 1908.) THIS paper is intended, in the light of recent investigations, to call attention to some of the essential features of good practice in sampling and mine-valuation. Mine-sampling may be divided broadly into two classes : (1) development-sampling and computation of ore-reserves; and, (2) stope-sampling and estimation of stope-values. DEVELOPMENT-SAMPLING. It is with development-sampling only that this paper deals-the term development being taken to mean the opening-up of a mine by cross-cuts, drives, connections, etc., preparatory to actual extraction of the ore for milling. In development-sampling the work of the sampler and the consequent estimation of the probable value of the mine or portions of the mine, as the case may be, cannot be verified until a later period, when the ground thus sampled has been exploited; hence, extreme care and accuracy are essential. Moreover, the development-sampler has to foresee and allow for contingencies which may arise when stoping is started. It follows that the sampling, in development-work, demands familiarity with the conditions likely to be encountered later in actual mining. In stopes, on the other hand, sampling is clone under actual working-conditions, and the sampler's valuations are immediately checked by the gold-output, or in some cases by a second sampling of the broken ore. A stope-sampler therefore need not be an experienced miner, though he should possess sound technical training and work accurately. Should he be somewhat inexperienced, be soon has his errors discovered for him, and learns to avoid them. Taking these facts into account, it is obvious that a mining company employing two or more samplers should

Jan 1, 1909

By George S. Rice

The extremely valuable papers and discussions on coal-cluat explosions by Bache,' Eavenson, Shurick, Mannakee,* and Raymond are of unusual interest to me, since it has been my duty to carry on investigations on this and allied subjects for the Technologic Branch of the United States Geological Survey during the past year, and to make a series of coal-dust tests in the explosion-gallery at Pittsburg during a short periocl when its use could be diverted from the pressing work of testing explosives. In several of the discussions mentioned, it has been assumecl that the Technologic Branch of the Survey mas not paying attention to certain phases of the general problem, or to the specific question of temperature and humidity of mine-air in its effect upon the explosibility of coal-dust. This assumption is not correct, since from the time of the establishment of the " Mine-Accident Division," with its subdivisions of explosives, rescue-work, and mine-investigation, the mining engineer.; of the last-named section have made the subject one of chief consideration. Their general duties, including mine-sampling for governmeilt purposes, emergency-calls for rescue-work, and esplosion investigations, have prevented undivided attention, but all of them have been supplied with instruments for such determinations, and special report-forms have been in use and a great many observations made in mining-districts to which the mining engineers have gone. With the idea of obtaining a uniform system of reporting these data, and in the hope that the work may be of service to mining engineers in the field, the collecting blanks in use are shown in Forrns 1 and 2.

Jan 1, 1911

By R. E. Cech, J. H. Hollomon

KURDJUMOV and Maksimova reported experiments with manganese steels and high carbon steels' and with an Fe-Ni-Mn alloy' in which mar-tensite was formed isothermally over a range of temperatures. They found in some cases that mar-tensite formation could be suppressed by rapid quenching to liquid nitrogen temperature. From their microstructural observations of martensite formed isothermally, they concluded that the rate controlling step is nucleation rather than growth. Kulin and Cohen,3 in an attempt to reproduce these experiments, found that with a steel having the same composition as that reported by Kurd-jumov and Maksimova, the transformation to martensite was essentially complete above the temperature range of Kurdjumov and Maksimova's isotherms. The possible reasons for this disagreement were not considered. Recent papers by Das Gupta and Lement4 and Kulin and Speich5 report the formation of isothermal martensite in a high chromium steel and in an Fe-Cr-Ni alloy, but neither paper can be considered a verification of the original Kurdjumov and Maksimova results. Further, in neither case were the authors able to suppress the formation of martensite entirely. Because of the important bearing the Kurdjumov and Maksimova results have to an understanding of the mechanism of martensite reactions it was felt that an experimental investigation directly concerned with checking the validity of their results was in order. This paper describes the results obtained on the isothermal transformation over the temperature range from —79" to —196°C of an alloy of iron, nickel, and manganese. Experimental Apparatus A 15 lb heat of an alloy containing 73.3 pct Fe, 23.0 pct Ni, and 3.7 pct Mn was melted by induction and cast under argon. The ingot was forged to 1-in. bar and a portion rolled to 1/16x1 1/2-in. strip. This strip was pack-homogenized 300 hr at 1100" in a helium-filled sealed iron tube. The composition after homogenization was 73.2 pct Fe, 22.94 pct Ni, 3.73 pct Mn, 0.05 pct C, and 0.015 pct N. The strips were cut to 1/2-in. width for dilatometer and metal-lographic specimens. Only the center portion of the 11/2-in. strip was used in the present investigation. The dilatometer employed was similar in design to one described by Flinn, Cook, and Fellows." A concentric fused auartz rod and tube assembly with hooks for holding the specimen was mounted so as to transmit the specimen dilation to a 1/10,000 in., 1/10 in. travel dial gage. The dilatometer proper was mounted by means of extension arms to a counterweighted sliding member on a vertical standard. This method of mounting permitted rapid transfer of the dilatometer from the austenitizing furnace to the quenching bath and low temperature chamber. A small electrical vibrator on the dilatometer kept frictional effects of the quartz members at a minimum. The austenitizing unit was a vertical, molybdenum-wound, hydrogen atmosphere furnace maintained at a constant temperature ±3°C by means of constant power input. A 12-in. stainless steel jacketed copper liner having 1/2-in wall thickness acted to equalize the temperature in the hot zone of the furnace. This liner, closed at the bottom end and open at the top to permit entrance of the dilatometer and specimen, was kept filled with dry nitrogen gas. A chromel-alumel thermocouple was placed inside the tube to determine the temperature. The 4-in. dilatometer specimens in the chamber varied less than 1/2° across the specimen length except for a 1 1/20 drop at the end nearest the open end of the furnace. The low temperature isothermal holding bath was a double Dewar arrangement similar to one described by Turnbull7. The outer bath was filled with a refrigerant at a temperature lower than the desired holding temperature. The inner bath was filled with Freon "11" or "12" or a mixture of both, depending upon the holding temperature. This inner bath which tended to be cooled by the outer bath was kept at a constant temperature by introducing a small amount of heat with a manually controlled electric heater. Stirring was accomplished by bubbling dry air through the bath. A Leeds and North-rup type K potentiometer was used to measure the inner bath temperature as indicated by a five element copper-constantan thermopile. The bath temperature was maintained within ±0.2°C of the desired temperature by occasionally adjusting the heater current so as to keep the Leeds and Northrup galvanometer at zero deflection with a constant setting of the potentiometer. Isothermal tests were usually continued for 300 to 400 min and another reading made at approximately 1000 min if the bath, unattended overnight, had not deviated in temperature more than 5°C. Transformation curves are drawn dashed (Fig. 1) through the time region where temperature was not controlled precisely. Experimental Procedure Dilatometer specimens of 1/2x1/16-in. strip were cut to 41/2-in. length and holes were drilled for the quartz hooks with proper spacing to give a 4-in. measured length. A thermocouple consisting of 0.012-in. diameter chrome1 and alumel wires was spot welded to the specimen and threaded between the dilatometer rods to binding posts near the dial

Jan 1, 1954

By Richard L. Stowe, Donnie L. Ainsworth

The static, rapid, and shock-loading response of rock is of interest to many in the field of rock mechanics. For example, the effects of loading rates on strength and stress-strain characteristics of rock are quite important in defining input loading for use in the design of underground protective installations. Response data on rock subjected to shock loading over a wide range of pressures are also vital in the field of underground nuclear testing. In addition, material response is needed in design of shock-absorbing liners. A number of investigators have reported that ultimate compressive strength of rock increases with increased rates of loading.1-6 Wuerker3 stated that modulus of elasticity, obtained from sonic determinations, increased in some rocks by a factor as high as 4.5. It appears that only a limited amount of work has been done on comparing static modulus and modulus obtained from impact testing.6 Phillips4 showed that with increased rates of loading the deformation decreased. More recent work by Atchley and Furr5 showed that when moderate-strength concrete is stressed with increased loading rates, the axial strain at failure increased. One would expect this to occur if we think in terms of purely elastic materials. For example, if E is constant and maximum strength [amax], increases, the axial strain would also increase. However, in rock testing thus far we have worked almost exclusively with nonhomogeneous and inelastic rocks. From this fact alone the increase in axial strain with increased load application is noteworthy.

Jan 1, 1972

By B. L. Averbach, Morris Cohen, S. A. Kulin

The martensitic transformation can be initiated by elastic stresses at temperatures above M. in a steel containing 20 pct Ni and 0.5 pct C. Shear strains and normal tensile strains acting on a potential habit plane promote the transformation but compressive strains oppose it. THE strain sensitivity of the martensitic transformation has long been recognized. It is now well known that martensite formation can be induced by plastic deformation at temperatures below, and not too far above, the M, point.' However, the exact role of elastic vs. plastic strains and of the concomitant stresses has proved to be a very elusive matter. Many investigators2-' of the steel hardening reaction have made use of the stresses resulting from volume changes to account for such observations as the existence of retained austenite and its variation across the section of a bar, the change in the amount of retained austenite with quenching rate, and the self-stopping nature of the transformation when the cooling is stopped. The effect of applied strain as an independent variable has been studied in 70 pct Fe-30 pct Ni alloys by Scheil7 and McReynolds.8 Some data are also available for lithium and Li-Mg alloys,9 Cu-Zn,10 Cu-Sn and Cu-A1 alloys," and aus-tenitic stainless steels."-" It has been found that martensite formation can be induced isothermally, even at temperatures above M,, by mechanical deformation of the parent phase. The competitive nature of plastic yielding by slip and by the martensitic reaction has been depicted by Scheil.7 Basically, Scheil postulated that: 1—at temperatures not too far above M, (where austenite is less stable than martensite), a critical resolved shear stress (within the elastic range) is required to promote the transformation to martensite; 2—at and below M,, the austenite lattice becomes mechanically, as well as thermodynamically, unstable and shears over spontaneously into martensite without the application of external stress; and 3—at temperatures sufficiently far above M8, the critical resolved shear stress for martensite formation increases to a level above that required for slip, and plastic flow then supersedes the transformation when external stress is applied. Scheil's concepts were tested by McReyno1ds8 ho found no change in the elastic moduli (measured both statically and dynamically) of 71 pct Fe-29 pct Ni alloys in the vicinity of the M, temperature. Since the moduli did not approach zero or become negative on cooling to M,, it was concluded that the austenite lattice does not become mechanically unstable at M2, as postulated by Scheil. McReynolds also reported that M, is not raised by elastic stresses. Accordingly, plastic deformation was considered to stimulate the transformation by generating martensite nuclei in the distorted regions of the parent phase. The above issues with regard to the role of applied stress require clarification, if a basic understanding of the martensitic transformation is to be obtained. For example, in the reaction-path theory,'"," it is postulated that strain embryos exist in the austenite that provide part of the activation energy for the nucleation process. These embryos are visualized to be regions of localized strain in which the atoms in the austenite lattice are displaced part way along the path to their ultimate positions in the martensite. Consequently, it would

Jan 1, 1953

By William L. Fink and

Although age-hardening in an aluminum-base alloy containing magnesium was observed by Alfred Wilml as early as 1911, it was not until 1919 that a theory of the mechanism of age-hardening was proposed. Merica, Waltenberg and Scott12 studying the age-hardening characteristics of Duralumin at the National Bureau of Standards, came to the conclusion that age-hardening of this alloy is caused by the precipitation of CuAl2, which is less soluble at the aging temperature than at the solution heat-treating temperature. It was postulated that the initially precipitated particles were of fine colloidal dimensions but that they gradually grew in size. Maximum hardness was assumed to be attained when the particles reached a critical size. This theory was at first received skeptically, since it was a widely accepted view at that time that the solid solution state should represent the hardest state of an alloy. Shortly thereafter, Jeffries and Archer proposed their well-known slip interference theory.'13 They suggested that age-hardening was accomplished by the precipitated particles acting as keys interfering with slip along crystallographic glide planes. They showed that this mechanical theory led direct,ly to the assumption that maximum hardening would be produced by particles of a critical size, which they suggested would have a diameter approximately 10 times that of the aluminum atom. This contribution amplified and explained so rationally the age-hardening theory of Merica, Waltenberg and Scott that wide acceptance resulted. Objections to Simple Precipitation Theory From time to time certain experimental facts have been cited against the simple precipitation theory. The first of these was the discovery by Fraenkel4 that, although the electrical resistivity of Duralumin decreased during aging at higher temperatures (from 100' to 250' C.), it increased during aging at ordinary temperatures. This was not in accordance

Jan 1, 1940

By William L. Fink and

Although age-hardening in an aluminum-base alloy containing magnesium was observed by Alfred Wilml as early as 1911, it was not until 1919 that a theory of the mechanism of age-hardening was proposed. Merica, Waltenberg and Scott12 studying the age-hardening characteristics of Duralumin at the National Bureau of Standards, came to the conclusion that age-hardening of this alloy is caused by the precipitation of CuAl2, which is less soluble at the aging temperature than at the solution heat-treating temperature. It was postulated that the initially precipitated particles were of fine colloidal dimensions but that they gradually grew in size. Maximum hardness was assumed to be attained when the particles reached a critical size. This theory was at first received skeptically, since it was a widely accepted view at that time that the solid solution state should represent the hardest state of an alloy. Shortly thereafter, Jeffries and Archer proposed their well-known slip interference theory.'13 They suggested that age-hardening was accomplished by the precipitated particles acting as keys interfering with slip along crystallographic glide planes. They showed that this mechanical theory led direct,ly to the assumption that maximum hardening would be produced by particles of a critical size, which they suggested would have a diameter approximately 10 times that of the aluminum atom. This contribution amplified and explained so rationally the age-hardening theory of Merica, Waltenberg and Scott that wide acceptance resulted. Objections to Simple Precipitation Theory From time to time certain experimental facts have been cited against the simple precipitation theory. The first of these was the discovery by Fraenkel4 that, although the electrical resistivity of Duralumin decreased during aging at higher temperatures (from 100' to 250' C.), it increased during aging at ordinary temperatures. This was not in accordance

Jan 1, 1940

By L. W. Reynolds, R. F. Campbell, K. G. Carroll, S. H. Ballard

Based on metallographic and X-ray data probable equilibrium conditions from 1340" to 2400°F are presented for the composition range investigated. These are correlated with investigations of Takei and Kuo. Provision is made for existence of Mo, C and (Fe, Mo)2,C, carbides heretofore omitted in diagrams in this composition range. These carbides and (Fe, Mo),C, in equilibrium with ferrite and austenite, restrict the austenite and austenite + ferrite fields to lower carbon contents and higher temperatures than in the Fe-C system. OK groups of steel of different molybdenum content, but with approximately the same carbon variation in each group, Table I, were studied. The alloys were melted in an acid-lined 30-lb high-frequency induction furnace, and the ingots forged to l-in. rounds, which were then ground to 0.75-in. diam. To minimize chemical segregations a short length of each bar was given a high-temperature heat treatment as listed in Table I. Specimens 0.062 to 0.125-in. thick were thereafter cut from the homogenized bars for heat treatment to determine phase changes in the alloy. Heat treatments were performed in either a lead bath or muffle-type electric furnace, depending upon the temperature required. Maximum temperature variation was * 2" F in the lead bath and * 5"F up to 1900°F and +8"F above 1900°F in the electric furnace. Adequate protection of the specimen from the oxidizing atmosphere in the electric furnace was obtained by sealing the specimen in either a pyrex or a silica capsule which was evacuated or not as circumstances required. For austenitizing followed by quenching to an isothermal temperature, the specimen was placed in a closed-end refractory tube the mouth of which, extending beyond the furnace, was connected to a vacuum pump. The specimen could be withdrawn from this tube and quenched as desired. IDENTIFICATION OF PHASES The principal method of identification of phases was metallographic examination conducted on surfaces well removed from exposure to lead or atmosphere effects during heat treatment. Hardness measurements (Dph/lOkg) were also made on these surfaces. The etching reagents used and their effects are listed in Table II. Martensite and ferrite were readily recognized by their customary appearance after etching with saturated picral or 1 pct nital. The presence of liquid at high temperature was indicated by curved interfaces characteristic of boundaries between solid and melt, by intergranular pools of dendritic areas in a carbide matrix, and by a black network coincident with austenite grain boundaries. Examples of these structures are shown in Figs. l(a) and (0). The types of carbide which were attacked by the electrolytic chromic acid etch and the alkaline KMnO, etch, Table 11, were identified by X-ray analyses of extracted carbides of five alloys. Carbide residue was obtained from a metallographic specimen by etching the specimen for 24 hr in hydrochloric-picric acid in alcohol (grain-size reagent), after which the carbides, standing in relief, were scraped from the surface of the specimen and mounted on a glass slide. Diffraction patterns were obtained with a Philips spectrometer using iron-filtered cobalt radiation. Figs. 2, 3, 4, and 5 illustrate typical etching behaviors with chromic acid and alkaline KMnO,. Chromic acid attacked Mo,C as shown in Figs. 2(a) and (b)-but not (Fe, Mo),C, Fig. 3(b). Alkaline KMnO, attacked (Fe, Mo),C, Fig. 3(c), and also attacked M0,C. In accordance with these behaviors, a typical differentiation of coexisting Mo,C and (Fe, Mo),~ is shown in Figs. 4(a) and (b). The carbide (Fe, Mo),,C, could not be obtained free of other coexisting carbides for etching experiments. The behavior. however. of coexisting carbides identified by x-ray diffraction as Mo,C ind (Fe, Mo),,C, indicate that (Fe, Mo)23C6 is not etched by chromic

Jan 1, 1961

By A. A. Bauer, H. A. Saller, F. A. Rough, J. R. Doig

among uranium alloy systems are the U-Zr, U-Ti,and U-Mo binary systems, in that an intermediate phase, of limited solubility range, exists in each system. These intermediate phase alloys offer a promising field of investigation, both as binary alloys and as ternary alloy combinations. While a limited amount of work has been done on the binary alloys, none has been done on ternary combinations of these alloys. As a possible prelude to more intensive investigation of these systems, a study of the phase relationships between the 6 phases of the U-Zr and U-Ti systems was undertaken. The solid-state portions of the binary constitutional diagrams for the U-Ti,1 Zr-Ti,2 and U-Zr3 systems are shown in Fig. 1. The solid to liquid transformations are not shown, as these transformations were not studied in the ternary system. The phase designations employed are based upon those shown in the uranium constitutional diagrams. The crystal structures of the various phases are given in Table I. Experimental Materials and Methods The ternary section presented is based upon the results of thermal, metallographic, and X-ray analyses of a series of alloys ranging in composition between the phases of the U-Zr and U-Ti systems (U-74 atomic pct Zr to U-35 atomic pct Ti). After preparation of the alloys, thermal data were obtained. On the basis of these data, specimens were heat treated and examined metallographically. Specimens were then selected from those heat treated for X-ray determination of the phases present in what were considered critical phase regions. Alloy and Specimen Preparation—Alloys for this investigation were prepared from selected as-reduced uranium and crystal-bar zirconium and titanium. For nominal compositions, see Table 11. Charges were arc melted from five to seven times under a helium atmosphere, finally being cast into wire-bar form for fabrication. A zirconium getter button was used to clean the furnace atmosphere prior to melting. The alloys were homogenized 20 hr at 900 °C and furnace cooled. A portion of each bar was then removed for thermal analysis; the remainder was saved for fabrication. Alloys 8 through 14 were placed in steel tubes along with zirconium getter chips and, after flushing with argon, the tubes were welded shut and rolled at 1500 °F to a 50 pct reduction in diam. The remaining alloys were annealed 24 hr at ll00°F to insure transformation of the ? phase to the 8-U2Ti phase and were then hot rolled from a salt bath. Attempts to roll alloys 6 and 7 at 1100 °F were unsuccessful. Following this, attempts were made to roll pieces of alloys 1 through 5, which had cracked rather severely during casting, at 1200 °F, but these efforts were also unsuccessful. These alloys were, therefore, further treated in their cast and homogenized condition. All alloys were pickled to remove oxide formed during rolling. Thermal data were obtained on these alloys by conventional differential thermal analysis at a heating rate of 30C per min, and by a high-speed heating method at rates of approximately 10°C per sec. For the latter method of analysis, a 36-gage Chromel-Alumel thermocouple is spot welded between two slivers of alloy, the sandwich being en-

Jan 1, 1958

By R. W. Hertzberg, F. D. Lemkey, J. A. Ford

The technique of unidirectional solidification has been applied to the A1-AI3Ni and A1-CuAl2 ezltectic alloy systems; the controlled microstructure of A1-A3Ni consists of parallel A13Ni whiskers emhedded in an aluminum matrix while the Al-CuAl2, system solidifies as parallel alternate lamellae of aluminm and CuAl2. Mechanical tests (tension and flexure) have shown that the particle-matrix interface bond formed during unidirectional solidification allowes for efficient load transfer from the matrix to the reinforcing phase. Strength and fracture characterstics of the alloys were studied as a function of phase orientation; the tensile strength of controlled Al-Al3Ni samples was 43,000 psi (three times greater than that of as-cast specime~s). The strength of the Al-CIA2, cutectic was studied as a function of lamellae-flexure stress-axis orientation using a microbend apparatus. A twofold increase in outer fiber stress was observed as the lamellae were rotated from a position of no reinforcing (perpendicular to the stress axis) to one of maximum reinforcing (parallel to the stress axis), The brittle reinforcing constituent in each alloy (A13Ni and CuAl2) was observed to he the nucleating phase for fracture. It is to be concluded from this investigation that unidirectionally solidified eutectic alloys have a potential use as reinforcing coniposite) materials. In a recent review article, Chadwick1 has described the work of many investigators in the area of controlled solidification of binary eutectic alloys. Thus far, more than twenty eutectic alloys have been unidirectionally solidified to produce structures consisting of two phases in the form of alternating lamellae or as whiskers embedded within a continuous matrix; in either case, the phases are aligned parallel to the growth direction. In particular, Kraft and Albright2 have shown that the Al-CuA1, eutectic alloy exhibits a lamellar microstruc- ture as a result of unidirectional solidification while Lemkey, Hertzberg, and Ford3 have studied the A13Ni whisker-aluminum matrix morphology in the A1-A13Ni system. Whether these alloys could exhibit reinforcing behavior would have to depend upon the strength and volume fraction of the reinforcing phase (AI3Ni and CuA12) and the nature of the bond between the reinforcing phase and the matrix. Hertzberg and Kraft4 have shown that the interface bond between chromium whiskers and a copper matrix in unidirectionally solidified Cu-Cr is particularly good while Lemkey and Kraft5 have reported chromium-whisker tensile strengths in excess of 1,000,000 psi. However, with the chromium volume fraction being less than 2 pct, no reinforcing behavior was displayed by the Cu-Cr alloy.6 In view of the encouraging results concerning the good interfacial bond and high strength of the reinforcing phase in the Cu-Cr system and preliminary mechanical test data reported by Lemkey et al.3 for the A1-A13Ni system, the A1-CuA1, and A1-A13Ni eutectic alloys containing 50 and 10 vol pct, respectively, of the reinforcing phase give promise of possessing mechanical reinforcing behavior. Based on this premise, the objective of the investigation was to study the strength and fracture behavior exhibited by these anisotropic microstructures. EXPERIMENTAL PROCEDURE Controlled Solidification Procedure. An extensive description of the technique of unidirectional solidification has been presented elsewhere.' However, the pertinent details for each eutectic system studied are discussed below. Al-A13Ni System. Several master heats of this alloy were prepared by vacuum-induction melting in an A12O3crucible and casting in SiO2 mold or by induction melting in alumina crucibles under argon. The initial heats were prepared from 99.99 pct A1 and 99.99+ pct Ni, while later heats used 99.999+ pct A1 and 99.99+ pct Ni. Specimens of these master heats were unidirectionally solidified using induction and resistance heating sources and argon atmospheres at growth rates from 1.1 to 28.7 cm per hr. The thermal gradients in the liquid were controlled in the range of 26" to 36°C per cm for these runs. Al-CuA12 System. The master heat and controlled

Jan 1, 1965

The regular activities of the Institute are financed mainly by income derived from members' dues, from advertising in MINING AND METALLURGY, and from the sale of publications to the public In addition, collateral activities and special projects are financed each year by means of income from endowed funds controlled and administered by the Institute In the aggregate, the amount of these endowments totals more than $830,000 An active effort is being made to increase the endowment funds of the Institute through a standing committee, of which Harvey S Mudd is Chairman, and

Jan 1, 1952

By Charles Hardy, George D. Cremer

POWDER metallurgy has been established for some time as a novel method for manufacturing a great variety of articles generally specialties that could not be made conveniently by any other method. In this category are such parts as porous hearings cutting tools of sintered carbides, and electrical contacts in which the metallic constituents do not tend to alloy in the proportions employed. The industry, however, is now at a stage where it has begun to manufacture articles that can and have been made by other methods. Thus, powder metallurgy is now preferred technique for the production of numerous machine parts characterized by high strength, accuracy of configuration, and substantial lack of porosity. Manufacture of such machine parts from metal powders has no doubt been stimulated by the war, although the trend in this direction has been apparent for several years. In any case, quantity production of high-density precision pieces by ponder metallurgy is now in full suing in the United States, and such production can and should continue to expand during and after the war.

Jan 1, 1942

By D. A. Dahlstrom

Although the liquid-solid cyclone is a relatively recent innovation in the field of coal preparation, various authors have already indicated three distinct applications to operations encountered in the modern tipple. Driessen's articles and work at the Dutch State Mines exhibited its possibilities in the bene-ficiation of fine coal with apparently high efficiencies down to the 200-mesh fraction. l,2,3,4 Other publications point out its use as a recoverer, thickener, and preliminary dewatering agent of fine coal which in too many cases is being wasted due largely to former high operating cost limitations.M Finally, certain authors have emphasized the practicability of applying the cyclone to the elimination of coarse and fine solids from water slurries in the tipple in order to produce a water suitable for recycle purposes containing only minute particles far below the 200-mesh size.6 This would result in a closed water system for many cases, with its accompanying economic and operating advantages to locations where water quantity and quality problems are severe; better operation of washing equipment using a consistent water devoid of injurious fine particles; lowered maintenance costs due to the lesser abrasion resulting when the fine abraiding particles are removed; and the greater ease of disposing of refuse streams of reduced volume containing greater percentages of solids. While the three applications have been indicated to be feasible, only a meager amount of data is available on capacities and solid elimination eficiency as a function of design variables and cyclone friction loss. This information is essential for the optimum use and construction of the solid-liquid cyclone, and accordingly, work on the problem was initiated at Northwestern University. For the unfamiliar reader, briefly the cyclone consists of a cylindrical section mounted above a truncated cone. The feed nozzle enters the cylindrical ring tangentially with the underflow nozzle permitting discharge of the concentrated solids located at the apex of the cone. The overflow nozzle through which the clarified water exits is centered in the cylindrical section at the top of the cyclone. Half sections of experimental cyclones will be found in Fig la and 6. The feed slurry enters with a tangential velocity, thus creating a spiral pattern of high centrifugal force. The solid particles of sufficient size and gravity are ejected outward to the walls and spirally discharge to the underflow. Most of the water with uneliminated fine solids moves radially inward along the path of the outer

Jan 1, 1950

By J. B. Darby

In an investigation of various equiatomic alloys of transition elements, Dwightl did not find evidence for the CsC1-type ordered structure in VMn. More recently waterstrat2 has reported the occurrence of this structure in the alloy his sample had been homogenized at 1150°C and subsequently reannealed at 725°C after crushing to powder. The purpose of this note is to confirm the existence of the CsC1-type ordered structure in VMn and to report a lattice parameter. An equiatomic alloy was prepared by arc-melting under an argon atmosphere vanadium and electrolytic manganese having a purity of at least 99.9 pet. The composition was obtained by charging an excess of manganese and monitoring the weight loss until the stoichiometric composition was achieved; the assumption that the loss during melting was due to the vaporization of manganese appears to be a valid one.' The calculated composition of the sample prepared was 49.85 at. pct V and 50.15 at. pct Mn. Specimens taken from the fractured as-cast buttons were wrapped in molybdenum foil and encapsulated in quartz capsules under vacuum; there was no evidence of a reaction between the alloy and foil. The X-ray patterns were obtained using filtered copper or chromium radiation with a Straumanis-type Debye-Scherrer camera having a diameter of 114.6 mm. The CrKa radiation is best suited for detecting the superlattice lines since it overcomes the slight difference in scattering powers of the components by making use of the anomalies in atomic scattering factor which occur when the frequency of the radiation is close to the characteristic absorption frequency of one of the elements.394 A specimen in the as-cast condition, as well as one which had been annealed for 1 week at 900°C and water-quenched, gave a diffraction pattern for the disordered bcc solid solution. However, when samples were homogenized at 900°C for 3 weeks followed by a subsequent anneal at 700°C for 6 to 7 weeks and then water-quenched or furnace-cooled, the (loo), (1111, and (210) Ka superlattice lines were observed in films taken with chromium radiation. An ao = 2.904 * 0.001 was determined by a Nelson-Riley extrapolation from a film taken with CuKa radiation. The author wishes to thank Mr. J. W. Downey and Mr. L. J. Norton for carrying out the experimental portion of this investigation. This work was performed at the Argonne National Laboratory under the auspices of the U.S. Atomic Energy Commission.

Jan 1, 1963

By W. W. Anderson, S. Razi

Electroluminescetrce of single crystals of terbium-(loped ZnS prepared by vapor-transport technique shows the sharp line specirum characteristic of the 4f— 4ft,ansitiotzs of the trivalent Tb3 rotz. V-I tt~easuverr~ents give evidence of space-ellarge-lirrlited curvent but the thrveshold for trap-filled law behavior is not iu agreement with Lampert&apos;s theory for. Single injection. Variations of &apos;brightness with applied voltage, the observation of double peaks its brightness because joms, and the spatial distribution oi electroLur?zir~escerrce indicate that the accelet~atiotz-collision mechanism involving the bst lattice and/ov shallow traps is most likely to be responsible fov excitation of&apos; electrolnminescence. Efficiency rtreusuver)~etits show the quantwn efficiency to be about 10 pct and powev efficiency about 0.05 pct. Effect of anr~eallng the crystal in sulfur vapor is to enluztzce llle rare-earth emission. It rs pvoposed tlzat sulfitv anr~ealing crreates acceptorr-lvpe defects with which the donor-type vare-eavtll ion can associate more readily vesulting in enhanced rare-earth emission. A&apos;o such e~zlznr~cerr~etrt is obserued when the crystal is atztrealetl in zinc vapor. Photolianinescence of ZnS doped nith a variety of rare earths also shows tile slurvp l~rze rwve-eavtlz erriission which in sorrretirr~es accompanied by broad band, stvuctureless lattice emission. Photo-atrd electrolutr~itzesce?~ce of ZIIS:Tb slw~rj do!rlit~unt rare-earth emission in the ~ticirzity of 54(3OA corre-sporrdit~g to the transition D* — Fj. Hoz~!el)er, the detailed line structuve of the luo spectvtr is cliffevet~t, irzdicutit~g that different sites are active in the two processes. Decay of rave-eartlr fluorescence in ZnS doped with any of sei!evul vuve eurtlzs car1 be described by a single exporleritial e.scepl joy ZrlS:lIo. Tl~is exceptiotr can be explaitred it~ tevrr~s of tlre closely spaced er~evgy 1e1:els Jov the HO~&apos; iorr. Decay lime measurertzekzts jov ZnS:Tb, using pulsed elect,-ical ar~d pulsed opticcll excitutiorzs, (11-e itz goor1 agrcetrier~t. LUMINESCENCE of rare-earth-doped materials has been a subject of interest for the past 20 years. Within the past few years there has been a considerable increase in rare-earth research motivated in search of new and more efficient laser materials and also due to the use of certain-rare-earth compounds in the preparation of color television screens. The purpose of this study has been to seek an understanding of some of the basic processes involved in exciting the rare-earth luminescence which is associated with transitions within the 4f shell of the trivalent rare-earth ion. Single crystals of ZnS doped with a variety of rare-earth ions have been prepared by vapor-transport technique described elsewhere.&apos; Photoluminescence was excited by a high-pressure short-arc mercury lamp together with suitable glass and chemical filters. For electroluminescence, sinusoidal and pulse excitations were used. 1) ELECTRICAL CHARACTERISTICS 1.1) V-I Measurements. Electroluminescence experiments were performed on crystals of terbium-doped ZnS. The samples were cleaned and etched and indium or In-Ga alloy contacts were alloyed on by heating in H2 atmosphere to 600°C for times ranging up to 10 min. Static voltage-current measurements were made on several samples. Fig. 1 shows the results for a typical sample. For voltage V < 20 v, the V-I relationship is linear giving a resistivity of 2.5 x 109 ohm-cm for this particular sample at room temperature. In the range of 20 to 250 v, I varies as V "3 and at still higher voltages (when electroluminescence is visible to the scotopic eye) current varies as Vs up to 600 v, all at room temperature. At 77"K, for V > 200 v, / I vge5 up to 1000 v. The V-I characteristics at room temperature follow reasonably well the behavior predicted by Lampert&apos; for one carrier space-charge-limited current in an insulator with traps although, as shown later, the expression derived by Lampert2 for the threshold for trap-filled law behavior Vtfl yields an unrealistically low value for trap density if we use the experimental value of 300 v for VtfL. Assuming the case for shallow trapping, the transition from Ohm&apos;s law behavior to space-charge-limited behavior occurs at voltage Vtr given by where no = thermally generated free carrier density, L = length of the sample, e = static dielectric constant, 6 = ratio of free to trapped electron densities, e = electron charge. For the ZnS:Tb crystal, L = 0.5 mm, E = 8.3 €0, Vtr - 20 v, and no = 5 x 10&apos; per cu cm, calculated from the ohmic behavior assuming electron mobility of 100 sq cm per v-sec. This results in 9 = 0= As more and more electrons are injected the Fermi level moves up in the forbidden gap toward the conduction band. If we assume a single-energy level for traps (which is not strictly correct, as we will show later), the current voltage characteristic is profoundly affected when the Fermi level crosses the trap level. The traps are now filled and injected carriers can no longer be immobilized in traps. Hence, current rises sharply with voltage. The transition from space-charge-limited behavior to the trap-filled behavior occurs at voltage VTFL given by

Jan 1, 1968

By B. C. Allen

Surface free energy and surface self-diffusion of solid molybdenum were studied in the temperature range 1600" to 2400°C using pressure-sintered bi-crystals. Comparison of groove angles formed in various surfaces perpendicular to the grain boundary indicate a maximum of 1 pct variation in surface free energy with crystallographic mientation. This anisotropy tends to decrease with increasing temperature. The surface diffusion of the bicrystals is equivalent to that of sheet with a mild (100) Preferred orientation. Anomalously low values found for bi-crystals with surface orientations of (OOl), (012), and (011) are rationalized in terms of anisotropy in surface free energy. THE effect of crystallographic orientation on surface free energy1,&apos; and surface self-diffusion3,4 has been primarily studied in fcc metals. The object of this work was to study the effect of orientation on surface diffusion and surface free energy of bcc molybdenum using pressure-sintered bicrystals. EXPERIMENTAL WORK Materials and Crystal Preparation. Arc-melted molybdenum rod was obtained commercially and electron beam zone refined at 50 cm per hr at 10- 5 torr to form single crystals about 8 cm long and 0.65 cm diam. Three crystals were prepared with axial orientations about 1 deg from [001.], [011], and. [111]. To reduce the carbon content, the crystals were annealed 2 hr in 1.4 atm flowing wet hydrogen at 2050°C. Then the oxygen content was reduced by annealing for 2 hr in -30°C dewpoint hydrogen at 2020°C. The resulting impurity analysis is given in Table I. Bicrystal Preparation. The single crystal rods were cut into transverse slices with a thin silicon carbide abrasive wheel to produce specimens about 0.6 cm long. They were mounted in epoxy and surrounded by stainless steel washers. Cutting in half was done longitudinally at various angles to known crystallographic planes containing the cylinder axis according to Fig. 1. To reduce surface deformation resulting from the cutoff wheel and thus reduce parasitic grain boundary formation on subsequent annealing, about 0.003 cm was manually ground off each cut surface with 600 grit paper. Care was taken to keep the surface flat. After removal from the mounts, one half was generally ro-tated 180 deg with respect to the other to give a po- tential symmetrical tilt grain boundary between the two halves. In the other cases when low misorienta-tion angles were desired, the crystals were not rotated. On the basis of symmetry, sufficient bicrystals were prepared to cover the entire range of misorientations for symmetrical tilt boundaries. The misorientations, +, ranged from 0 to 45, 0 to 90, and 0 to 60 deg for [001], [011], and [111] bicrystals, respectively. One [Ill] twist bicrystal was prepared from 2 single crystal discs rotated 17 deg relative to each other. Each specimen consisted of two pieces which were placed in a cylindrical tantalum can. Sharp edges were rounded and the fit was made as snug as possible to reduce subsequent deformation during bonding. The assembly and crystals generally were vicuum outgassed at 900" or 1700°C and then electron beam welded in the can at l x 10-4 torr. After being leak checked, the samples were placed in an autoclave and hydrostatically gas-pressure bonded5 in four batches under helium at 10,000 to 18,000 psi at 1650°C for 3 hr. Satisfactory bonds were obtained in many cases, and most of the crystals bonded after two exposures. The results did not appear to be affected by the various pressures used, preannealing conditions, crystal orientation, or time-pressure-temperature route taken to the final bonding condition. After bonding, the tantalum cans were selectively removed in cold concentrated HF. Measurements indicated overall deformation was under 1 pct. The bicrystals were metallographically ground and polished flat and perpendicular to the axis. Examination showed the boundaries were straight and almost free of parasitic grains caused by extraneous local deformation. Annealing. In preparation for thermal grooving, the bicrystals were cleaned and annealed by outgassing at 10-5 torr at 1900°C and heating at 2300°C under 1 atm 99.996 Ar for 0.5 hr. The crystals were held in a closed 4-deck box made of molybdenum sheet, and were heated in a Ta-1OW resistance furnace. The ar-

Jan 1, 1970

By I. W. L. Finlay, H. R. Ogden, R. Jaffee, D. J. Maykuth

Aluminum has been found to be soluble in a titanium to about 26 pct, and to raise the temperature range of transformation from a to 8. Two intermediate phoses exist in the system, a new face-centered tetragonal phase, designated as 7, which occurs between 34 and 46 pct Al, and TiAI Metallographic and X-ray diffraction data were used to determine the diagram. ALUMINUM is one of the useful alloying additions for titanium-base alloys. It is an important constituent in two commercial alloys released in 1950, one containing 4 pct A1 and 4 pct Mn, and the other containing 3 pct A1 and 5 pct Cr as the metallic alloy content. It is probable that, in the future, other titanium alloys containing aluminum will be released. In the course of alloy development work, some of the features of the constitution of titanium-rich Ti-A1 alloys have been developed. These are described in the present paper. The only available information on the Ti-A1 system concerns the aluminum-rich end of the diagram. In summarizing the early work, Hansen&apos; indicated that there is little solubility of titanium in aluminum, and that a two-phase region exists up to the intermediate compound, TiA1a (37.2 wt pct Ti), which has a congruent melting point at 1355°C. Fink&apos; has reported that a peritectic reaction occurs at 665°C between the melt and TiA1³ to form the terminal solid solution of titanium in aluminum. The solubility of titanium in aluminum has been established at 0.24 pct at 510°C and is estimated to be between 1 and 1.5 pct at the peritectic temperature. No information has been published on the diagram above TiAl,, other than a statement- hat aluminum raises the a to ß transformation temperature of titanium. Materials and Preparation of Alloys Iodide-type titanium and high-purity aluminum were used in this investigation. Typical analyses of these materials are given in Table 1. Alloys were prepared by arc melting on a water-cooled copper hearth, using the techniques described previously for contamination-free melting.&apos; Weight losses during melting were very low, ranging from 0 to 0.02 g in a 10 to 15 g charge, indicating that little or no loss of aluminum occurred during melting. This permitted the use of nominal compositions for these alloys. Alloy ingots containing from 0 to 7.5 pct A1 were rolled in air at 850°C to 0.040 in. sheet. Alloys containing from 7.5 to 16 pct A1 could only be rolled about 25 pct reduction at temperatures up to and including 1200°C. Alloys containing above 16 pct A1

Jan 1, 1952

By F. W. Jessen, N. Mungan

The plastic flow characteristics of clay water suspensions were first recognized by Binghaml in 1916 and further studied by Ambrose and Loomis&apos; in 1931-1932. Many physical and chemical properties of clay suspensions are deeply influenced by particle size distribution as well as by the exchange cation. Separation of a number of fractions is particularly helpful in making a detailed study of a suspension. A limited amount of work has been conducted on the effects of particle size on specific properties of mud drilling fluids. A knowledge of particle size distribution is essential for a better understanding of fundamental properties of clay suspensions, particularly viscosity and gel strength. The purpose of this investigation was to study the effect of particle size on the viscosity and gel properties of clay suspensions and to determine any changes in mineral composition accompanying variation in size. EXPERIMENTAL PROCEDURE Numerous methods may be used to fractionate a poly-dispersed clay suspension into nearly mono-dispersed systems. Particles larger than 44 micron (325-mesh) may be obtained by wet screening. For sizes from 44 to 1 micron, the fractionation may be accomplished by sedimentation, elutriation plus microscopy,&apos; or turbidimetry.6 When dealing with systems which contain a major portion of particles below l-mi-cron size, slit ultramicroscopy,7 centrifuge methods of Svedberg,5 eams,9 and McBain,l0 and X-ray methods" may be employed. The Sharples super centrifuge has been proved suitable for the economical separation of suspended particles of submicron range.12,13,14 Three-hundred and forty-five gal of 1 per cent by weight Aquagel suspension was mixed in a high-speed 5-gallon mud mixer and the suspension was allowed to stand in 5-gal glass bottles for 30 days to assure full hydration. All particles larger than 44 microns were removed by wet screening. Five fractions in the range of 44 to 1 micron were obtained by sedimentation. The Sharples super centrifuge method described by Hauser and Reed.13 Norton and Speil,14 Ha user and Lynn,15 and Hauser and Schachman,10 as employed to yield five fractions of submicron size. A super centrifuge operation chart prepared by Fancher, Oliphant, and Houssierre12 was used to determine operating conditions for the desired particle ranges. The clay suspension, after removal of all material above 1-micron size, was centrifuged. During rotation of the bowl, the particles were deposited on a thin plastic liner which fit closely inside the bowl. After passing 5 gal of suspension through the centrifuge, the liner was removed and the particles on the liner were scraped off, remixed and recentrifuged. This was repeated five times. Thus, Fractions H, I, K, and L were obtained. The finest fraction, M, remained in an approximately 0.4 per cent suspension. Water was evaporated under reduced pressure from this suspension to yield the finest particles. Results of the fractionation are included in Table 1. In order to achieve proper orientation and humidity conditions for X-ray diffraction analysis, the following procedure was employed. The dilute suspensions were placed on glass slides and allowed to dry at room temperature. This was repeated several times. Previous to measurement each sample first was placed in a water desiccator for half an hour and then allowed to stand for 15 minutes at room temperature. thus yielding an equilibrium condition of about 52 per cent relative humidity. The relative humidity of the room in which the X-ray diffraction was carried out was maintained at approximately 50 per cent so that there was no change in the spacings during the X-ray examination. Samples

By C. G. Dunn, F. W. Daniels, M. J. Bolton

IN recent investigations1. a data on relative grain boundary energies in silicon iron have been obtained. The present investigation is a continuation of this work along similar lines for the purpose of obtaining additional information on boundary energies. One aim has been to extend the data for a (110) series, which was formerly studied,&apos; to small differences in orientation, since a recent theory3,4 based on a dislocation model of grain boundaries predicts a E. ? (A-ln ? )* form of energy curve for small differences in orientation. Another aim has been to provide data for a (100) series—all grains having (100) planes parallel with the surface of sheet speci-mens—and to compare the results with the predicted form of energy curve. Finally, information on the nature of approach to equilibrium conditions was to be obtained through observations on the movements of grain boundaries. Recently published papers2,3,5 have treated the mathematical problem of expressing equilibrium conditions for grain boundaries for the general case when boundary energy depends upon boundary orientation. The equilibrium equations, which relate equilibrium angles and grain boundary -free energies, contain additional terms that express the variation of energy with boundary rotation. It has been necessary in the present investigation, however, to neglect these additional terms since no data for their evaluation were available. Dropping the additional terms leads to the following approximate equations which were used: E12/sin ?3 = E18/sin?2 =E28/sin?1 or those of the form: E28 + E13 cos ?3 + E12 cos ?2 = 0 Experimental Procedure Except for sample F1, all specimens were made from two lots of silicon iron called C and L respectively. The compositions of these are listed in table I. In the preparation of 12 three-grain specimens for the present investigation, a controlled grain growth technique, which has been described previously,&apos; was used. After preparation, the specimens were notched at the boundaries to shorten them as a means of providing more rapid approach to equilibrium in the anneals. Initial grain boundary angles were determined from micrographs taken at X500. Seven anneals, totaling about two days at 1300°C and two to four days at 1400°C, were run in an atmosphere of pure dry argon as described in a previous paper.&apos; Boundary changes generally could be followed without metallographic surface preparation, because thermal etching occurred during the anneals. These observations indicated angle changes of as much as 35" in some instances, with the major changes occurring during the first anneal. Annealing was discontinued when the angles reached stationary values. Fig. 1 gives a schematic diagram for purposes of defining crystallographic and boundary directions for three grains with common plane parallel with the sheet. Differences in orientations (expressed by A) are the angles between [00l] directions. They can be calculated with the aid of the ?&apos;s. The grain boundary angles can be calculated from the w&apos;S. Results Results are given in table 11. Types of crystal-lographic planes refer to those parallel with the sheet. Orientations are given in terms of [00l] direc-tions as shown in fig. 1. Except for grain 3 of specimen H6, which was unintentionally tilted 7" out of a (100) plane, all grains generally were within 1" or 2" of the plane specified. The directions of boundaries were measured generally to within 1" on both upper and lower surfaces, and the average direc- tions were determined for use in calculating the grain boundary angles. Variations in grain boundary directions refer to deviations from average directions. Boundaries near the junction point for all specimens generally were very close to perpendicu-

Jan 1, 1951