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By R. Woods

INTRODUCTION Gaudin (1), considered that "the mechanism of mineral collection is the central problem of flotation theory". From their work on adsorption phenomena in flotation for more than fifty years (2), Gaudin and his co-workers have made major contributions to the elucidation of this mechanism, a fact acknowledged by their contemporaries in the field (e.g. (3)). Their investigations in a number of areas, e. g. demonstrating that submonolayer quantities of adsorbed collector species can induce flotation, or identifying the influence of substances which can activate or depress (1), have helped to provide the basis for and the development of a theory to account for the interaction between mineral and collector. It was established in the late fifties (4) that the presence of oxygen is a prerequisite for the flotation of metal sulfides with thiol collectors. This fact rendered inadequate the concept of simple ion adsorption being responsible for rendering these minerals hydrophobic. The requirement for oxygen has been explained (1) in terms of this species being a component in chemical reactions between the sulfide surface and the collector. Gaudin has been instrumental in identifying the products of chemical reactions at sulfide minerals (1, 5), a line of work which has been continued particularly by Finkelstein and co-workers in South Africa (6). In presenting a detailed mechanism of the mineral-collector- oxygen interaction, it has often been considered that the mineral surface [must] first be oxidized to a metal oxy-sulfur species which

Jan 1, 1976

By John D. Roach

DURING an investigation of the yttrium-hydrogen system aimed at producing solid yttrium hydride specimens containing various amounts of hydrogen, it was observed that yttrium containing approximately 2 wt pct H exhibited a tendency to crack and crumble to a powder on standing in air at room temperature. It was also observed that longer hydriding times, at a given temperature, increased this susceptibility to powdering without an increase in hydrogen content, also that moisture in the air is necessary for the observed effects to occur. Even if cracking did not occur, there was a continual formation of a light gray powder on the as-treated surface of the hydrided yttrium and this reaction continues until the entire hydrided piece has been reduced to a powder. Storing the hydrided specimens in the absence of air or removing the surface either by machining or grinding were effective means of preventing this disintegration of hydrided yttrium. In an attempt to determine the reason for this powdering phenomenon, the surfaces of a number of as-hydrided yttrium specimens were examined by X-ray techniques. In all cases the X-ray pattern obtained showed the major phase to be yttrium hydride (YH2) as would be expected since the specimens contained 2 to 2.2 wt pct H. In some cases a trace of yttrium oxide was observed. There was also a third phase present on the surface of these hydrided specimens which could not immediately be identified. This unknown phase was a face-centered cubic material, NaC1-type structure, with a lattice parameter of 4.855 and a calculated density of 5.914 g per cc. Very slight hand polishing of the surface of the hydrided yttrium specimens completely removed both the unknown phase and the traces of oxide so that only the yttrium hydride pattern remained. Based on the X-ray patterns the quantity of the unknown phase on the surface of the hydrided yttrium appeared to be directly related to the susceptibility of the material to powdering. Work at the Denver Research Institute on the yttrium-nitrogen binary system showed that yttrium nitride (YN) is a face-centered cubic material with a lattice parameter of 4.878A and a density of 5.890 g per cc. They also noted that this compound rapidly disintegrated to a powder on standing in air. The unknown pattern observed in the above specimens corresponds very closely to that of the nitride—fcc structure, 4.885 parameter, and density of 5.914 g per cc. The presence of this thin film of nitride on the surface of the hydrided specimens probably accounts for the observed powder formation and crumbling. The nitride reacts with the water vapor in the air (verified by private communication from Dr. C. Huffine, General Electric Co.) to yield yttrium oxide and ammonia. Ammonia is readily detected when hydrided yttrium specimens were allowed to stand in bottled moist air. The powder formed on the surface of the hydrided specimens was shown by X-ray analysis to be yttrium oxide. This reaction appears to occur primarily at the grain boundaries since discrete particles of yttrium hydride separate from the specimens during this powdering process. The reaction of yttrium nitride with water vapor is believed to be as follows: 2YN + 3H2O - Y203 + 2NH3. Despite the fact that this nitride is present only as an extremely thin surface film, if the above reaction is not prevented by removing this film from the surface either by machining or grinding, the reaction continues until the entire hydrided piece has been reduced to a powder. To account for this continuation of the reaction, it is believed that the following reactionalsooccurs: 2YH2 + 2NH3 -2YN + 5H2. The nitride produced by the latter reaction reacts in turn with water vapor. The reaction therefore becomes autocatalytic and continues until the hydride has been consumed and the entire piece reduced to oxide powder. The amount of nitride required to initiate this reaction is quite small and this nitrogen contamination can occur from a number of sources—hydrogen gas employed, minute leak in the hydriding apparatus or even from degassing of the reaction vessel itself. Longer hydriding times increase the possibility of nitrogen contamination and this is especially true when a dynamic gas system is employed in the hydriding process. The production of stable, solid hydrided yttrium is dependent on the complete absence of nitrogen contamination during processing. If such contamination does occur, powdering of the hydrided product can be prevented by removing the nitride from the surface e.g. grinding or by preventing access of air to the pieces e.g. sealing in wax or plastic. The author wishes to express his appreciation to General Electric Co. for sponsoring this research and for permission to publish the results of work under Subcontract AT-93.

Jan 1, 1962

By R. R. Harvey

Recent evidence has indicated that the films forming at crude oil-water interfaces are the result of naturally occurring su,rface-active components in the crude oil rather than oxidation products resulting from exposure of the oil to air. Since the presence of such films may strongly affect the displacement of oil from porous media, a knowledge of their composition and formatiorz is of fundamental importance. In this study, without the assumption of a specific mechanism, film formation is considered to involve the transfer of a weakly surface-active species in solution in the oil to the oil-water interface. The film gas law and the Gibbs adsorption equation may thus he applied to the system. The rigid films which form upon decrease of the oil-water interfacial area are considered a result of the very slow rate of re-solution of the film material into the oil. When the shrinkage of the interface is very rapid, the amount of film material leaving the interface is negligible, and the system may then be treated in the same manner as an insoluble film. The combination of these treatments leads to a new concept of the film factor in which it may he quantitatively related to the concentration of film-forming material at the interface. Suitable experiments can then yield the following information. 1. Original concentration of surface-active species in the oil. 2. The change of concentration with time. 3. Concentration of surface active species at the interface. 4. The change of this concentration with time. 5. Interfacial tension of the system at any time and its variation with time. 6. Equilibriunz concentrations for a given set of conditions. 7. Concentration requirements for solid film formation. 8. Molecular diameter of the adrorbed species. These calculations may be made solely from data taken from measurements on the original oil. While using the pendant drop method of interfacial tension measurement of crude oil-water systems, Bartell and Niederhauser' remarked on the thin, transparent, pliable, solid films or membranes which were often formed at the interface between crude oil and water. As a result of their observations of these films, Bar-tell and Niederhauser concluded: 1. The films are formed by material from the oil adsorbed at the interface rather than by a reaction at the interface. 2. The adsorption of the film-forming material is a rather slow process. 3. There are good indications that the process is reversible; i.e., the film constituents may re-dissolve into the oil. 4. The film will spontaneously spread over any new surface formed. 5. The process of adsorption at the interface is accompanied by a decrease in the interfacial tension of the system. Although they had proved that film formation was not due to a reaction at the interface, Bartell and Niederhauser were unable to demonstrate conclusively whether these film-forming constituents were present in the original reservoir oil or whether they were formed as oxidation products after exposure of the oil to air. Recent evidence seems to point to their presence in the original oil. Although the true nature of the film-forming constituents is still uncertain, these conditions would be fulfilled by the assumption of the presence of a slightly soluble surface-active component in the original oil. This component would migrate to a freshly formed interface because it would thereby lower the free energy of the system. The component at the surface constitutes an expanded monolayer which would become compressed, upon retraction of drop, due to the slow rate of re-solution of the film components into the bulk of the oil. It is possible, by involved extraction processes, to separate or at least concentrate this film-forming fraction. However, using the assumptions made above, the technique of monolayer film study may be applied to reveal much of the nature of the adsorbing material without resorting to an actual separation of the film-forming constituents.

Jan 1, 1952

By M. S. Smani, P. Blazy, J. M. Cases

Investigations on the beneficiation of sedimentary Moroccan phosphate ores are reported in four separate parts. The first discusses electrochemical phenomena at the solid-aqueous solution interface for minerals of the apatite group. In the second part, the electrochemical properties of the calcareous gangue associated with the phosphate ore are investigated. The flotation of calcite and phosphate oölites, using both anionic and cationic collectors, is discussed in the third part. It is shown that selective recovery could be obtained for precise concentrations of collectors and pH ranges. Results of work aimed at optimizing depression of phosphate oölites during inverse flotation of the carbonate gangue with sodium oleate are reported in Part 4.

Jan 1, 1976

By William E. Ford, Edward Salisbury Dana

A. Anhydrous Carbonates The Anhydrous Carbonates include two distinct isomorphous groups, the CALCITE GROUP and the ARAGONITE GROUP. The metallic elements

Jan 1, 1922

Jan 1, 1891

By R. Schuhmann, O. W. Moles

at temperatures of 1150°, 1250°, and 1350°C for liquid copper sulphides ranging in composition from saturation with Cu to about 21.5 pct S. From the experimental data, activities of Cu, S, and Cu2S in the melts were calculated. Also, the results furnish a new determination of the location of the curve showing the compositions of Cu-saturoted sulphide melts in the CU-S constitution diagram. THE "white metal" made as an intermediate L product in converting copper matte to blister copper is a matte approximating Cu2S in composition. It has long been known that liquid Cu-S mattes can exist at a given temperature over an appreciable composition range, with S:Cu ratios ranging up from a lower limit corresponding to equilibrium of the matte with a liquid Cu phase. The Cu-saturated matte has somewhat less S than Cu2S, although previous measurements of the compositions of Cu-saturated mattes are not in good agreement.'" As the S:Cu ratio approaches and slightly exceeds that for Cu2S, the partial pressure of sulphur gas over the liquid matte rapidly approaches atmospheric pressure,5 so that Cu-S melts with much more S than Cu2S can be made and handled only in pressure apparatus. Accordingly, the liquid copper sulphides of interest in copper smelting and readily studied experimentally fall in a relatively narrow composi- tion range about Cu2S. The freezing points of Cu-S mattes show a maximum of 1129°C near the composition Cu, dropping to 1105oC for Cu-saturated mattes having less S and dropping under 1100°C for higher S mattes.' Estimates of the sulphur pressure and other thermodynamic properties of liquid Cu,S were made by Kelley; based on extrapolations from equilibrium measurements on solid Cu,S. More recently, Cox and coworkers' studied equilibria in the reaction: Cu,S + H, e 2Cu + H,S over the temperature range 700" to 1250°C. They measured equilibrium H2S:H2 ratios presumably with both the sulphide and the metal phase present, so that their data above the melting point of Cu,S should measure the S pressure of Cu-saturated sulphide melts. The sulphur pressures of liquid copper sulphides of several S:Cu ratios at their freezing points can be estimated from data given by Posnjak, Allen, and Merwin." However, beyond these few data, no systematic study seems to have been made of the relations of sulphur pressure and other chemical properties of liquid copper sulphides to composition and temperature. The measurements of the sulphur pressures of liquid copper sulphides reported in this paper represent a first step in the quantitative study of the chemistry of mattes. Similar work is planned on

Jan 1, 1952

Oct. 8-13, 1917 COMMITTEE ON ARRANGEMENTS H. A. BUEHLER, Chairman, W. E. McCOURT, Secretary, C. J. ADAMI, Southeastern, J. A. CASELTON, Entertainment, F. W. DEWOLF, Book, , E. F. GOLTRA, Finance, L. L. HUTCHISON, Tulsa, E. McAULIFFE, Transportation, P. N. MOORS, Reception, MRS. P. N. MOORE, Ladies, J. D. ROBERTSON, Publicity, VICTOR RAKOWSKI, Joplin., H. A. WHEELER, Excursions, A. THACHER.

Jan 7, 1917

By Y. G. Shiau, J. E. Burke

Several attempts have been made to account for the fact that grains in a fully recrystallized metal will coarsen on annealmg. Two fundamentally different hypotheses have been advanced, with several variations of each. One school considers that the cause of grain growth is the surface energy of the grain boundary. The theory of Jeffriesl that grain size and grain size contrast control growth is an example. The second school considers the primary cause to be a difference in the lattice energy on either side of the grain boundary. The more perfect grains are considered to grow at the expense of their less perfect neighbors, with a decrease in free energy, and a consequent increase in the stability of the system. Burgers2 presents the case for the strain or lattice imperfection theory of grain growth quite completely. The latter theory is attractive, since there is no doubt that differences in lattice energy or lattice perfection can cause grain boundary migration. A recrystallization nucleus grows for this reason as was beautifully demonstrated by van Arkel and Ploos van Amstel.³ Although it is generally stated that grain growth does not precede recrystallization, several workers4-5 have shown that small strains may induce a single crystal in an aggregate to grow, or at least may cause some grain boundary migration. The origin and nature of the imperfections that are considered to be responsible for normal grain growth have never been clearly described, but it is generally agreed that they are a consequence of the difficulty of atomic reorientation in the solid state, and that they do not seriously differ from the type of imperfection that can be introduced by mechanical deformation. An object of the present work was to determine whether the introduction of mechanical deformation would cause grain growth in a specimen which would not otherwise show it under the conditions used. It was also planned to determine the effect of such deformation on the rate of growth in a specimen which would show growth in the unstrained condition. Recently Maddigan and Blank6 have indicated that slightly strained specimens of alpha brass can undergo considerable growth prior to recrystallization if annealed at low temperatures. French7 has reported that at deformations less than 17 pet it is very difficult to detect the beginning of recrystallization in the same material. It was therefore planned to broaden this work to include a complete study of the microscopic behavior of alpha brass under condition~ of temperature, time, grain size and deformation such that recrystallization would not occur or would occur to only a Small extent. Experimental Procedure The brass was prepared by melting cathode copper and commercially pure zinc

Jan 1, 1949

By J. E. Burke, Y. G. Shiau

Several attempts have been made to account for the fact that grains in a fully recrystallized metal will coarsen on annealmg. Two fundamentally different hypotheses have been advanced, with several variations of each. One school considers that the cause of grain growth is the surface energy of the grain boundary. The theory of Jeffriesl that grain size and grain size contrast control growth is an example. The second school considers the primary cause to be a difference in the lattice energy on either side of the grain boundary. The more perfect grains are considered to grow at the expense of their less perfect neighbors, with a decrease in free energy, and a consequent increase in the stability of the system. Burgers2 presents the case for the strain or lattice imperfection theory of grain growth quite completely. The latter theory is attractive, since there is no doubt that differences in lattice energy or lattice perfection can cause grain boundary migration. A recrystallization nucleus grows for this reason as was beautifully demonstrated by van Arkel and Ploos van Amstel.³ Although it is generally stated that grain growth does not precede recrystallization, several workers4-5 have shown that small strains may induce a single crystal in an aggregate to grow, or at least may cause some grain boundary migration. The origin and nature of the imperfections that are considered to be responsible for normal grain growth have never been clearly described, but it is generally agreed that they are a consequence of the difficulty of atomic reorientation in the solid state, and that they do not seriously differ from the type of imperfection that can be introduced by mechanical deformation. An object of the present work was to determine whether the introduction of mechanical deformation would cause grain growth in a specimen which would not otherwise show it under the conditions used. It was also planned to determine the effect of such deformation on the rate of growth in a specimen which would show growth in the unstrained condition. Recently Maddigan and Blank6 have indicated that slightly strained specimens of alpha brass can undergo considerable growth prior to recrystallization if annealed at low temperatures. French7 has reported that at deformations less than 17 pet it is very difficult to detect the beginning of recrystallization in the same material. It was therefore planned to broaden this work to include a complete study of the microscopic behavior of alpha brass under condition~ of temperature, time, grain size and deformation such that recrystallization would not occur or would occur to only a Small extent. Experimental Procedure The brass was prepared by melting cathode copper and commercially pure zinc

Jan 1, 1949

For the seventh consecutive year, the value of nonfuel mineral raw ma¬terials produced in the US set a new record, reaching almost $20 billion-up 12.7% from the previous year, according to the US Bureau of Mines. The Bureau reported that this mineral production was the basis for about $200 billion worth of processed materials of mineral origin.

Jan 5, 1979

By E. R. Gilbert, D. E. Munson

Creep of copper under 75 to 1.50 kg per sq cm stresses at temperatures near the melting point was found to he a complex reaction controlled by three mechanisms acting in parallel. In order of appearance with decreasing temperature, the auerage activation energies, Qc , are 168, 79, and 24 kcal per mole. Stress dependence of the minimum creep rate was found to he an exponential for the two high-Qc processes and a power law for the low-Qc, process. Transition of control occurs from one mechanism to another. The relative transition temperature depends upon the applied stress, and the range oiler which the transition occurs depends upon the difference in the activation energies of the mechanisms. The creep behavior at high temperature is explained by the climb of dislocations through thermally or mechanically formed jogs, CREEP in pure fcc metals at temperatures in excess of one half the absolute melting point is normally controlled by dislocation c1imb.1,2 A climb model which seems applicable, according to an extensive analysis of data,3 was derived by Weert-man.4 This model assumes jog-saturated dislocations and predicts an activation energy, Qc, nearly equal to that for self-diffusion, USd. Although the requirement of jog saturation is restrictive, agreement between theory and experiment seems adequate. Many other theoretical treatments, including an early model by Mott,5 include detailed consideration of jog formation as an initial requirement for climb. These models predict activation energies for creep which differ from those of self-diffusion. Seeger6,7 postulates an observed activation energy related to the stacking-fault energy. Thus, Usd <Qc<5Usd + Uj where Uj is the jog-formation energy. Seeger incorporated qualitatively the influence of the relative numbers of thermal and athermal jogs. Expanding this concept, Shoeck8 explicitly states a function based on formation mode and relative numbers of vacancies: e r ci exp {-uf/k T} exp {-Um/k T} [ 1 ] where Uf and Um are energies of vacancy formation and migration, respectively. The concentration of jogs, Cj, depends upon the manner of jog production. For intersection jogs, Cj is not sensitive E. R. GILBERT, Junior Member AIME, formerly with De- to temperature; for thermal jogs, Cj is proportional to exp {—Uj /k T}. Schoeck regards each mode as a distinct mechanism; therefore, the mechanisms may act together.299 The diversity predicted by theory, surprisingly, has not been substantiated by experimental results. A significant investigation must include the extremes in stacking-fault energy. Extensive creep studies of aluminum10 and nickel,11,12 high stacking-fault energy metals, have been made. Comparable studies on a low stacking-fault energy metal, such as copper, have not. It is the purpose of this paper to report the results of an investigation of the creep of copper under conditions which favor thermal jogs. EXPERIMENTAL Cylindrical compression creep specimens (0.240 in. in diameter by 0.400 in. long) were machined from cold-drawn rods of electrolytic tough-pitch copper containing 0.0007 Mg, 0.002 Fe, 0.001 Ni, 0.0005 Ag, Cd < 0.005, and Pb < 0.005 wt pct impurities. Undetected spectrographically was a nominal 0.04 wt pct 0, which occurs as a Cu2O constituent distributed discontinuously at grain boundaries. Vacuum annealing at 900°C for 15 min produced a stable 0.03-mm average grain diameter. Testing was carried out using apparatus similar to that described by sherby,13 modified by enclosing platens and a push rod in a vacuum cylinder. Normally this arrangement resulted in pressures less than 10 Only a slight surface tarnish, less than 0.0005 in. in thickness, occurred during the test. The applied stress, corrected for atmospheric pressure, was maintained within 2 pct of the desired true stress by the addition of lead shot at fixed strain increments. Creep strain was measured with dial gages as a relative displacement of the upper and lower platens; accuracy of measurement was 0.0001 in. Two creep-test methods were used, the differential or cyclic temperature14 and the isothermal, to obtain creep data at stress levels of 150, 100, and 75 kg per sq cm over the temperature range of 620° to 1032°C. Minimum creep rates were used from both test methods; this was considered proper because comparable temperature tests or cycling to the same temperature gave the same creep rate, within experimental error. The cold vacuum test chamber, with the unstressed specimen in place, was heated to temperature by placing a preheated furnace over the chamber. Temperature equilibrium was attained within 30 min. For the cyclic tests, the stress was removed during the 5 to 10 min necessary to effect the temperature change and re-

Jan 1, 1965

By J. E. Korski

Quebec Cartier Mining Co. operates a low-grade iron ore deposit at Lac Jeannine, Que., which is located at the extreme southwestern end of the Quebec-Labrador Trough (Fig. 1). The facilities there include one of the world&apos;s largest concentrators, a hydroelectric plant of 60,000 hp, and the modern town of Gagnon. A 191-mile railway connects the mine to the seaport situated at Port Cartier on the St. Lawrence River. The harbor in Port Cartier can accommodate both lake and large ocean- going vessels and is also the site of a major grain terminal. The concentrator in Lac Jeannine is designed to produce some 8 million tpy (tons per year) of specular hematite concentrate, containing 66% Fe, from approximately 20 million tons of crude ore. An annual stripping program of 5 to 8 million cu yd of waste rock uncovers the required ore. Approximately 100,000 tons of ore and waste are removed daily on a year-round, 24-hr-per-day basis. The open pit covers an area with maximum dimensions of 8,000 x 2,600 ft (Fig. 2). Present mining plans call for an ultimate depth of approximately 1,000 ft below original ground level. The rocks in the pit area are typical of this portion of the Grenville geological province. They consist largely of marble, quartz rock, pegmatite, iron formation, and stratigraphically older gneisses. The iron-ore formation consists of medium to coarse-grained specular hematite in a quartz host rock. The formation occurs locally in a plunging, anticlinal form. The topography in the vicinity of the mine is rolling, with maximum variations in elevation of about 400 ft. The hills are thickly wooded with birch, black spruce, and some tamarack. Lakes and muskeg areas are plentiful. Precipitation and temperature records have been kept since 1959. The annual rainfall has ranged between a low of 23 in. and a high of 42 in.

Jan 1, 1969

By T. E. Leontis

The importance of magnesium alloys in the manufacture of aircraft engines has been realized for many years. A concentrated effort has been exerted in the laboratories of the Dow Chemical Co. to develop magnesium-base alloys with improved properties at elevated temperatures whereby greater advantage can be taken of the inherent lightness of these materials. In an earlier publication, attention was called to the superior tensile properties and resistance to creep of cerium-containing magnesium alloys at temperatures up to 700°F. More recently, additional data on magnesium-cerium alloys have been furnished by Murphy and Payne. The present paper deals with the mechanical properties of sand-cast magnesium alloys containing zinc, at temperatures up to 500°F. The study includes binary Mg-Znt alloys containing up to 10 pct zinc and several ternary and some polynary alloys based on the Mg-~Zn,f Mg-jZn, Mg-4.4211, Mg-6Zn, and Mg-roZn binaries. The compositional variation in the tensile properties, hardness, and creep at room and elevated temperatures, and in the electrical conductivity at room temperature has been determined. The results of these tests show that many zinc-containing magnesium alloys exhibit resistance to creep at elevated temperatures significantly higher than that of present commercial magnesium alloys, but lower than that of alloys of the magnesium-cerium type. In addition, they have high tensile properties at room and elevated temperatures and high conductivity. This combination of mechanical properties and conductivity indicates that these alloys should be given serious consideration for commercial applications involving exposure to temperatures higher than those to which present commercial magnesium alloys can be taken safely. Literature The properties of sand-cast magnesium alloys containing zinc as a principal alloying element have not been investigated very extensively, especially at elevated temperatures. The following remarks, which are confined to sand-cast alloys, will serve as a summary of the most important references on this phase of magnesium metallurgy. The first recorded determination of the tensile strength of binary Mg-Zn alloys as a function of zinc content was reported by Aitchison. Similar studies, including the measurement of elongation and hardness, have been performed by Gann and Winston,&apos; Dumas and Rockaert, Haugh-ton and Prytherch,6 and by Spitaler, the last being reported by Beck.&apos; A slight superiority in the tensile strength of chill-cent Mg + 8 es alloy Over the sand-cast alloy was noted by Maybrey.8 That alloys of magnesium containing 5 to 12 pct zinc

Jan 1, 1949

By Nicholas J. Grant, Yoichi Ishida, Arthur W. Mullendore

An inert particle -marker technique was developed to provide a direct measurement of grain boundary sliding during creep in tile interior of aluminum specimens. Groin boundary sliding in the interior as measured by this technique was found to be nearly the same or slightly lower than that measured on the surface. These data disagree with those obtained by the grain-counting technique developed by W. A. Rachinger. In tests where both techniques were used, the grain-counting technique gave large and variable values of grain boundary sliding. It is shown that the grain-counting technique gines erroneozcs results because of a preferred direction of grain growth during creep. A question which has long plagued those who study grain boundary creep is whether or not surface measurements yield a true representation of the deformation in the interior of a specimen. The first attempt to measure grain boundary sliding in the interior of a creep specimen was made by Rachinger.1 His technique is based on the measurement of average grain diameters, both parallel and perpendicular to the tension axis, followed by a calculation of the grain elongation from these measurements. The difference between total elongation, Et, and the calculated grain elongation, Eg, is considered to be the result of grain boundary sliding, Egb. Surprisingly large values of Egb/ Et (90 pct) came out of creep tests of pure aluminum at temperatures above 250°C (480°F), utilizing the Rachinger technique. This value is very much larger than the values of 8 to 15 pct measured by numerous investigators on the surface of a specimen. Values obtained on the surface by Rachinger, however, were consistent with those measured by the displacement of reference scratches. Chaudhuri and Grant2 repeated the Rachinger method for A1-10 pct Zn, and found the same large values of Egb/Et inside the specimen. However, from purely geometrical considerations, grain boundaries cannot slide without grain deformation; therefore, they questioned the validity of Rachinger&apos;s grain-counting method. Several authors2,3 have suggested that grain boundary migration during high-temperature deforma- tion may tend to restore the equiaxed grain shape in order to minimize interfacial energy. Rachinger4 tested this hypothesis and found that it did not hold. Grains elongated by rapid extension retained their elongated shape during a subsequent slow creep test. It appears that some other factor was responsible for the abnormally high values of .Egb/El. McLean and Gifkins5 made a survey of the effect of grain size on the ratio Egb/ Et, and proposed that the high values were the consequence of a small grain size. They failed to explain the low value of Egb/Et which Rachinger observed on the surface of the small grain size specimens, and suggested that the surface effect on the value of Egb/Et does not always occur. Many of the disagreements arising from Rachinger&apos; s work stem from the fact that there was only one method of estimating grain boundary sliding in the interior of a specimen, and that method was of an indirect nature. If a marker line of some sort could be introduced inside of the specimen, one could make measurements just as clearly as on the surface. In the present investigation, the technique utilizes a layer of finely dispersed oxide particles inside the specimen introduced by hot press-bonding of two pieces of aluminum. Rachinger attempted a similar technique but with the oxide distribution he obtained there was so much interference with grain boundary motion that quantitative measurements were not attempted. MATERIALS AND EXPERIMENTAL PROCEDURE A) Materials Preparation. Two pure aluminum rods (impurity content in percent: Si, 0.002; Cu, 0.004; Fe, 0.002; Mg, 0.000: Zn, 0.000; V. 0.001) 1 in. diameter by 2 in. high were jacketed end to end in a 4-in.-diam cylinder of commercial-purity aluminum, 4 in, high, which had a l-in.-diam hole. The l-in.-diam mating surfaces of the two pure aluminum rods had been electropolished initially. This interface provided the oxide-film internal -marker plane after the hot press-bonding process, see Fig. 2. The composite was annealed for 1 hr at 900 F and hot upset more than 50 pct reduction in height in one step. The resulting slab was then rolled to 1/2 in. thickness; the pure aluminum test material was cut out and annealed at 900°F for 1 hr and cold cross-rolled to 1/8 in. thickness. The sheet was machined into specimens with a 1/8-in.-square cross section and a 1/2-in. gage length. The following heat treatment was given the specimens: Specimen 5A series: annealed at 1000°F for 5 min for grain-size control and then cooled to 700°F and held for 15 hr for stabilization.

Jan 1, 1965

Jan 1, 1879

By A. E. Bibb, F. W. Kunz

A platelet form of zirconium hydride was found in zirconium and ZY-1 wt pct U single crystals containing hydvogen in the range of 50 to 100 ppm. The habit planes for the hydride plateletg in the zirconium crystals were found to be {1012}, {1121}, and (1122) while the habit planes in the ZY-1 wt pct U crystals were {1121), {1123}, and (1.231). With the exception of the {1231} plane, these have been identified as the twin planes in zirconium. The different effect of hydridc on the mechanical puoperlies of zirconium at high and low strain rates has been explained in terms of the defornzation mechanism and the habit planes of the hydride. ZIRCONIUM and zirconium-base alloys are considered for many nuclear reactor applications, where the pick-up of hydrogen is a distinct possibility, if not an unavoidable occurrence. An excellent example of this is the absorption of hydrogen by zirconium and its alloys during aqueous corrosion, where as much as 30 pct of the hydrogen produced by the corrosion reaction is picked up by the metal. Under the proper conditions, the absorbed hydrogen would precipitate as zirconium hydride and consequently alter the mechanical and physical properties of the zirconium-base alloys. The knowledge obtained from a study of the crystallographic aspects of the hydride precipitation is necessary for the interpretation of the observed changes in the mechanical and physical properties. The deleterious effect of hydrogen on zirconium-base materials is not always observed in normal slow tensile loadings at or above room temperature, but becomes immediately obvious under impact loading, where the energy absorbed decreases with increasing hydrogen content,&apos; and in notched specimens where triaxial loadings are attained.2 This effect is of serious consequence in the applications of Zr where shock loads or high strain rates are experienced. The purpose of this investigation was to determine the crystallographic planes of the Zr and Zr-1 wt pct U matrices upon which zirconium hydride will precipitate and to correlate these with the observed mechanical property data. PROCEDURE Crystal Growth—Single crystals of Zr and Zr-1 wt pct U alloy were used in this investigation. The Zr crystals were made from sponge that had a typical analysis as shown in Table I. Large grains were grown in this material by rapid cooling an arc-melted charge in a button furnace. This operation undoubtedly lowered the volatile impurity concentration below those listed in Table I. Back-reflection Laue patterns of these large grains showed the presence of considerable substructure and lattice distortion within each grain. It was found that an 18-day anneal in an argon atmosphere at 840°C (a high a heat-treatment) followed by furnace cooling, resulted in the further growth of the large grains and removed the detectable substructure ana lattice distortion. Equiaxed crystals with a cube edge of approximately 1/4 in. were produced in this manner. Large crystals of Zircaloy-2 (1.5 wt pct Sn, 0.15 wt pct Fe, 0.1 wt pct Cr, 0.05 wt pct Ni) also were produced by this technique but the large amount of substructure formed could not be removed by annealing. The Zr-1 wt pct U solid solution crystals were produced from an alloy rod consisting of Westing-house-Foote hafnium-free crystal bar zirconium and depleted uranium. The chemical analysis of the alloy is also given in Table I. Three-inch sections were cut from an 0.120-in.-diameter alloy rod and tapered to a fine point on each end. The samples were sealed in evacuated Vycor tubes, homogenized for 5 hr at 1000°C in the ß-phase field, and slowly lowered in a gradient furnace into the a phase (800°C), annealed for 60 hr at this temperature to produce grain growth and furnace cooled. This technique produced large grains approximately 1/2 to 3/4 in. long over the entire cross section of the rod, which were free of observable substructure. Metallographic examination up to X1000 did not show the existence of any second phase in these grains. Crystal Orientation—All crystals were abraded on fine emery paper to produce two flat surfaces 90 deg to each other. The crystals were analyzed by the standard Back-Reflection Laue technique to obtain the crystal orientation with respect to the abraded surfaces. The Back-Reflection Laue Photograms of all crystals used were clear and sharp indicating that the crystals were free from any substructure that could be detected by these means. Hydriding—Only samples containing four to five grains were used for hydriding. Samples that exhibited fine grained patches were discarded. The samples were hydrided using the following technique. After etching in 48 vol. pct HNO3, 48 vol. pct H2O, and 4 vol. pct HF (48 pct) solution, the specimens

Jan 1, 1961

Jan 1, 1930

By C. R. Austin

In a recent paper1 the authors presented information on a refinement of the Rohn type of creep test with data on pure iron that exemplified the behavior of the apparatus. The present paper extends that work to nickel and cobalt and to the examination of the comparative high-temperature properties of two widely used commercial alloys. Careful attention has been given to the effect of plastic deformation in decreasing creep rate, using the elementary metals iron, nickel and cobalt for purposes of analysis. In any endeavor to develop new materials for high-temperature service one soon encounters a real problem as to the sort of test to be employed in evaluating the commercial possibilities of the alloys. At one time it was considered satisfactory to determine the tensile strengths at elevated temperatures and to use these values as a criterion for the commercial application of the alloys. More recently the proportional limit has been used as a means of classification. Data obtained from a form of bend test have also shown promise as providing a "merit index" of high-temperature serviceability.= Koerber and Pomp3 state that since the influence of testing time is small, the high-temperature yield point of steel furnishes sufficiently reliable and concordant values up to about 350" C. (and a maximum of 450" C. for alloy steels) to serve as a suitable basis for comparison evaluations of materials and for judging permissible stresses. Koerber and Pomp consider, however, that above these temperatures creep limit must be substituted for the high-temperature yield point as the all-important characteristic of the material. Whatever may be their definition of

Jan 1, 1934