Search Documents

By R. W. Fountain, M. Korchynsky

The precipitation phenomena occurring in cobalt-tantalum alloys have been investigated in the temperature range frm 500" to 1050°C by correlating the results of metallographic, X-ray, micro-and macrohardness, and electrical resistivity studies. The property andmacrohardness,changes were found to depend on 1) general precipitation, and 2) lamellar precipitation. Two new intermetallic phases have been identified: 1) a Co3Ta, a metastable ordered face-centered-cubic compound, and 2) a stable ß Co3Ta phase of hexagonal structure. In addition, the previously reported Co2Ta phase was found to exist in two allotropic modifications: the hexagonal MgZn,-type and the cubic MgCu2-type Laves phases. SINCE a large variety of structures can result as a consequence of the decomposition of a solid solution, predictions on the nature of property changes are difficult, if not impossible, to make. For any rational attempt to correlate properties and structures of a precipitation-hardenable alloy, a detailed understanding of the kinetics of decomposition and morphology of phase separation, as well as knowledge of phase relationships, appears to be prerequisite. Information of this type has been accumulated in the past for many alloy systems, both of theoretical and pastforpractical importance.1,2 Although the presence of intermetallic compounds has been reported in cobalt-base alloys,3 the amount of published information on precipitation-hardenable cobalt-base systems is very limited. A survey of the binary phase diagrams of cobalt indicates that cobalt-tantalum alloys might be of interest as typical of other cobalt-base systems in which Laves phases of the A,B type can be precipitated from solid solution. The present work has been undertaken, therefore, to study the kinetics and morphology of the precipitation reaction in this system and to establish a base for a correlation between the structural aspects and properties in this class of alloys. PREVIOUS WORK The only available phase diagram of the cobalt-tantalum system is based on the work of Koster and Mulfinger. According to these authors, the maximum solubility of tantalum in cobalt is about 13 pct (at 1275°C) and. less than 7 pct at room temperature. Tantalum additions lower the temperature of allotropic transformation of cobalt (about 420°C), and at 7 pct Ta, the high-temperature face-centered-cubic modification (ß cobalt) is retained at room temperature. The precipitating phase was originally designated as Co5Ta2 compound (55.2 pct Ta, about 1550°C melting point), but subsequent investigations by wallbaum5" identified this constituent as the A,B-type Laves phase. Wallbaum's data indicate that there are two modifications of this intermetallic compound: one richer in cobalt (Co2.2 Tao.8)of the hexagonal MgNi, type; and another of a higher tantalum content (Co2Ta) of the cubic MgCu, type. On the other hand, Elliott7 found that the cobalt-rich alloy (CO2.10,Tao.~l) was predominantly the cubic MgCu, type at 800°C and a mixture of both the MgCu2 and the hexagonal MgZn,-type Laves phases at 1000°C. At 1200°C, Elliott found only the MgZn, type while at 1400°C, he observed only the MgCu2 type. At the stoichiometric composition, Co2Ta, Elliott reported only the cubic MgCu2-type Laves phase in the temperature range of 600oto 1600°C. The precipitation of the cobalt-tantalum intermetallic compound is accompanied by a marked increase in hardness. According to Koster's4 data, the Brinell hardness of an 8 pct Ta-Co alloy increases from 230 to 340 upon short-time aging at 800°C. EXPERIMENTAL PROCEDURE The binary cobalt-tantalum alloys investigated contained 5, 10, and 15 pct Ta. The range of tantalum additions was thus slightly broader than the reported minimum and maximum solid solubility limits of tantalum in cobalt (7 and 13 pct, respectively)4 The alloys were vacuum-induction melted in a magnesia crucible using cobalt rondelles and technically pure tantalum sheet as raw materials. Deoxidation of the melt was accomplished with carbon, and the chemical analysis of the alloys is given in Table I. The effect of isothermal aging treatments on the progress of precipitation was studied on samples cut from cast ingots. These samples were solution treated for 2 hr at 1250°C and water-quenched. Aging was conducted in the temperature range from 500" to 1050°C for periods between 15 min and 1000 hr and followed by water-quenching. To prevent contamination from the atmosphere, all samples were sealed in evacuated Vycor or quartz tubes for heat-treatments. For solution treatment, argon at 0.2 atmospheric pressure was introduced prior to sealing of the capsule to prevent collapse at high temperature, and titanium sponge was placed at one end of the capsule to act as a getter. MACROHARDNESS The effect of aging on Vickers hardness (Dph) of

Jan 1, 1960

By R. H. Harrington

THE definition of precipitation-hardening is well understood and its principles have been subjected to study for some time. However, the variation of properties with double aging, combined with strain-hardening, gives promise of new and useful phenomena. Properly to understand these distinctions, definitions of these terms are necessary. Precipitation-hardening. This reaction usually depends upon an alloy having such a composition that it consists only, or in major part, of but one phase, a solid solution, at a high temperature, so that, upon slow cooling, this solution breaks down into two phases, the solid solution and a compound. By heating the alloy to the temperature for maximum solution of the compound, quenching, and reheating to a suitable lower temperature for critical dispersion precipitation of the compound, useful properties of the alloy may be developed to a maximum. Double Aging.-Dougle aging in part is similar to the process for precipitation-hardening, consisting of heating a precipitation-hardening alloy to the solution temperature, quenching, reheating to the suitable dispersion precipitation temperature, strain-hardening, and again reheat-ing to a suitable temperature to develop certain useful properties to a maximum. The strain-hardening may be induced by two methods: (1) a rate of cooling from the first reheat may suffice to strain the solid solution lattice of the alloy; or (2) the alloy may be cold-worked to a suitable degree between the two reheats.

Jan 1, 1936

By J. J. Beeson

Page I. PrimaRy MineRalization............................................ 356 1. Introduction..................................................... 356 a. Scope of Work............................................... 356 6. Acknowledgments............................................ 357 c. The Bingham Mining District.................................. 357 d. General Geology.............................................. 357 e. The Utah Copper Mine....................................... 376 f. Field and Laboratory work.................................... 376 2. The Quartz-Monzonite.............................................. 377 3. The Pegmatite Veins.....:.•..............................:........ 378 4. The High-Temperiture Veins.........................-.............:379 5. The Early Hydrotherma! Period.:......................'............ 380 6. The Late Hydrothermal Period................................:.... 385 II. Seconuabx EnricuMENT............................................. 386 . 1. Author's Note;..... ;............................................ 386 2. Factors Governing Enrichment.....................................: 387 3. The Enrichment.................................................... 387 4. The Covellite Enrichment.; .:; ; : .:. :..:.............:................ 388 a. The Capping above the Covellite Ores:....:....................... 388 b. Conclusions Drawn. from a Study of the Capping and the Covellite ; Ores....,.............................................:.....,.' 392 c. Covellite Enrichment at Depth......:......................... : : 392 5. The Chalcocite Enrichment...................................... ; 394 a. The Carbonates in the Capping above .the Chalcocite Ores...:..: .. 394 ' b. The Formation of Chalcocite..:.:. :......:..'...........,......' 394 . 6. Secondary Chalcopyrite and Bornite.............................i ..., . 395 111.- Conclusions as to the Formation.oF the ImpoRtant Sulphides in the Disseminated ORes........................................,.......... 400 I. Primary MiNeRalization 1. Intrdtiction a. Scop. e .of Work.. first part bf, this paper deals with the country rock, its pegihatitic, pneumatolytic', and hydrothermal altektions; also, the effect of magnetite on the deposition of the primary chalcopyrite and pyrite.

Jan 1, 1916

By L. N. Cosner, C. F. Austin, J. K. Pringle

Shock wave attenuation studies with three igneous rocks are reported: a spessartite elastic to low intensity stress waves, a diorite which is anelastic, and a porous scoria. The level of shock passing through a given thickness of rock has been determined by donor-receptor methods using high speed streak cameras. Test results include shock propagation velocities adjacent to the explosive, the dependence of test results upon test geometry, and the variation in shock wave attenuation in these three common igneous rock fabrics. The utilization of explosives in rock mediums involves a series of interrelated phenomena. These are the detonation process within the explosive, the transfer of energy from the explosion products to the surrounding rock medium in the form of sound or shock waves, and the transmission of the resulting waves through the rock medium to create damage in the rock mass. The subject of this paper is the investigation of the close-in phenomena that occur in the intensely shocked rock immediately adjacent to a detonating explosive charge. The specific aspects studied include propagation velocities, free surface velocities, pulse attenuations and the determination of the equations of state for the test rocks employed. In addition, a series of tests was conducted to indicate the dependence of the test results on the charge diameter utilized. The equations of state for rock materials have been investigated by a number of authors. Lombard, ',' Grine and Fowles,3,4 Zharkov, 5 Hughes and McQueen,6 Chabai,7,8 and Bass9 have published equation of state data on rock materials similar in structure or composition to those investigated in the study. Rinehart10 has prepared a voluminous tabulation of the available equation of state data from these publications and in addition has included data on many other rock and earth materials in his compilation. The reader is referred to the extensive tabulations and bibliography of Rinehart for classical studies involving the equations of state of metals and other solids and liquids. The problems of sympathetic detonation between drill holes in hard rock environments has received the attention of various authors as have the problems of explosive rock coupling close-in stress wave attenuation and explosive diameter effects. 11 to 18 In particular, the don or-receptor geometry of this report is very similar to the geometry of adjacent drill holes in rock. Most published pulse velocity and attenuation work in rock has been devoted to velocities at much lower amplitudes than those associated with close-in explosives phenomena and most attenuation studies involve transmission and attenuation across large spans of material. Excellent work in these areas has been published by the U. S. Bureau of Mines and by many individuals and groups engaged in geophysical research. SELECTION OF TEST MATERIALS A series of rock materials was desired for this test program that would be meaningful industrially, and that would also yield comparative results from a research viewpoint. Based on the availability of a series of related petro-fabrics, for which there is an abundance of physical and shock response data, three igneous rocks were selected as the test series. The applicability of shock wave studies in igneous rocks to industrial problems is readily demonstrated. Hardrock tunneling, quarrying and mining take place in igneous rocks throughout the world. From a research point of view, igneous rocks exhibit a wide variety of grain sizes and considerable diversity in composition, yet the common igneous rocks have a general similarity in their constituent mineral structures and in their grain-to-grain bonding mechanisms. The latter is based on the concept of rock formation through magma cooling and crystallization, though the degree of liquefaction in granitic plutonic masses is indeed a moot question. Three specific types of igneous rock were chosen for study. A coarse-grained granitic textured rock was desired that would serve as a representative of the widely abundant intrusive rocks of the world and

Jan 1, 1967

By E. P. Klier, Taylor Lyman

THE structures formed when austenite is quenched to subcritical temperatures and allowed to transform isothermally have been the subject of intensive study since the work of Davenport and Bain.1 Isothermal transformation diagrams summarizing the results of many of these studies have become widely familiar to metallurgists. Of particular interest to metallographers are the dark-etching, acicular products of transformation formed by isothermal reaction below the temperature of maximum velocity of the austenite to pearlite reaction. These products, the bainite structures, can be formed in a wide variety of steels. However, the constitution of bainite is not well understood and divergent views have been expressed as to its mode of formation. In this investigation a combination of dilatometric, microscopic and X-ray methods has been brought to bear upon the problem in the hope of some elucidation of the bainite reaction as it occurs in hypoeutectoid steels. MECHANISM OF BAINITE FORMATION Davenport and Bain1 considered the mechanism of bainite formation to consist of two steps-an allotropic transformation followed by precipitation of carbide. The separation in time of the two processes was considered to be very slight in the upper temperature range and very marked at temperatures just above the martensite range. This concept has been restated by Vilella, Guellich and Bain2 in the form of a definition of the acicular mode of transformation as: The successive, abrupt formation of flat plates of supersaturated ferrite along certain crystallographic planes of the austenite grains; this supersaturated ferrite begins at once to reject carbide particles, (not lamellae), at a rate depending upon temperature. In effect, this is the acicular mode of transformation, even though the temperature be such as to limit the actual life of the quasi-martensite to millionths of a second. The investigation of isothermal decomposition of austenite in certain alloy steels (notably those containing chromium or molybdenum) has revealed that there may be a range of temperatures between the pearlite and bainite reactions in which the austenite decomposes at a relatively low rate.3-6 Further, it is characteristic that in the second region of fast reaction the decomposition of the austenite is incomplete by one reaction but goes to completion by a second reaction. These observations led Wever7 to a description of the bainite reaction as follows: I. The reaction takes place by initial precipitation of a martensitic intermediate structure. 2. In the upper temperature range this structure readily decomposes into ferrite and cementite. In the lower range carbon separates in an unknown form. 3. In the upper temperature range the precipitated carbide nucleates the precipitation

Jan 1, 1944

By S. Pearson, L. Rotherham

Measurements have been made of the variation of internal friction with temperature for OFHC copper, and for a series of binary solid solutions of high purity copper with zinc, gallium, germanium, arsenic, and silicon, to investigate the effect of alloying elements in substitutional solid solution on grain boundary viscosity. The apparatus used was of the torsional pendulum type developed by KeI1 modified so that the specimen and vibrating system could be maintained in a high vacuum. The activation energy for grain boundary relaxation in copper is 33,000 cal per mol. This is considerably less than the activation energy for self-diffusion, in disagreement with Ke's theory2 that the two activation energies should be the same. All the alloying elements increase the activation energy for grain boundary relaxation to approximately 44,000 cal per mol, increasing the grain boundary viscosity at a given temperature. The results are discussed in terms of theories of grain boundary slip put forward by MottQ and by Ke.+ Experiments have also been made to investigate the effect of small amounts of oxygen on the variation of internal friction with temperature for copper. EXPERIMENTS described in this paper were designed to investigate the effect of alloying elements in substitutional solid solution on grain boundary viscosity, using the internal friction method developed by Kel for measuring grain boundary viscosity. It was considered desirable to make the system under investigation as simple as possible by making experiments on a comprehensive series of binary alloys of copper with the elements immediately following it in the periodic table, i.e., zinc, gallium, germanium, and arsenic. These elements are within the range of favorable size factor for solution in copper, and the systems have been extensively investigated by Hume-Rothery and his co-workers." Later, in order to study the effect of the relative size of the solvent and solute atoms, experiments were made on a series of Cu-Si alloys, since the apparent atomic diameter of silicon when in solution in copper—the diameter necessary to assign to the solute atoms to account for the lattice constant of the solid solutiona—is almost the same as that of the solvent atoms. Preliminary experiments were made to measure the variation of internal friction with temperature for pure copper. Ke made some experiments on copper," but found that the curves obtained were not reproducible in consecutive experiments. The copper was maintained in an argon atmosphere during the experiments but, although he took precautions to exclude oxygen, Ke came to the conclusion that the precautions were not entirely satisfactory, and that the lack of reproducibility must be due to absorption of oxygen by the copper. Experiments have therefore been made to investigate the effect of small amounts of oxygen on the variation of internal friction with temperature for copper. Previous Work Ke made extensive experiments on high purity aluminum and was able to show that the grain boundaries behave viscously at high temperatures. He found that the relative rate of movement of the grains could be expressed as Where v is the velocity of slip, A is the constant, u is the shear stress, H is the activation energy, R is the gas constant, and T is the absolute temperature. By assuming that the boundary layer was one atom thick, he obtained a value for the viscosity which, when extrapolated to the melting point, gave good agreement with the viscosity of the liquid metal at the melting point. Ke also made experiments on other high purity and commercial purity metals:, and on some alloys." " He came to the conclusion2 that the activation energy for grain boundary slip was the same as for self or volume diffusion and for steady state creep in single crystals. Rotherham, Smith, and Greenough" have shown recently that the activation energy for grain boundary slip is not the same as for self-diffusion in tin,

Jan 1, 1957

By R. M. Goldhoff, H. J. Beattie

The high-temperature properties of a 1 Cr-Mo-V forging steel are described. A series of controlled heat treatments was designed to delineate the effects of austenitizing and tempering treatments, temile strength, and grain size on such properties. Studies indicate that the mechanical properties and their varlations under creep can be described by the initial metallurgical structures and their changes droPirzg exposure, particulavly the carbide reactions. Such structures are described and correlated with the mechanical properties. FOR many years the large steam turbine industry has relied on the 1 Cr-Mo-V type forging steel for critical applications. Because of its adequate heat resistance and relative economy, it is currently in use in the temperature range up to about 1050" to 1075°F. Attempts to understand the large property variations attainable in this steel involve the structural modifications due to the wide latitude in its heat treatment. The heat treatment essentially involves a two-step process which includes solution-ing of carbides in the austenite range followed by a suitable tempering treatment below the critical to adjust the level of properties. The latter step is referred to as "secondary hardening" and is basically an ordinary aging reaction involving carbides. In the commercial heat treatment of large components of such steel, homogenizing and stress-relief anneals may be included and have some importance in determining subsequent properties. Several studies of the engineering properties of this steel as a function of transformation micro-structure have been reported.1 3 However, in this steel the carbide reactions, which are a function of composition and heat treatment, appear to be the property-controlling factor rather than the micro-structure defined as a transformation product. Thus, the tempering resistance of this type of alloy steel is mainly a function of the size and distribution of alloy carbides.4 However, it is also necessary to consider the stability of the microstructure and the effect of dynamic carbide reactions on subsequent properties. It is the purpose of this paper to show the interdependence of properties and corresponding structures, particularly carbide reactions, developed for a limited set of controlled heat-treatment conditions applied to 1 Cr-Mo-V steel. MATERIALS AND PROCEDURE The material for this work came from a large production forging whose chemical composition is shown in Table I. The property data accumulated on this steel as a function of heat treatment were room-temperature tensile and smooth- and notch-bar creep rupture at 1050°F (notch geometry: shank diameter = 0.357 in., depth = 50 pct, radius = 0.005 in.). To achieve controlled structural variations the temperature and time of both austenitizing and tempering were varied in a manner to produce a series of eight steels each at one of two grain sizes and one of two hardness levels in proper combination for valid comparisons. This will be clear upon examination of Table I. Structural studies involved the use of optical and electron microscopy as well as X-ray and selected-area electron diffraction. To reveal the nature of carbide precipitations, electrolytic extraction techniques were used with subsequent analysis of the residue by X-ray diffraction. Weight losses of the steel specimens during electrolysis were measured and successive chemical fractionations of the residues were applied and checked by X-ray examination. The details of fine structural distributions, morphologies, and crystallography of the precipitates as well as dislocation distributions were investigated by examining in the electron microscope three common types of preparation. a) Relief Replicas. Mechanically polished sections were etched in 2 pct nital, replicated with a nitrocellulose film, which was shadowed by chromium vapor deposited at a glancing angle of 20 deg. b) Extraction Replicas. A thin film of carbon was vapor-deposited on the polished and etched surface. The carbon films were then etched off and gathered on electron-microscope supporting grids. The carbides left imbedded in the carbon replica in their original distribution were then examined crystallo-graphically by selected-area electron diffraction. c) Thin Films. Specimens were mechanically ground and polished down to a thickness of 0.001 to 0.003 in. Final thinning was done electrolytically in a "chrome-acetic" electrolyte. When holes began to appear in the foil, the voltage was interrupted and applied in several 1-sec bursts. Sections of the foil between holes were thin enough to pass a 100-kv electron beam that carries an image.

Jan 1, 1965

By A. P. Watts

FOREWORD THE problem of concentrating the disseminated lead ore of southeast Missouri is extremely simple. The economic mineral is galena and the gangue is dolomite. The ore assays from 4 to .6 per cent. lead, 80 per cent. of which is recovered in a 70 per cent. concentrate.. The ore is crushed through 10 mm. and sized on 2 mm.; the oversize is jigged, while the undersize is tabled, and the slime treated by flotation. The district is the largest lead producer in the world, the output of metallic lead for the year 1915 being 183,906 tons. Over 20,000 tons of ore a day is being treated. The management of the different properties is progressive and always open to suggestion. The fact; however, must not be lost sight of, that the problem in the district is peculiar, as the average ore contains but 80 to 120 lb. of a metal normally selling at about 4c. per lb. The low price of the metal limits the treatment system to one of low cost. An elaborate treatment is neither logical nor advisable. An engineer first visiting the district may consider that the method of treatment could be improved

Jan 10, 1917

Jan 1, 1924

Jan 1, 1910

By W. F. Chiao

With the application of dislocation damping theory an attempt was made to determine whether the generation and extension of dislocations is inherently more difficult in a brittle steel than in a ductile steel. A ductile steel was compared with a brittle stee1 by simultaneously measuring the ultrasonic attenuation and velocity during tensile test, and the density of free dislocations and their mean loop length were then calculated as a function of strain. The results showed that in the ductile steel there was always a large generation of dislocations and great extension of loop length occurring at some stage within the early plastic region. In contrast, the brittle steel showed very little or no such sudden changes in dislocation dynamic states after the onset of plastic deformation. Furthermore, a strong temperature dependence of dislocation dynamic states was also observed in the ductile steel and a hypothesis was suggested that a thermally activated process of dislocation rearrangement could occur at higher deformation temperatures. The activation energy of dislocation rearrangement at room temperature was estimated as about 2030 cal per mole.C. DUCTILITY is an indispensible property in the application of engineering materials, especially steel. During the past two decades the theoretical and experimental approach to the understanding of flow and fracture of metals has been constantly undergoing changes and progress." while the fracture behavior of metals can be influenced by many factors such as chemical Composition,3 second-phase particle mor-phology,4 and dislocation arrangement,5 it is now a general belief that the fundamental understanding of the ductile-brittle fracture phenomena of solid materials must stem from the study of dislocation dv-namics developed under stress conditions.6,7 Most of the traditional ductility tests, such as Charpy impact test, slow bend test, and tensile fracture test, cannot by themselves reveal directly the mechanisms of ductile to brittle transition of materials. In the experimental investigation of tensile ductility it would be ideal to be able to study directly the dynamics of dis-locations in a bulk specimen during the process of deformation. Since the ultrasonic pulse technique is the only satisfactory method for studying dislocations and the fine details of deformation characteristics in metals in the course of a tensile test, it would appear that a comparative study of ultrasonic attenuation changes during tensile tests of metallic materials exhibiting different ductility might be very informative. So far no work comparable to this study has appeared in the literature. Recent progress in both theory and experiment has indicated the feasibility of studying the dislocation mechanisms of ductility behaviors by ultrasonic measurements during tensile test. Granato and Lucke8 have developed a quantitative theory that enables the calculation of dislocation density and their average loop length from the measurements of ultrasonic attenuation and velocity, and several investigators, including Chiao and Gordon,9'10 have shown that simultaneous ultrasonic measurements can be successfully made during a tensile test. Furthermore, many investigators11-13 have repeatedly proposed in the past several decades that deformation and fracture are mutually self-exclusive, and that the ability or inability of a material to deform plastically, i.e., to generate dislocations, is a major factor in determining whether the material will be ductile or brittle. Thus, in the present work an attempt was made to determine whether the generation and extension of dislocations is inherently more difficult in a brittle steel than in a ductile steel. This article is principally concerned with the study of the relation between the propagation of ultrasonic waves and tensile deformation in a steel series which displays quite different toughness at room tempera-turk. changes in attenuation and velocity of ultrasonic waves have been measured as a function of strain during the deformation process. The results have been interpreted in terms of the vibrating string model for dislocation damping as developed by Granato and Lucke, and it has been found that some of the more subtle predications of the model are in good agreement with the experiments. This would be especially meaningful because most of the previous experiments in testfying the model were carried out with single crystals of high-purity materials and little work has been done with polycrystalline steel alloys. EXPERIMENTAL PROCEDURES AND RESULTS Specimen Materials. The tensile specimens used throughout this experiment were of two compositions selected from a series of Fe-Mo-0.77 pct Mn-0.22 pct C steels prepared for a ductile-brittle fracture transition study. One steel contains 0.21 pct Mo and the other 1.03 pct Mo. These two compositions were chosen for the present study because they possess quite different toughness properties at room temperature. The 0.21 pct Mo steel is quite ductile while the 1.03 pct Mo steel is rather brittle, as measured by the standard Charpy impact test. The alloys had been prepared by vacuum induction melting and chill casting in steel molds. The ingots were hammer-forged into 1/2-in.-sq bars from which tensile specimen blanks were cut. These blanks were first normalized under argon atmosphere at 1700°F and then reaus-tenitized and isothermally transformed at 1050°F to a bainitic microstructure. The chemical compositions, heat treatments, hardness measurements, and Charpy transition temperatures of the two steels are listed in Table I.

Jan 1, 1970

By G. D. Rieck, M. M. P. Janssen

Chemical diffusion coefficients and heats of activation for diffusion in the NizAh fy), NiAl (6), and Ni3A1 (E) intermetallic phases and the solid solution of aluminum in nickel (( phase) were calculated from layer growth experiments. No finite diffusion coefficient for the NiAl3 ((3) inter metallic phase could be calculated. The values of the diffusion coefficients are dependent both on the method of calculation and the type of diffusion couple. The heat of activation for diffusion in the y phase was found to be 47 kcal per mole in the temperature range oj 428" to 610°C. Heats of activation of 41, 12, and 48 kcal per mole were found for diffusion in the 6, E, and ( phases, respectively , in the temperature range of 655" to 1000°C. Experiments with markers in the diffusion zone demonstrate a very pronounced Kirkendall effect. It appears that only aluminum atoms take an active part in the diffusion process during the formation of the 0 and y phases at temperatures of about 600°C. During the formation of the 6, E, and < phases at higher temperatures only nickel atoms are moving. It is suggested that the great stability of the intermetallic compounds in the Ni-A1 system governs the Kirkendall effect. SOME factors controlling layer growth during inter-diffusion in the Ni-A1 system (phase diagram, see Fig. 1) were studied by Castleman and Seig1e.l&apos;~ They found the NiA1, ((3) and NiAl3 (y) intermetallic compounds to appear in the diffusion zone of Ni-A1 couples at annealing temperatures of 400" to 625°C; the NiAl (6) and Ni3A1 (E) intermetallic compounds appeared in y-Ni couples at annealing temperatures of 800" to 1050°C. These authors carefully examined metallographically Ni-A1 couples after 340 hr annealing at 600°C. Besides the (3 and y phases they found very thin layers of the 6 and E phases. ~n~erman~ and Castleman and Froot4 observed a much more rapid growth of the 5 and E phases at 600°C in Ni-A1 couples in case a crack was present at the /3-A1 interface. Numerous layer thickness measurements carried out by Castleman and Seigle on the y phase prove that the layer growth of this phase obeys the parabolic law after a certain transient period. From this they concluded that the layer growth of the y phase is controlled by volume diffusion. The growth of the 13, 6, and E phases appeared to be volume-diffusion-controlled also. The authors estimated that at 600°C and at atmospheric pressure Dp was 1.8 x lo-"ll sq cm per sec, D, 9.1 x 10" ™ sq cm per sec, Qp 27 kcal per mole, and Qy 31 kcal per mole. The present work was carried out to obtain more quantitative data about the kinetics of growth of the phases of the Ni-A1 system and the reactions that occur during the formation of these phases. Because in this system the diffusion process results in the formation of several distinct intermetallic compounds, the current term reaction diffusion is used in the title of this paper. In order to obtain layers of the fl phase compound of uniform thickness, a new technique for preparing diffusion couples was developed. The kinetics of growth of the y phase in 6-Al, E-Al, and Ni-A1 diffusion couples was studied at different temperatures. The kinetics of growth of the 6, c, and ( phases in Ni-y, Ni-6, and Ni-c diffusion couples was also studied at different temperatures. The calculation of the diffusion coefficients Dp and Dy by Castleman and Seigle are critically considered in this paper; by means of a revised method of calculation more reliable val-ues of , and Dg were found. These values are in good agreement with the values of the diffusion coefficients obtained by the method of Boltzmann-Matano. From the temperature dependence of the diffusion coefficients the heats of activation for diffusion were calculated by means of an Arrhenius-type equation. The investigation of the Kirkendall effect has been used to obtain information about the ratio of the intrinsic diffusion coefficients of the separate atoms5 and the mechanism of diffusion. Moreover porosity as a result of a distinct Kirkendall effect would be of practical importance in connection with the bonding of diffusion coatings. The analyses of the diffusion couples were carried out by metallographic methods. The values of the concentrations at the phase boundaries and the concentration profile in each of the phases, which are needed for the calculation of diffusion coefficients, were obtained by electron-pro be X-ray microanalysis. EXPERIMENTAL PROCEDURE A) Materials for Diffusion Couples. The intermetallic compounds 6 (50 at. pct Ni) and E (74 at. pct Ni) were prepared from the pure metals by high-frequency induction melting in argon atmosphere. Use was made of aluminum wire (99.99 wt pct Al) and nickel sheet (99.95 wt pct Ni). The 6 and E phase melts and the nickel shiet (thickness 0.1 and 0.5 mm) used for preparing diffusion couples were annealed for 64 hr at 1200°~ for homogenization and grain coarsening (final crystal size 1 to 3 mm). composition and homogeneity of the intermetallic compounds were checked by mi-crohardness measurements and X-ray diffraction. From the 6 and E phase melts discs of 0.5 mm thickness were prepared by means of a water-cooled rotat-

Jan 1, 1968

By S. E. Maddigan, B. R. Zimmerman

MOST metallurgists are well acquainted with the contributions already made to the study of metals by the use of X-rays. On the one hand, the radiographic method is constantly becoming of increasing importance as a nondestructive testing procedure, while on the other hand the techniques of X-ray diffraction have made important contributions to the fundamental knowledge in studies of phase diagrams, stress conditions, and so forth. A new, and as yet insufficiently appreciated, tool exists in the application of microradiography to alloy structures. This new procedure in many respects bears the same relation to microscopic investigations as does macroradiography to the gross inspection of metal surfaces. However, where macroradiography is used ordinarily to determine the soundness of the metal structure, this new method can be used not only to investigate micro-un- soundness but also to examine the distribution of the constituent elements within the body of the alloy. When combined with the normal procedures of microscopy, this offers a technique of great potential power. It is intended in this paper to present a few examples of copper-base alloys demonstrating the additional knowledge obtained from the microradiographic method. The detection of micro-unsoundness is demonstrated in Fig. I for a sample of cartridge brass as cast. The sample appeared quite sound when viewed under the microscope, but in reality was thickly populated with minute voids. This application of the method conforms in every way to conventional large-scale radiography and has been used to facilitate production of metal for vacuum devices.1 The true potentialities of the technique are revealed by comparison of Figs. 2a and 2b, which are respectively a- photomicrograph and a microradiograph of a cast alloy of 80 per cent Cu, 10 per cent Sn, 10 per cent Pb. The remarkable difference in appearance provides some concept of the additional information this method may yield when used as an auxiliary to microscopic studies. Fig. 2a shows two phases of the copper-tin system plus segregated lumps of lead. Fig. ab, on the other hand, shows with exceptional clarity the actual dendritic growth of the metal crystals during solidification. The practice of microradiography depends upon the laws of absorption of X-rays. This subject has been reviewed thoroughly in recent papers,1,2,3 but for clarity a brief resumé will be given here. As is well known, when a beam of X-rays is transmitted through a layer of homogeneous material, the intensity is reduced according to the relation:[ ]

Jan 1, 1944

By W. H. Bassett

IN the commercial annealing of cartridge brass there are four points regarding which definite data are essential. They have to do with the correct interpretation of grain count in its relation to annealing temperature and, incidentally, to Brinell hardness. These points are: 1. The comparison of the grain sizes of two cartridge-brass mixtures: 69 copper, 31 zinc, 0.376-in. (9.5-min.) gage; and 68 copper, 32 zinc, 0.130-in. (3.3-mm.) gage. 2. The comparison of the grain sizes of annealed metal that had previously been reduced by rolling varying amounts; for instance, 20.0, 36.6, 50.9, and 59.1 per cent. 3. The determination of standards for grain sizes on annealed brass of the composition 68 per cent. copper, 32 per cent. zinc and 69 copper, 31 zinc. 4. The comparison of grain size with Brinell hardness on identical samples of annealed metal. In their comprehensive and thorough investigation of the recrystallization of cold-worked alpha brass on annealing,1 Mathewson and Phillips have discussed the relations between temperature of anneal, degree of deformation, and structural alteration in alpha brass. They have also shown certain comparisons between the ordinary physical properties and the grain size of annealed brass. The purpose of the present investigation is mainly concerned with the grain size of cartridge brass, its relation to Brinell hardness, and the publication of sufficient data to enable those engaged in the inspection of such material to have a correct foundation upon which to work.

Jan 1, 1919

Jan 1, 1924

By D. B. RUSHNORE, K. A. Pauly

Bulletin. No. 41, May, 1910. ABBREVIATIONS. Al. Ind., Alliance Industrielle. am. Elec&apos;n.. . American. Electrician. Bull. As. Ing. El Bulletin de la Association des Ingenieurs Elec- triciens Sortis de l&apos;Institut Electrotecnique Montefiore, Liege. Bull. Soc. bilge d&apos; Elec Bulletin de la Société belge d&apos;Electriciens. Col. Guard . Colliery Guardian. D. R. P Deutsches Reich Patent. Dilly. Poly Dingler&apos;s Polytechniches Journal. E. & DI. Jour Engineering and Mining Journal. E. T. Zeit Electrotechnische Zeitschrift. Eclair. Elec. L&apos;Eclairage Electrique. 1;1. Anzeig Elektrotechnischer Anzeiger

Sep 1, 1910

By C. W. Mac Gregor

THE observation of the flow layers, or Lüders&apos; lines, produced in mild steel when it is stressed into the plastic range often provides considerable useful information for the study of the fundamental mechanism of the plastic flow of metals. It is a commonly observed fact that as soon as the upper yield point in a mild-steel tensile-test bar is passed, an unstable condition develops in the equilibrium, the load on the testing machine drops and if the surface of the bar contains a thin mill scale, the latter begins to flake off rapidly along certain "lines." These "lines" are in reality traces of wedge-shaped flow layers and are formed by the inter-section of the surface of slip and the surface of the bar. The average plane of the flow layer itself has been found by various experimenters to be inclined at an angle of approximately 45° to the axis of the test bar. In a body of more complicated shape stressed in possibly a different man-ner, the tangent planes to the flow layers at each point seem in many cases to be inclined at an angle of about 45° to the principal stress direc-tions at the point. As examples of the latter may be mentioned the logarithmic spirals obtained when a rivet hole is punched in a thin plate, or when a tube is subjected to internal pressure only, etc. Hence if the flow layers are observed from an experiment, the shear-stress trajectories -the curves indicating the directions of the maximum shearing stresses in the body-may then be approximately determined. Since the latter curves form an orthogonal family inclined at 45° to the principal stress directions, the family of principal stress trajectories, or the curves indicating the directions of the principal stresses, may be thus obtained. Although the discontinuous type of stress-strain curve at the yield point is commonly observed with test bars of mild steel, more recent experiments have shown that this is by no means the only material exhibiting this interesting phenomenon. A similar effect has been observed in zinc and cadmium crystals, by Becker and Orowan;1 in duralumin and other aluminum alloys, by Dawidenkow, 2 Elm,3 Fell, and others; in copper and copper-alloy crystals, by Elam;5 in brass by Elam;5 in brass wires by deForest;6 in higher carbon steels, tungsten steel, and

Jan 1, 1939

By B. J. Nelson, F. N. Rhines

STUDIES upon the rates of oxidation of copper alloys containing small quantities of the alloying elements1,2 have shown that steady growth of the scales at predictable rates is limited to a small concentration range and to oxidation at temperatures above 600°C., whereas unpredictable and erratic oxidation behavior is found with alloys containing larger additions and at low temperatures. The latter effect appears to be associated with modifications in the geometric disposition of the oxides. At low concentrations of the baser elements in copper, simple oxidation behavior is typified by the formation of two distinguishable zones of oxidation: (I) a subscale composed of discrete particles of the oxide of the alloying element embedded in substantially pure copper and (2) an external scale composed of cuprous oxide (with a very thin layer of cupric oxide on the outside) containing particles of the alloying element distributed throughout, but most profusely near the metal surface (see Fig. 2, top). Where this type of structure obtains, the process of oxidation is believed to be implemented, and its rate controlled, by the diffusion of the reacting elements through a continuous matrix phase; i.e., through the alpha (metallic) phase of the subscale and alloy and through the cuprous oxide of the external scale. Copper is presumed to diffuse outward through the cuprous oxide layer to the external surface, where it reacts to form more cuprous oxide.3 Oxygen is delivered at the oxide-metal interface in the form of cuprous oxide, which may either react with the alloying element to deposit its oxide at the inner limit of the external scale or dissolve in the pure copper that remains at the surface of the metal after it has become depleted in its oxidizable alloying elements. Both oxygen and the alloying element are pictured as diffusing through the metal, the oxygen inward and the alloying element outward. Where the two diffusion streams meet in sufficient concentration the oxide of the subscale is precipitated. If the oxygen pressure in the surrounding atmosphere is just below the pressure necessary to maintain cuprous oxide undecomposed, no cuprous oxide can be formed; some of the oxide of the alloying element may form upon the external surface, if its decomposition pressure is below that of cuprous oxide, and a normal subscale appears. This condition is achieved experimentally by enclosing the alloy with cuprous oxide together with an excess of copper. At temperatures from 600°C. downward, the precipitated oxides tend to accumulate at the grain boundaries of the alpha phase where a relatively small volume establishes a partial barrier to the diffusion of the reactants. Similarly, large volumes of precipitated oxides tend to deposit in various patterns that may be expected to hinder the diffusion of the reactants in the

Jan 1, 1943

By C. G. Fink

A PART of the material incorporated in this paper was derived from the studies of ancient metal objects examined by the authors from time to time for museums and art collectors, in order to verify the age of the metal. The data thus accumulated during several years was systemati-cally surveyed and supplemented by new material obtained in examina-tion of a set of specimens that had been generously donated by museums for this study. Altogether over 50 pieces were investigated, of which 39 are described in this paper. Every specimen is at least 1500 years old; the majority are considerably older. Exact conditions to which these bronze and copper objects were exposed in the ground are unknown, unfortunately. It is safe to say that exposure conditions did not remain constant during the long period of burial. Although the collection of specimens studied represents a wide variation in location and age of the objects, as well as in the nature and composition of the metals, it was found that fundamental changes in the metal structure were limited to a few typical cases which, therefore, may be considered as more or less general. The authors&apos; chief aim was to study the effect of extremely long periods of corrosion on the structure of bronze and copper.

Jan 1, 1936

By V. K. Zworykin

THROUGHOUT its development the science of electronics, like so many other branches of science and industry, has been indebted to the metallurgist. Metallurgy has provided the electronic engineer with conductors to furnish passage for electrons outside of his evacuated tubes; it. has developed tungsten and tantalum wires for emitting electrons and various refractory metals and alloys suitable for electrodes, which receive them and guide them within the tube; finally it has supplied ferromagnetic materials both for shielding electronic apparatus from magnetic fields and for forming highly concentrated fields for deflecting electron. Today I hope to show you that electronics has not been altogether ungrateful for this assistance. More specifically, I wish to indicate the role that the electron microscope, a relatively recent product of electronic engineering, is playing even now in the study of metals. Every metallographer is aware that his standard metallographic microscope will not reveal structural detail involving separations smaller than a certain fraction -at the very best a tenth-of a micron. This limitation, he knows, is a consequence of the fact that the wave length of the light employed for forming the image has this order of magnitude. It is, thus, a property of all light microscopes. To overcome it, a different imaging medium must be employed. The existence of such a medium became evident when it was demonstrated&apos; that electrons were affected by magnetic and electric fields with axial symmetry just as light is affected by glass lenses; these fields, or "electron lenses," curve the paths of electrons reaching them from an object located on the axis of symmetry in such a manner that, at some distance from the lens, a true image of the object is formed, Furthermore, although a wave length may be ascribed to electron beams just as to light, this wave length is much shorter; for electrons accelerated through a difference of potential of 50 kilovolts, the wave length is only about one hundred-thousandth of that of visible light: Actually, the electron microscopes of today resolve detail one hundred times as fine as light microscopes; if the maximum useful magnification of the latter is taken to be a thousand, that of electron microscopes is one hundred thousand. THE ELECTRON MICROSCOPE Since electrons are readily scattered and absorbed by matter, specimens to be observed in the electron microscope must be prepared to be less than one micron in thickness; furthermore, the interior of the microscope must be carefully evacuated to prevent collisions between the electrons and the molecules of the air. Finally, as already mentioned, the lenses are formed by axially symmetric magnetic or electric fields, and not by any material medium such as glass. Apart from these basic differences, the electron microscope closely

Jan 1, 1943