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By E. H. Rennhack

THE presence of sufficient quantities of the inter- stitial elements, carbon, oxygen, and nitrogen is known to decrease sienificantly the room-temperature impact resistance of titanium. Hydrogen does not appear to affect impact strength unless present

Jan 1, 1959

By M. Metzger, T. A. Read

The strengthening effects of hydroxide and synthetic plastic films in the creep of cadmium crystals were studied. The results were broadly consistent with the naive mechanical model. The dislocation-film reactions and the significance of the macroscopic stresses associated with certain films are discussed in the paper. ALTHOUGH a number of studies have been made of the strengthening of metal single crystals produced by surface films usually less than Ir in thickness, the mechanisms underlying this phenomenon are still not well established. Roscoe' first noted the critical shear stress and the stress-strain curve of cadmium-crystal wires to be raised markedly by the presence of oxide films. Subsequent investigators2-' have confirmed the effect for several metals and films including a number of plated poly-crystalline metal films. Harper and Cottrel1" with a thin film on zinc presumed to have formed as a result of etching and washing treatments) did not detect a strengthening at a rate of extension of 10 ' sec-' although at lo-' sec-' the film raised the flow stress by 7 pct at 0.0005 extension and by 22 pct at 0.01. The marked strengthening at moderate strains thus appeared to be the result of a dynamical effect dependent on the rate and amount of plastic flow. Other observations'. " are consistent with this view. The large effects of films on the tensile creep rate of single crystals were demonstrated when the observation of Andrade and Randall," that the creep rate of cadmium crystals increased on immersion in aaueous solutions of certain cadmium salts. was shown' * to hold only when a film was present initially and was explained by the attack of the solution on the film and the removal of its strengthening effect. Changes in the creep rate interpretable in terms of the removal or formation of various films have been observed by other investigators.3,8,10,31 Certain other experiments have been reported in which surface films or liquid environments have affected the plastic behavior of metals, but these are not within the scope of the present work which deals with the effect of solid films on single-crystal slip. Possible Mechanisms—The film does not significantly reduce the load on the crystal in the elastic range (the elastic moduli of film and crystal are not too greatly different and the area of the film is so small that its share of the load is negligible) so that the initial operation of internal Frank-Read sources is not affected. However, as Fisher" has pointed out, a source terminating on the crystal surface requires only half the critical stress for activation as an internal source of the same length and a relatively strong film (even a very thin one) would pin the free ends of the surface sources so that, if surface sources are responsible for plastic flow in a clean crystal. such a film would increase by roughly a factor of 2 the stress necessary for generation of dislocation loops at a given rate. The other general possibility is that the film resists the passage to the crystal surface of internally generated dislocations. If the film is adherent and if the specimen is to exhibit slip of the same general character as a film-free specimen, the film must deform or fail in shear at the active slip planes. The naive mechanical model, where shear-stress relaxation across the active slip plane in the metal transfers load to the film until it deforms and slip can

Jan 1, 1959

By R. C. Hall

In the present work, experiments determined the effect of the order-disorder reaction near the Fe,Al composition on the magnetic anisotropy and magnetostriction of single crystals: A material which is more difficult to magnetize in one direction than in another is anisotropic, that is, the ease with which ferromagnetic materials may be magnetized in different crystallographic directions determines the anisotropy. Two investigators have estimated the anisotropy of the aluminum iron alloys from poly-crystalline data;1,2 however, these estimates should not be considered to be reliable. Siege1 made two accurate measurements of the first anisotropy constant* from single crystals of 4.0 wt pct aluminum sotropy throughout the ferromagnetic aluminum iron system have been described. These results were obtained from single crystals which were slowly cooled and, hence, ordered near the Fe3Al composition.4 Magnetostriction is the change in length per unit length resulting from applying a magnetic field to a ferromagnetic material. For the purpose of this report, the saturation magnetostriction was measured. This is the maximum change in length per unit length resulting from an applied magnetic field large enough to saturate the material in a given direction. This magnetostriction has been determined for single crystals of the ferromagnetic aluminum-iron system.' The crystals near Fe3Al were also ordered. PREPARATION AND TREATMENT OF SINGLE CRYSTALS Single crystal bars were grown from the melt.' The raw materials for these crystals were iron and aluminum of 99.9 and 99.99 pct purity, respectively. From the bars, crystals in the shape of circular

Jan 1, 1959

By M. L. Bleiberg, L. J. Jones

THE changes induced in metals by subjecting them to heavy particle bombardment have received widespread attention in recent years and results of such investigations have been reviewed by Slater, Dienes, Seitz, Siegel, and Seitz and Koehler.1-5 Of particular significance are the property changes that occur in uranium-base alloys due to neutron bombardment. These changes are generally very severe since fission occurs within the sample, resulting in the formation of heavy, highly energetic fission fragments, as well as additional neutrons born in fission, which serve to intensify the damage generally associated with fast neutron bombardment of nonfissionable metals. The property changes in uranium and uranium alloys due to neutron bombardment have been summarized by Paine and Kitte16 and Billington.7 It has also been shown that neutron bombardment of

Jan 1, 1959

By R. H. Hiltz, N. J. Grant

Solid-solution titanium-cerium alloys, cold-rolled to 0.010 in. thickness, and annealed, were internally oxidized to produce a fine dispersion of CeO2 in the titanium matrix. The oxidized alloys were cut into sheet stress-rupture specimens and tested for rupture life at 426°C, 538°C, and 648°C. There was noted an improvement in strength over that of pure titanium, the internally-oxidized alloys showing curves with a much flatter slope on the log stress—log rupture life plot, characteristic of the stable metal-metal oxide class of alloy. DURING the past 3 to 4-year period, considerable interest has been aroused by the potential offered by alloys based on the ultra fine dispersion of a hard insoluble phase in pure metals or solid-solution alloys for use at high temperatures. The impetus to the development and study of such alloys was provided by Irmann.1,2 through the success of the SAP alloy (pure aluminum containing about 12-14 volume pct of A12O3). Subsequent studies have thrown light on the true potential of this class of alloys in the remarkable stability of the structure after long-time exposure at high stresses and high temperatures.a-5 While success has been demonstrated in the special case of the A1-Al2O3 alloys, and more recently

Jan 1, 1959

By J. O. Betterton, W. K. Noyce

I~JTI HE work on the indium-zirconium system is part A of a larger investigation of zirconium-phase diagrams with solute elements, silver, cadmium, indium, tin, and antimony, which is intended to provide information about the metallic valency of zirconium. The system is an ideal example since the sizes of indium and zirconium atoms are nearly the same, and the solute and the solvent occur in the same row of the periodic table. The only solute from Group III B which had previously been added to zirconium was aluminum' and it was of interest to determine whether indium would dissolve appreciably in alpha zirconium and stabilize this phase in the same manner. Experimental Procedure The experimental methods used in the indium-zirconium work are similar to those described earlier for the silver-zirconium system.' Reactor Grade 1 iodide zirconium of 99.95 pct purity* and indium of 99.99 pct purity were arc cast, and then annealed in silica capsules with protective foils of zirconium and molybdenum. Annealing temperatures and times varied from 1 day at 1300°C to 43 days at 550°C. The compositions of the alloys were detcrmined by an individual chemical analysis of the actual metallographic specimens after two or more sides were examined to insure the absence of solute segregation. Calibrated thermocouples were used throughout, and the uncertainty in the temperatures reported, owing to calibration error, furnace control, and measurement of the small temperature gradients near the capsules, is believed to be well within ±2°C. Typical impurities in the alloys were found by a neutron-activation analysis of four different alloys. The results are given in Table I where the ranges of values will be seen to compare favorably with the usual values for iodide zirconium. With respect to oxygen, nitrogen, and carbon the absence of any unusual amount of these elements compared to iodide zirconium was indicated by the alloy hardness values, since hardness is increased sharply by the presence of these elements. The value for 0 pct In, obtained from the alloys by extrapolation, is 58 2 7 VPN which is in the low part3 f the normal range of hardness found in iodide zirconium. Results of the Experiments The zirconium-rich portion of the indium-zirconium phase diagram from the annealing experiments is shown in Fig. 1. Indium was found to

Jan 1, 1959

By C. Feng, C. Elbaum

AN investigation of the fracture plane of titanium, containing 0.5 wt pct to 1 wt pct 0, was recently carried out by fracturing large-grain specimens in a tensile jig at room and liquid air temperature. The experiments were carried out on strip specimens 3 to 5 mm wide, 2 mm thick and about 100 mm long. The specimens consisted of several grains, each occupying the entire cross section of the strip and extending over a length varying from 5 to 20 mm. These specimens were prepared by a method described elsewhere.' The identification of the fracture plane was possible in twelve of approximately thirty crystals investigated. In some cases severe local distortion took place near the fracture and prevented identification of the fracture plane. On, at least, three occasions the fracture propagated in a helicoidal

Jan 1, 1959

By E. Buehler

THE floating zone refining technique'-4 has provided a new experimental approach to the purification of many reactive materials especially those with high melting points. The method avoids the use of a crucible and at the same time it provides a favorable environment at the freezing interface for the growth of single crystals. The physical properties of refractory metals such as molybdenum and tungsten have long been of interest from a fundamental point of view and because of their widespread use in vacuum electronics. Single crystals are often desirable for study. Prior to the introduction of zone melting, single crystals of these metals were usually grown in the solid state at temperatures slightly below the melting point by strain annealing3 or by the application of a moving temperature gradient. Such techniques were mostly successful on wires, ribbons, or rods having small cross-sectional areas. Recently, larger samples of tungsten, molydbenum, tantalum, columbium, and rhenium were melted by a floating zone process which employed electron bombardment heating.10 Single crystals were obtained in many instances

Jan 1, 1959

By D. J. Mack, J. A. Mullendore

MUCH work has been reported in the literature on the separation of individual phases in steels, particularly carbides, for the purpose of structural and chemical analysis. The methods usually involve separation of the particles from the matrix by chemical or electrochemical methods. Another .group of alloys ideally suited for selective removal of one phase is the commercial aluminum bronzes. Al-

Jan 1, 1959

By J. W. Downey, A. E. Dwight, R. A. Conner

THE MgCu2 type Laves phase has been observed in the thorium-osmium and thorium-iridium systems. Two alloys corresponding to ThOs2 and ThIr, were prepared by arc melting a 5-g charge in a water-cooled copper. crucible. Both alloys were very brittle. X-ray diffraction patterns were obtained using filtered Cu and Co radiation with a Straumanis-type DeBye-Scherrer camera. The a,, values as shown in Table I were obtained by a Nelson-Riley extrapolation. The structure belongs to space group Oh7-Fd3m. Relative intensities for each of these compounds were calculated on the assumption that 8 thorium atoms occupy (a) positions and 16 osmium or irid-ium atoms occupy (d) positions. The calculated intensities are the product of the Lorentz-polarization factor, the multiplicity factor and F', the structure factor, neglecting the absorption and temperature

Jan 1, 1959

By J. Gurland

A method of measuring the degree of contact between grains of 1 phase in multiphase alloys is derived. It is shown that the number and the areas of contacts between grains can be determined by metallographic analysis. Relations between the area of contact, the dihedral angle, and the mean free path of spherical grains are developed. It is often desirable to interpret the mechanical behavior of alloys in terms of structural features observed under the microscope. A quantitative description of microstructures of multi-phase alloys requires consideration of geometrical and topological factors such as the amount, size, and shape of grains and the continuity of the constituent phases. The number and dimensions of grains may be metallographically determined by applying the methods described by Smith and Guttman1 and by Fullman2 in their papers on the measurement of particle size and internal boundaries in opaque bodies. The present work intends to show that the same methods can be used to evaluate some of the parameters of phase continuity; namely, contiguity, number of contacts, and areas of contact of grains of 1 phase. Application of these concepts may be found in such problems as the calculation of electrical conductivity of 2-phase alloys, or the evaluation of the length of fracture path through grains of a brittle phase more or less dispersed in a matrix. I. DEFINITION The following definitions and derivations are based on 2-phase alloys, of which the units of the microstructure are grains. Contiguity is defined as the average fraction of surface area shared by 1 grain of a phase with all neighboring grains of the same phase. In a 2-phase alloy of constituents A and B, the contiguity of phase B is given by: where Cb is the contiguity of phase B aBß is the average area of contact between a grain of phase B and adjacent grains of phase B, aAB is the average area of contact between a grain of phase B and adjacent grains of phase A. The contiguity is a measure of the average degree of contact between grains of phase B in the 2-phase mixture. The ratio varies from 0 to 1 as the distribution of phase B changes from a completely dispersed to a fully agglomerated structure. Smith and Guttman1 have shown that the relative areas of 2 kinds of interfaces in a volume are directly proportional to the number of interfaces of each type intersected by a random line. The lines can be drawn on metallographically prepared surfaces as long as these are random sections through the sample. The contiguity ratio of phase B then becomes: Cb = 2NBB/NAB + 2N BB [2] where: NBB is the number of interfaces between grains of phase B intersected per unit length of random line. Nab is the number of interfaces between grains of phase B and grains of phase A intersected per unit length of random line. The factor 2 arises in Eq. [2] because each interface B-B is associated with 2 grains of B-phase, whereas the interface A-B involves only 1 grain of B-phase. The calculations of contiguity do not require any assumption as to volume fraction, particle size, or particle shape. Fullman2 gives an expression for the determination of the mean free path between particles. The mean free path PB between grains of phase A, i.e., through grains or agglomerates of phase B, may be found from the relationship: PB =2fB/NAB [3] where fB is the volume fraction of phase B. The calculation of the mean free path is independent of grain shape and grain size distribution. The calculation of the number of contacts and of contact area requires assumptions of grain shape, grain size, and contact shape. In the case of uni-

Jan 1, 1959

By T. F. Archbold, R. E. Grace

The use of interference colors is the simplest experimental way to determine the approximate film thicknesses of oxides formed on liquid metals. A large amount of data under various conditions can be obtained easily and quickly. The oxidation of liquid lead has been studied by several investigators,1-5 but all of these studies have been on thick oxide scales. The purpose of- this work was to investigate the range 0 to 2000A film thickness. By using electron diffraction, Jenkins concluded that the surface of the film on molten lead is yellow orthorhombic PbO with the (110) planes parallel to the surface. This was confirmed by Shimaoka and Yamai7 who also used electron diffraction from the melting point of lead to 730°C. The rate of oxidation of pure liquid lead in air was measured in the temperature range 453" to 643°C by observing the rate of formation of interference colors on static lead baths. Approximately 170 g of lead (99.999 pct) was placed in a refractory dish having a working surface of 13.8 sq cm. This crucible was placed in a deep-well resistance furnace open to the atmosphere. The times at which the various interference colors appeared, after wiping the molten metal surface clean, were determined at 453", 533°, 600°, and 643°C. Elapsed times were determined with a stopwatch and all colors were estimated by eye. Bath temperatures

Jan 1, 1959

By E. S. Bartlett, D. N. Williams

QUANTITATIVE values for the oxidation rate of unalloyed molybdenum in air at temperatures above the melting point (1460°F) of the characteristic oxide are contained in the literature as a result of previous investigations. Lustman' reported values corresponding to 0.36 in. penetration per day (IPD) at 1500 and 1600°F in still air, noting essentially no variation in rate with temperature. Jones, Spretnak, and Speiser' reported values corresponding to 0.14 and 0.13 IPD at 1500 and 1800°F, respectively, in still air, attributing the decreased oxidation rate at higher temperatures to a lesser accumulation of the corrosive molten oxide on the surface at the higher temperature as a result of increased volatilization rate. Harwood3 ecently summarized work in the field, presenting generalized data corresponding to 0.48 to 0.96 IPD at 1800°F and 0.55 to 0.83 IPD at 1700°F in slowly flowing air. In a recent program at Battelle, it became desirable to know more about the characteristic oxidation behavior of molybdenum under varying conditions of temperature and atmosphere. Using oxidation-test apparatus designed for dynamic, continuous recording of weight change during testing,' values for the oxidation rate of molybdenum were obtained at temperatures from 1400 to 2150°F. In addition the effect of air flow on the oxidation rate was studied briefly at temperatures of 1600, 1800, and 2000°F. Exhaust of the contaminated atmosphere from the oxidation chamber was effected by an impeller pump attached to a 3/16-in.-diam opening in the oxidation chamber. The volumetric exhaust rate (cubic feet per hour) was normally maintained slightly in excess of the input rate to avoid condensation of MOO,,' on the sample suspension rod. The entering atmosphere was preheated prior to admission to the oxidation chamber by a 1 1/2-in.-diam cup packed with shredded asbestos. The experimental data are presented in Table I. Comparing conditions 2 and 3 (taking into account the temperature difference) and conditions 8 and 9. shows that in the absence of forced exhaust an atmosphere of moving air results in greater oxidation rate than a stagnant atmosphere. The use of forced exhaust, as shown by comparing conditions 3 and 4 and conditions 14 and 15: resulted in an even greater increase in oxidation rate. By virtue of the size of the atmosphere input and exhaust openings, it was calculated that the exhaust velocity was about 60 times that of the input velocity for essentially equal volumetric flow rates. Because of the proximity (about 3/4 in.) of the exhaust port to the specimen, it is logical to assume that cleansing of the atmosphere immediately surrounding the specimen was accomplished much more efficiently by the exhaust flow than by the input flow at a constant-volumetric flow rate. Also, it can be seen by comparing conditions 4 through 7 and conditions 11 through 13 that increasing the rate of atmosphere flow (by increasing input velocity with a proportional increase in exhaust velocity) above some optimum value has little, if any, further effect on the oxidation rate. These results suggest that there is a maximum oxidation rate for molydenum at a given temperature which is obtained when conditions are maintained such that the partial pressure of MOO3 in the atmosphere surrounding the specimen is at a low value. By controlling the partial pressure of MOO, surrounding the sample, it is possible to control the rate of volatilization of MOO:, from the surface. This, in turn, affects the rate of oxidation, since the thickness of the MOO3 layer determines the amount of oxygen which will be able to reach the reaction surface.' When the liquid oxide layer is less than some critical thickness, i.e., when the volatilization rate is high, enough oxygen is transported to the active surface to permit oxidation to proceed at the maximum rate allowed by the kinetics of the oxidation reaction. However, if the volatilization of MOO,, is suppressed, the thickness of the layer of MOO3 on the surface increases, and the diffusion of oxygen through the oxide layer becomes the rate-controlling step in the oxidation process. The lack of agreement between the present results and those of previous investigators is presumed to be due to differences in removal of the oxidation product (Moo,) from the immediate vicinity of the sample. By comparing conditions 1, 5, 10, 12 and 14, it is seen that when forced exhaust was used the oxidation rate of molybdenum increased with increasing temperature. A rapid increase was observed between 1400 and 1600°F, attributable to the effects

Jan 1, 1959

By G. A. Alers, J. H. Bechtold, R. W. Armstrong

The plastic flow of .molybdenum was studied in compression at temperatures below the ductile-to-brittle transition in tension. It was found that molybdenum is ductile in compression at all temperatures down to —196°C. The yield strength continues to increase with decreased test temperature below the tensile ductile-to-brittle transition temperature range. Difficulty was encountered in correlating the experimental data obtained with existing theories of plastic flow. THE body-centered-cubic metals as a group show a rapid and very large increase in yield strength with decreasing test temperature.' It is of fundamental importance to the understanding of the mechanism of plastic flow that this general characteristic be examined carefully. Unfortunately, brittle fracture in low-temperature tensile tests prohibits the measurement of flow properties at very low temperatures. By using compression tests, fracture can be suppressed and the plastic flow of the bcc metals can be studied far below their tensile ductile-to-brittle transition temperatures. Molybdenum was chosen for this experiment and plastic-flow data were obtained down to —196°C.

Jan 1, 1959

By A. H. Daane, J. J. Hanak, F. H. Spedding

IN the study of the rare earths at the authors' laboratory, methods have been devised for preparing pure rare earths and many of their properties'- have been determined. Very little information is found in the literature on the properties of europium metal, which is the scarcest of the naturally occurring rare earths. Since we have recently prepared europium in appreciable quantities, we are in a position to describe the method of preparation and some of the properties of this metal. Metallic europium was first prepared by Klemm and Bommera who heated a mixture of potassium and EuC1, in vacuo, and obtained a mixture of europium metal and KC1. Trombe4 electrolyzed a molten salt bath containing NaC1, KC1 and EuC1, and collected the europium in a molten cadmium electrode as a cadmium alloy. The cadmium was subsequently removed by distillation but the properties of the residue indicate that pure europium was not obtained. McCoy attempted to prepare europium by distilling mercury from europium amalgam, but the stability of the Eu-Hg intermetallic compounds prevented a complete separation of the two elements. The preparation of very pure europium was suggested by previous work in this laboratory1,4 in which ytterbium and samarium were prepared by heating the respective sesquioxides with lanthanum metal turnings to 1450°C in vacuo and distilling ytterbium and samarium from the mixtures. Subsequently, small test experiments showed that europium also could be prepared by this method. A more detailed version of this method was used in the work described in this paper. Klemm and Bommer3 found by X-ray analysis that the metal had a body-centered cubic (bcc) structure with a lattice constant a, = 4.582 ± 0.00A (converted from kX units), with a bulk density of 5.30 g per cu cm. They found europium to be weakly ferromagnetic. Spedding, et al,7 found a,, to be 4.606 ± 0.001A. Barrett' showed that the bcc structure persists throughout the temperature range 5 to 3000K with a, at 5, 78, and 300°K being 4.551 ± 0.003, 4.551 ± 0.004 and 4.577 ± 0.001A, respectively. Trombe1 described europium as being iron gray, malleable, about as hard as neodymium and not attacked by cold water. He estimated the melting point to be between 1100 and 1200°C. Previous work in this laboratory indicated that the pure metal corrodes rapidly in air and reacts with cold

Jan 1, 1959

By P. G. Shewmon

THIS paper is concerned with the redistribution of a second phase when a 2-phase solid alloy is annealed in a temperature gradient. An examplewhich is discussed in detail is the redistribution of carbides in a low-carbon steel annealed in a temperature gradient below the A, temperature. As a result of such a treatment, the carbides at the cooler side disappear, while those at the hotter side grow in size. In a 2-phase zirconium-hydrogen alloy, the amount of hydride will tend to be increased at the cold end and decreased at the hot end. That such an effect can occur in solid metals is not generally appreciated. However, with the trend toward higher operating temperatures in several fields, this phenomenon will become more important in the future. Actually, such effects should be rather common in the core of nuclear power reactors where the op- erating temperatures are high and temperature gradients are an order of magnitude higher than in previous high-temperature applications of alloys. The purpose of this paper is to point out the existence of such an effect, to discuss the manner in which one predicts whether a redistribution of the second phase will occur in a given system, whether the second phase will go to the hot or cold face, and how rapidly the redistribution will occur. One aspect of this problem has been discussed by pfam and Wernick2who considered the migration of a small liquid region through a solid. In their case, the diffusivity in the liquid is much larger than in the solid and they have interpreted their results successfully by using Fick's first law as the flux equation in the liquid, and taking the concentration gradient from the phase diagram. For all the cases they have studied, this type of analysis predicts that the liquid region migrates from the cold end to the hot end, and this is what was found ex-

Jan 1, 1959

By C. F. Lucks, W. D. Wood, H. W. Deem

Electrical ad thermal-conductivity measurements have been made on titanium and 6 titanium alloys. The following expression was derived for calculating thermal conductivity from electrical conductivity: K = 0.626 s T x 10-8 + 0.00497 where K = thermal conductivity, cal sec-1 cm-2 cm0C-1 s = electrical conductivity, ohm-]cm-1 T = temperature, deg Kelvin This expression gives calculated thermal conductivities that are not in error by more than 10 pct from the experimental values for A-55 titanium and the 6 alloys measured; namely, A-110AT, C-110M, C-130AM, Ti-6A1-4V, Ti-140A, and Ti-155A, which are representative of the presently available commercial a and a-ß titanium alloys. THERMAL conduction in a gas, liquid, or amorphous solid is accomplished by energy diffusion from molecule to molecule, a relatively slow process. For this reason, such materials have a low thermal conductivity. When a material is crystalline, there is in addition a heat transfer by a vibratory motion of the crystal lattice. It is the contribution of this lattice conduction that makes a crystal a much better conductor than a gas, liquid, or amorphous solid. Lattice conduction, however, does not account for the relatively enormous thermal conductivity of a metal. This is explained by the freedom of movement of valence electrons. These valence electrons drift in a metal in a direction toward a drop in electrical potential and, to a lesser degree, in the direction of a temperature drop in the metal. Since both the electrical and thermal conductivities of a metal depend upon the presence of free electrons, there should be some relatively simple relation between the two. Wiedemann and Franz in 1653 first discovered an approximate empirical relation between the electrical and thermal conductivities of metals at the same temperature. This ratio varies appreciably with temperature, and Lorenz later showed that by adding a temperature factor to the Wiedemann-Franz ratio the value of K/(s T) is constant, where K is the thermal conductivity, a the electrical conductivity, and T the absolute temperature. This Lorenz constant shows a greater variation for alloys than it does for pure metals, the ratio increasing slightly as the proportion of alloying elements increases. It has been found that a closer correlation between electrical conductivity and thermal conductivity can be obtained by means of an equation of the form K - A s T + B [1] where A and B are constants. In this equation the term AsT is the contribution of electronic conduction and B is the contribution of lattice conduction to the total thermal conductivity. EXPERIMENTAL MEASUREMENTS Electrical-resistivity and thermal-conductivity measurements on titanium and 6 titanium alloys serve as the basis for determining an equation of the K = AsT + B type for titanium alloys. Specimens—Table I shows nominal compositions of the 7 metals on which electrical and thermal conductivities were measured. The electrical and thermal-conductivity specimens for each alloy were taken from stock of the same heat and usually from adjacent sections of the same bar. All of the metals measured were in a mill-annealed condition. Methods—Electrical-resistivity measurements

Jan 1, 1959

By J. D. Lubahn

Prior tests by Dorn, where the strain rate in a tensile test was suddenly changed, have shown a small, but definite rate-history effect to exist. If this effect is neglected in the calculation of creep behavior from tensile results, an error should occur. This paper shows that this error is smaller than experimental scattering, and may be neglected. DORN1 has found a small but definite rate-history effect in the aluminum alloy 250 at liquid-nitrogen temperature using tensile tests with an abrupt rate change. Such an effect is to be expected from pre-viously known2-6 temperature-history effects, together with general experience that effects produced by changing the strain rate a large amount are simi-

Jan 1, 1959

By J. R. Mihalisin

THE application of electron-microscopic techniques to the study of high-temperature alloys has met with considerable success.' This is due, in part, to the submicroscopic nature of the phases present in this type of alloy making high resolution methods necessary. Among the newer electron-microscopic techniques, the extraction replica technique' has proven to be of great value. This method allows phases to be examined in situ in the microstructure and permits their identification by means of electron diffracticn. This technique eliminates the main disadvantage in using electrolytic extraction of residues as a means of phase identification; that is, the uncertainty of relating constituents identified in the residue to what is observed microscopically. In the course of work relating microstructure to physical properties in high-temperature alloys, it was found that the formation of grain-boundary carbide phases and intragranular precipitation of an age-hardening phase greatly influence strength and ductility at high temperature. In particular, it was found that the formation of a unique type of precipitate in Nimonic "80" (a nickel-chromium-aluminum-titanium type superalloy), termed cellular or discontinuous precipitate, at grain boundaries generally results in poor ductility at high temperature. However, this mode of precipitation can be avoided by proper heat-treatment. It was assumed that this precipitate was either a carbide or age-hardening type of phase. Using various etching procedures with the extraction-replica technique, it was found to be possible to extract the carbide and age-hardening phases preferentially in this type of alloy. Previous work has shown that the age-hardening phase can be extracted from superalloys merely by

Jan 1, 1959

By J. O. Betterton, D. S. Easton

A detailed investigation was made of the phase diagram of silver-zirconium, particularly in the region 0 to 36 at. pct Ag. The system was found to be characterized by two intermediate phases Zr2Ag and ZrAg and a eutectoid reaction in which the -zirconium solid solution decomposes into a-zirconium and Zr2Ag. It was found that impurities in the range 0.05 pct from the iodide-type zirconium were sufficient to introduce deviations from binary behavior, and that with partial removal of these impurities an increase in the a-phase solid solubility limit from 0.1 to 1.1 at. pct Ag was observed. The phase diagram of the silver-zirconium system is of interest as an example of alloying a transition metal from the left side of the Periodic Table with a Group IB element. Silver would normally act as a univalent metal, its filled 4d-shell remaining undisturbed during the alloying. However, there is a possibility that some of the 4d electrons might transfer to the zirconium. An insight into such a question can occasionally be obtained by comparison of phase diagrams. The silver-zirconium system forms part of a more complete review of various solutes in zirconium in which these valency effects were studied.' Earlier work on the silver-zirconium system was done by Raub and Enge1,2 who investigated the silver-rich alloys. After the start of the present experhents, work on this system was reported by Kemper3 and by Karlsson4 which for the most part agrees with the phase diagram presented here. EXPERIMENTAL PROCEDURE The alloys were prepared by arc casting on a water-cooled, copper hearth with a tungsten electrode and in a pure argon atmosphere. Uniform solute composition was attained by multiple melting on alternate sides of the same ingot. Progressive improvements in the vacuum conditions inside the apparatus during the course of the experiments reduced the Vickers hardness increase of the pure zirconium control ingot from 10 to 20 points, observed initially, to negligible amounts at the end of the experiments. Such hardness changes in zirconium are a well known indication of purity. For example, -01 wt pct additions of oxygen, nitrogen, and carbon increase hardness by 6, 10, and 3 VPN respectively. '9' Further verification that the final casting technique did not add a significant quantity of impurities was obtained when pure zirconium was arc cast and then isothermally annealed in the vicinity of the allotropic transition. The transition was always observed to take place over the same temperature range as in the original crystal bar. The alloy ingots were annealed in sealed silica capsules for times and temperatures which varied between 1 day at 1300°C and 60 days at 700°C. The best method found to prevent the reaction of the zirconium with the silica was foil wrapping of molybdenum or tantalum. With this method, samples of pure zirconium were found to be unchanged in hardness after annealing for 3 days at 1200°C. In most of the experiments the protection of these foils was supplemented by an additional layer of zirconium foil inside the molybdenum or tantalum foil. The alloys, foil, and the capsule were outgassed at pressures in the range 10 to l0-7mm Hg in the temperature range 800" to 1100°C before each anneal in order to remove hydrogen and other impurities, and to provide a suitable container for the high purity, inert atmosphere, which is essential in the annealing of zirconium. The temperature measurements were made with Pt/Pt + 10 pct Rh thermocouples calibrated frequently during the experiments against the melting points of zinc, aluminum, silver, gold, and palladium. For the longer anneals the sum of various temperature errors was generally well within ± 2°C. For short-time anneals and during thermal analysis the overall temperature error is considered to be within ± 0.5°C. The compositions of the alloys from the quenching experiments were determined by chemical analysis at Johnson Matthey and Company, Ltd., under the direction of Mr. F. M. Lever. The actual metallo-graphic samples were individually analyzed in every case, and prior to the analyses two or more sides of each specimen were examined to insure that the specimen was not segregated. The sum of the solute and solvent analyses was in each case within the range 99.9 to 100.1 pct. In the course of the experiments, minor impurities in the range 0 to 500 ppm were found to have significant effects on the zirconium-rich portion of the phase diagram. Similar effects had been encountered previously in other zirconium phase-

Jan 1, 1959