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By J. H. Keeler

The tensile characteristics of Zr-Fe binary alloys containing up to 5 atomic pet Fe are reported for the temperature range —195o to 500°C. A linear relation between stress at constant strain and volume fraction ZrFe2 was found. The coincidence of data from unalloyed zirconium with zero volume fraction ZrFe2 indicated little or no solid solution in IL-zirconium, in agreement with Hayes, Roberson, and O'Brien.' The absolute increase in strength due to particle strengthening, although decreasing with increasing temperature, was a greater percentage at high temperature. ZIRCONIUM is of interest for use in nuclear reactors because of its low neutron absorption, good corrosion resistance, and high melting temperature. The alloying of zirconium with iron has been considered as a means of increasing the mechanical strength of zirconium without seriously increasing the neutron absorption " over that of unalloyed zirconium. The Zr-Fe system has a eutectic at 934°C and 23.7 atomic pet Fe and a eutectoid at 800°C and 4.0 atomic pet Fe. The zirconium-rich end of this system determined by Hayes, Roberson, and O'Brien' is shown in Fig. 1. The intermetallic phase coexisting with a and ß-zirconium has been reported as Zr2Fe3,1 with a,, - 7.04A, and as ZrFe2,2 face-centered-cubic, the phase being of the MgCu, type (C-15). It was more recently agreed to be ZrFe33 and face-centered-cubic. This intermetallic phase is essentially insoluble in a-zirconium, the hexagonal-close-packed modification; and the lack of solubility makes the system particularly interesting metallurgically. The system is strengthened almost exclusively by particles of in-termetallic phase, and there is very little solid solution strengthening to complicate the analysis. The Bureau of Mines, in its survey of zirconium-base alloys,' reported the tensile properties of Zr-Fe alloys made from chloride process zirconium, induction melted in graphite crucibles, and swaged at 850°C in iron sheaths. Additional data on these ZrFe alloys were given by Hayes, Roberson, and O'Brien.' Chubb and Muehlenkamp% have recently reported the strength of Zr-Fe alloys containing up to about 1.25 atomic pet Fe. These alloys were made from crystal bar zirconium, hot worked at 820°C, annealed at 700°C, and tested both at room temperature and at 316°C. Experimental Procedure The materials used in this investigation were high purity zirconium made by the iodide process and electrolytic iron. The composition of the zirconium as obtained by spectrographic analysis and the nominal composition of the iron are given in Table I. Alloy ingots weighing 100 g were made by are-melting under argon in water-cooled copper crucibles in a multihearth furnace."

Jan 1, 1957

By R. W. Guard, W. R. Hibbard

As part of a program to determine the deformation characteristics of pure metals, the tensile creep properties of high purity aluminum (99.994 pct Al) have been determined using a constant stress loading. The effects of prior annealing temperature, stress, and temperature were studied. These data have been examined on the basis of present-day theories. None of the methods was found to give an accurate representation of the data over all conditions studied. The structures of several specimens after creep were examined by optical and X-ray metallography. Marked changes in the grain size were found and it was substantiated that the magnitude of these changes was greatly influenced by the applied stress (or creep strain). The changes in fiber texture which occur on creep deformation are shown to be consistent with those observed on deformation at lower temperatures. UNDERSTANDING the properties of new mate-*-J rials under different experimental conditions must be based on considerable knowledge of the mechanisms involved in plastic flow and how to control them. Tension and creep tests have been carried out on a number of pure metals in order to document their deformation characteristics. It is hoped that a comparison of these data with theory will indicate those areas where work on mechanisms and their relation to metallurgical variables may be most fruitful. Since aluminum has often been used for studying the qualitative aspects of creep behavior, it was chosen for detailed study. The effects of annealing temperature, stress, and test temperature on the constant stress creep behavior were studied. Although most of the tests were made above one half the absolute melting temperature, a few tests were made at lower temperatures to give qualitative comparisons. Some specimens were examined metallo-graphically and by X-ray diffraction techniques after creep. Since only a limited number of specimens could be made from the available material, no effort was made to study detailed changes in sub-grain structure or other metallographic features. Servi and Grant3 nd Dushman et al.' have previously published studies of high temperature creep of high purity aluminum. In both cases, a great deal of emphasis was placed on the minimum creep rate.

Jan 1, 1957

By R. P. Carreker

ONE of a series, this report is concerned with the experimental documentation of the deformation behavior of pure metals over a wide range of temperature. Previous reports in the series describe the creep behavior of platinum1 and aluminum2 and the tensile behavior of copper,&apos; silver,& and molybdenum.5 Subsequent reports will describe the creep behavior of copper and silver and the tensile behavior of platinum, thus providing extensive creep and tensile data on four face-centered-cubic metals tested under comparable conditions. Material Two lots of high-purity aluminum were used in this investigation. One lot, designated AC, was obtained direct from the Aluminum Co. of America as high-purity aluminum in notch-bar form. A typical analysis of this type of aluminum is 0.006 pct Si, 0.015 pct Cu, 0.006 pct Fe, and 99.975 pct Al. Lot BS was obtained from J. I. Hoffman of the Bureau of Standards as a portion of the Aluminum A described in his report of the redetermination of the atomic weight of aluminum.6 Lot BS originated with the Aluminum Co. of America also, and was received in the standard notch-bar form. A detailed chemical and spectroscopic analysis of lot BS has been published.6 Principa1 impurities were 0.006 pct Si, 0.003 pct Fe, 0.002 pct Cu, and >99.987 pct Al. Samples were processed identically by room-temperature swaging and drawing to 0.030 in. diam wires from approximately ½ in. diam rods that were machined from the notch-bar pigs. Samples were cut from the cold-drawn wire, placed in a grooved graphite block in groups of thirteen, and annealed in air for 1 hr at each of several temperatures, with the results shown in Table I. The two lots of aluminum responded quite differently to annealing, the behavior of aluminum AC suggesting that its grain growth was controlled by impurities. The recrystallization texture of aluminum of similar purity comparably treated has been reported as predominantly <111> fiber with some <100> fiber present.7 Testing Procedure The testing procedure and method of analyzing the data have been described previously.2,3 Annealed specimens of 5 in. gage length were tested in an Instron tensile testing machine at a constant head motion corresponding to a strain rate of 0.04 min-&apos; at a number of temperatures. Instantaneous tenfold rate changes were employed in some tests to deter-

Jan 1, 1958

By R. P. Carreker

GERMANIUM is a member of that group of ele-nents—carbon, silicon, germanium, and tin-— that are currently of particular interest because of their interesting electrical properties. Near room temperature these materials of diamond cubic crystal structure are characteristically brittle. However, there have been several recent reports that germanium is not brittle at elevated temperatures These observations may be summarized as follows: 1—Germanium single crystals were observed to creep in bending under a maximum fiber stress of 6 kg per mm&apos; (8500 psi) at 500°C and above.&apos; The creep curves were sigmoidal, with an induction period that decreased with increasing temperature. 2—Germanium crystal bars have been bent 90" about radii ten times the bar thickness at 550°C.&apos; 3—-Audible clicks were heard when l/4 in. thick crystals were bent drastically at 450" to 500°C. No microstructural evidence of mechanical twinning was observed 4—Slip lines have been observed in tensile specimens deformed at 600°C, and are coincident with (111) planes, the close-packed planes of the diamond cubic structure.] The slip direction has not been established. 5—A germanium crystal has been compressed approximately 20 pct along a <110> axis at 700°C." The results reported subsequently were obtained in exploratory experiments intended to indicate some aspects of the behavior of germanium single crystals when tested in simple tension. Specimen Preparation A single crystal bar of n-type germanium was used. It was of purity corresponding to a resistance range of 2 to 20 ohm-cm. Tensile specimens were prepared by cutting small rectangular bars

Jan 1, 1957

By R. P. Carreker, R. W. Guard

True stress-true strain data are given for nominally pure molybdenum (99.95 pct) over the temperature range -196° to 1540°C (0.027 to 0.63 T/T). Strain rate sensitivity was determined by rate change tests and stress relaxation tests. Inhomogeneous yielding and strain aging effects were observed. The yield stress and tensile strength depend markedly on temperature below 400°C and are insensitive to temperature in the range 400&apos; to 800°C. The ductile to brittle transition range is +25O to -25OC. Strain rate sensitivity and stress relaxation effects are very large near room temperature. THIS report presents data which were obtained as a part of a continuing program designed to document the deformation behavior of nominally pure metals over a wide temperature range.&apos;" This research program is rooted in the conviction that sufficient tensile and creep data of appropriate quality and detail does not exist to provide precise phe-nomenological descriptions of the deformation behavior of pure metals as a function of temperature, strain rate, and grain size. Such knowledge is needed to provide a firm base for alloy design and an experimental framework to which theories of deformation must conform. These data also complement a parallel exploratory program investigating the properties of refractory metals, including molybdenum, and their alloys. Several important reviews and research reports dealing with the tensile properties of molybdenum as a function of temperature have appeared."-lo To a great extent the work to be reported in the present paper parallels that of Bechtold and of Pugh. The authors&apos; results are in general agreement with those of Bechtold and Pugh. The principle differences be- tween the present experiments and those published previously are as follows: 1—Wire specimens prepared from arc cast molybdenum were used in the present experiments, bar specimens were used by others. Pugh used arc cast material while Bechtold&apos;s most extensive results" were obtained on powder metallurgy specimens. 2—Several different grain sizes were tested in the present experiments. 3—The testing range was extended to cover from -196°C to 1570°C (0.027 to 0.63 T/Tm). 4—The combination of wire specimens, high speed autographic recording, and a stiff testing machine permitted more detailed study of initial yielding. The principle contribution of the present paper is the demonstration of the relatively consistent mechanical properties of nominally pure arc cast molybdenum specimens of differing metallurgical history. Bechtold and Scotts have shown that arc cast and powder metallurgy molybdenum samples have comparable tensile properties in the range —75" to 200°C when processed to have the same re-crystallized grain size. Material The molybdenum used in this investigation was supplied by the Climax Molybdenum CO. Its metallurgical history is outlined in Table I; the final products were 0,030 in, diam wire specimens annealed at 1100&apos; and 1375°C and 0.060 in. diam wire

Jan 1, 1957

By R. P. Carreker

THE experiments described in this report were conducted as a part of a general program designed to document the deformation behavior of pure metals over a wide range of temperature. Material The silver used in this investigation was obtained from the Handy and Harmon Co. as 99.97+ pct pure. It is from the same length of wire as that designated H-H silver by Hoffman and Turnbull in their investigation of self-diffusion in silver.&apos; The silver was drawn at room temperature from 0.060-in. diam to 0.020-in. diam wire without intermediate annealing. Seven-inch lengths of cold-drawn wire were annealed in groups of ten by placing them in individual quartz tubes spaced circumferentially about and near the vertical furnace axis. Dry nitrogen was passed through the furnace to lower the partial pressure of oxygen. The three annealing temperatures used were 700°, 800°, and 900°C, producing grain sizes of 0.017, 0.040, and 0.250 mm diam after 30 min at temperature. Testing Procedure The details of the testing procedure have been described in the reporting of results of a similar investigation of Copper. Annealed specimens of 5-in. gage length were gripped between grooved plates. Tests were performed using an Instron tensile-test-

Jan 1, 1958

By John Nunes, Frank R. Larson

Temperature-dependent functions of various ten-sile flour stress and fracture parameters were investigated on iron and low composition alloys of Fe-C, Fe-Cr, Fe-Mn, and Fe-Ni. Data were obtained over a range of test temperatures from 200" to -196"C at an initial strain rate of 0.01 min". It was found that the strain-hardening exponent, n, for the Fe-Ni alloys increased significantly at very low testing temperatures with increased alloy addition. Also, the stress and strain at fracture followed a predictable transitional type of behavior. The flour stress-temperature dependence was evaluated employing the relation: M where it is shown that composition can strongly influence the constant M. An empirical formula describing this effect is as follows: M = MI (equivalent at. pct Ni)-&apos;" Some preliminay, tests were run on several commercial steels to test the generality of this relationship. One of the phenomenological explanations of the ductile-to-brittle transition encountered in bcc metals is based upon the difference in the temperature dependence of the flow and fracture stresses.&apos; It states simply that the fracture stress is relatively insensitive to temperature and that the yield or flow stress rises rapidly with decreasing temperature so that, when the flow stress reaches the fracture stress, brittle fracture occurs without appreciable plastic deformation. This behavior is illustrated schematically in Fig. 1 where the ductile-to-brittle transition is shown to occur at two possible temperatures. The lowest temperature, TB, represents essentially completely brittle behavior while the higher temperature, TN, represents "notch" brittle behavior. Support for this theory is obtained when it is observed that fcc metals, which do not commonly have a ductile-to-brittle transition, also do not commonly exhibit any strong temperature de- pendence of the yield stress. This relationship of yield or flow stress behavior with the ductile-to-brittle transition has been generally observed for most metals representative of these two crystallo-graphic structures. One exception to this generalization appears in the case of tantalum where no apparent ductile-to-brittle transition has been observed. However, this may simply be an indication that this type of behavior is not directly related to the flow stress and that some other mechanism such as twinning must also be operative. A comparison of the typical temperature dependence of the yield stress of some representative metals is shown in Fig. 2. In the study of the temperature dependence of the strength of metallic materials, considerable ef-

Jan 1, 1963

By J. W. Spretnak, G. W. Powell, J. H. Bucher

Tlze room-temperature tensile fracture oj smooth, round specitnens of three ultrnhigh- strength steels tempered to a wide range of strength levels was studied by means by light and electron-microscopic examination of the fracture surfaces. The fracture of AISI 4340 and 300 M at all the strength levels studied, and H-11, except after tempering at 1200° and 1300°F, occurs in three stages. The initiation of fracture is internal (except in some lightly tcmpeved specimers in which fracture is initiated at surface flaws), and is nucleated largely by separation at metal-second phase intevjaces. TIze voids grow and, coalesce to form a crack. When the crack has reached a sufficienl size, rapid propngutio~z ensues. Failure in this stage of fracture usually occurs by dimpled rupture of inicroshear stefis. In the case of H-11 tempered in the 1125° to 1300°F range, fracture in the shear steps is predominantly by concentrated deformation without void formation. The termination of fracture is usually occomplished by the formation of a shear lib in which fracture occurs by shear dimpled rupture. In the case of H-11 tempered at 1200° and 1300°F, no shear lip was obserued, and the radial elelments extend to the surface—a true termination slage does not exist. ThE tensile fracture of several metals and alloys has been investigated.2-4 In the case of polycrystal-line materials, cup-cone fracture usually results. The mechanism of cup-cone fracture may be summarized as follows.5 Cavities are formed in the necked region of the specimen. They usually are initiated by inclusions or second-phase particles. The cavities extend outwards by means of internal necking, and a crack lying about perpendicular to the length of the specimen is formed in the necked region. Subsequent crack growth occurs by the spread of bands of concentrated plastic deformation inclined at an angle of 30 to 40 deg to the tensile axis. Cavities are formed in the bands of concentrated deformation. The deformation bands zigzag across the bar with the net result that mac-roscopically the crack extends about perpendicular to the specimen axis. The final separation, or cone formation, appears to occur by continued crack propagation along one of the deformation bands out to the surface of the specimen. The micromechanics of the tensile fracture of ultrahigh-strength steels have not been thoroughly investigated. Larson and carr6,7 studied the tensile-fracture surfaces of AISI 4340 with a low-power microscope and reported that three stages of fracture could be observed in general. A centrally located region characterized by circumferential ridges, an annular region characterized by radial surface striations, and a peripheral shear lip were found. It was first pointed out by 1rwin8 that the central region is very probably one of fracture initiation and slow growth, and that the annular, radially striated region is one of rapid crack growth. Presumably the crack grows slowly, assuming roughly a lenticular shape, until it is large enough for the initiation of rapid propagation. In this investigation, it was attempted to determine the fine-scale aspects of the room-temperature tensile fracture of some ultrahigh-strength steels, and to relate the variation in fracture mode with microstructure. The steels studied were AISI 4340, 300M, and H-11 tempered to a wide range of strength levels. I) EXPERIMENTAL PROCEDURE The compositions of the steels studied are given in Table I. The steel was received in the form of hot-rolled bar stock 5/8 to 1 in. in diameter from which oversized specimens were machined and heat-treated. The heat treatments employed are given in Table 11. Subsequent to heat treatment, the specimens were ground to the final dimensions and stress-relieved by heating for 1 hr at 350°F (with the exception of the as-quenched steel). Standard smooth round specimens of 0.252-in. diameter and 1-in. gage length were tested in a Tinius Olsen Universal Testing Machine using a cross-head speed of 0.025 in. per min. The relatively coarse aspects of the fracture topography were determined by light-microscopic examination of sections through the fracture surface of nickel-plated specimens. A direct carbon-replication technique9 was used in the electron-microscopic study of the fracture surfaces. The replicas were examined in the electron microscope, and stereo pairs of electron micrographs were taken. The stereo pairs were then examined using a Wild ST4 Mirror Stereoscope. Carbide and inclusion particles extracted in the replicas were analyzed by selected-area electron diffraction. II) EXPERIMENTAL RESULTS The mechanical testing data are summarized in Table 111. The values reported are the average of

Jan 1, 1965

By W. V. Kotlensky, H. E. Martens

In an earlier communication, Ref. 1, the authors presented the tensile properties and the changes in crystal structure and microstructure for pyrolytic graphite tested at 2750°C. It was suggested that changes in the structure may have an effect on the tensile properties of this material. Accordingly, tensile tests were run on a series of specimens, from one block of pyrolytic graphite, some of which were hot-worked by being strained at 2750°C, others of which were tested in the as-deposited condition. According to the manufacturer, General Electric Co., the polycrystalline material was deposited at 2100°C on a substrate of synthetic graphite by decomposition of a dilute natural gas mixture. The as-deposited material had a density of approximately 2.19 g per cc, and as shown in Fig. 1 had a structure of medium size growth cones which were regenera-tively nucleated. All specimens had a gage section 0.75 in. long by 0.06 in. wide by 0.1 in. thick with the thickness direction perpendicular to the basal planes. The testing method and equipment used have been previously described, Ref. 2 and 3. For both the hot-working and the testing the load was applied parallel to the basal planes; a cross-head speed of 1.5 x l04 in, per sec was used. All heating was done in an atmosphere of bottled helium.

Jan 1, 1962

By H. C. Doepken

Wrought Hadfield steel was tested in axial tension at from 100° to —196°C, to determine flow and fracture stresses as well as conventional properties. Ductility and related properties, such as fracture stress, decreased continuously with temperature. Peculiarities during straining indicated possible martensite formation or mechanical twinning. THE mechanical properties of austenitic manganese steel have been summarized by Avery and Day,&apos; including the low temperature tensile properties reported by Hadfield and coworkers2,3 " These data show the material to be brittle at both liquid oxygen and liquid &apos;hydrogen temperatures. Avery and Day report considerable ductility measured by Charpy V-notch tests at —100°F (-73°C). Since the review of low temperature properties by Seigle and Brick&apos; show face-centered cubic structures, other than Hadfield steel, to be ductile at low temperatures, it was felt that a study of tensile properties, between the reported intervals, should show whether the decrease in ductility was a gradual decline or an abrupt drop beginning at some intermediate point. X-ray examinations were made in an attempt to shed further light on the mechanism of work hardening of austenitic manganese steel. The presence of a magnetic constituent has been found by Hall" and by Avery, Homerberg, and Cook." Krivobok7 demonstrated the growth of magnetic constituents on reheating cold-worked specimens. Hadfield2,3 &apos; and Goss found no a iron. The steel from which the specimens were taken was commercial hot rolled % in. diam rod of electric furnace steel. The ladle analysis was as follows: C, 1.40 pct: Mn, 12.11; P, 0.060; S, 0.012; Si, 0.120; Ni, 0.080; and Cr, 0.090. Fifteen specimens, 6 in. long, were cut from a

Jan 1, 1953

By R. L. Smith, J. L. Rutherford

ALTHOUGH considerable effort has been devoted toward the determination of the mechanical properties of pure metals, it is extremely difficult to compare the results of such work. This is because of differences in method and type of test, grain size, heat treatment, and purity of material. There is very little information on the properties of most of the high melting point reactive metals with purities greater than 99.9 pct. Previous evidence concerning the tensile properties of unalloyed metals show that: body-centered-cubic metals are brittle at low temperatures; face-centered-cubic metals remain ductile; some hexagonal metals remain ductile, others become brittle. It may be that traces of impurities decrease the low temperature ductility of the body-centered-cubic lattice. This would suppose that a body-centered-cubic lattice free from impurities would exhibit nearly the same ductility as a face-centered-cubic structure. The present investigation has been an evaluation of the tensile properties of zone refined iron in order to determine whether higher purity material than that previously evaluated behaves differently. Experimental Procedure In this investigation, high purity iron was zone purified using the floating zone method.&apos; The starting stock was obtained from two sources. The first,

Jan 1, 1958

By O. W. Simmons, L. W. Eastwood, C. M. Craighead

The results of a preliminary study of 113 ternary titanium-base alloys are described. The compositions investigated were as follows: 1. Ternary titanium-carbon alloys containing copper, silicon, vanadium, chromium, manganese, iron, or cobalt. 2. Ternary titanium-nitrogen alloys containing chromium. 3. Ternary titanium-chromium alloys containing additions of vanadium, molybdenum, tungsten, cobalt, or nickel. 4. Ternary titanium-manganese alloys containing additions of silicon, chromium, tungsten, or iron. Tensile properties, minimum bend radii, hardnesses, response to heat treatment and aging treatment, and phase relationships for these alloys were determined. THIS is the second in a series of papers describing some of the results of the exploratory work on titanium alloys being conducted at Battelle Memorial Institute for the Wright-Patterson Air Force Base. The work with the binary alloy systems has been described in a previous publication.&apos; This paper describes ternary titanium-base alloys as follows: 1. Ternary titanium-carbon alloys containing copper, silicon, vanadium, chromium, manganese, iron, or cobalt. 2. Ternary titanium-nitrogen alloys containing chromium. 3. Ternary titanium-chromium alloys containing additions of vanadium, molybdenum, tungsten, cobalt, or nickel. 4. Ternary titanium-manganese alloys containing additions of silicon, chromium, tungsten, or iron. As the preparation and fabrication of the 0.5-Ib experimental titanium alloy ingots to 14-ga sheet, as well as the methods of testing used in this work, have been described in the first paper of this series, Titanium Binary Alloys, no further discussion of these techniques will be given. Ternary Alloys Containing Carbon The following alloys of titanium, containing carbon and one other metal, were investigated: 1. 0.25 pct carbon; 1, 2.5, and 5 pct copper. 2. 0.25 pct carbon; 1, 2, and 3 pct silicon. 3. 0.25 pct carbon; 1, 2.5, and 5 pct vanadium. 4. 0.25 pct carbon; 2.5, 3.5, 5, 10, and 15 pct chromium. 5. 0.5 pct carbon; 2.5, 3.5, and 5 pct chromium. 6. 0.25 pct carbon; 1, 1.75, 2.5, 3.5, 5, 7.5, and 10 pct manganese. 7. 0.5 pct carbon; 1.75, 2.5, 3.5, 5, 7.5, and 10 pct manganese. 8. 0.25 pct carbon; 0.25, 0.5, 1, and 2 pct iron. 9. 0.25 pct carbon; 1, 2, and 3 pct cobalt.

Jan 1, 1951

By J. W. Downey, S. T. Zegler

A study has been made of the occurrence of Phases having the Cr30-type structure in ternary alloys having the general composition where B cept for iron and copper all the B components are known to form binary compounds having the type structure. The occurrence of the struc-ture type in the ternary alloys was investigated using optical and X-ray metallography. The results indicate that iron can replace virtually none of the cobalt in at 900°C, but substitutes for nickel in up to the approximate composition tually insoluble in V3Ni and 2) and exhibit complete solubility. At 100 nickel can be substituted for cobalt in VSCo up to the composition solubility; and 3) in the system , two solid solutions having the type structure occur. At the same temperature, 1000°C, V3Si exhibits limited solubility for V3Co but is completely mis-cible with V3Ga. Lattice parameters have been determined for all the binary and ternary phases. The solid solutions are discussed in terms of the atomic size and electronic factors which are thought to influence the occurrence of Cr30-type phases. The binary compounds V3Co and V3Ni were found to form peritectoidally from o phases and a vanadium, the former compound at a temperature between 1000" and 1050°C, the latter between 800°and 900°C. The substitutions of nickel and rhodium for cobalt in V3Co were determined not to raise the superconducting transition of VSCo to a temperature as high as 2°K. The occurrence and stability of binary intermediate phases having the Cr30- or A15-type structure (space group o;, Pm3n) have been discussed in several recent papers . Vanadium forms the struc -ture at or near V3B stoichiometry with all the mem -bers of the cobalt and nickel groups in the periodic table and also with Au, Ga, Si, Ge, Sn, P, As, and Sb. This paper describes a survey of the occurrence of the CrsO-type structure in ternary alloys having the general composition V3(B, B&apos;) where B and B&apos; are Fe, Co, Ni, Cu, Rh, Ir, Si, and Ga. The purpose of the survey was to clarify the connection between the position of the B partner in the periodic table and its CrsO-type phase-forming tenden~y.~,~ The survey is part of a broader program concerned with the influence of B-partner substitutions on the superconducting transition temperatures of CrsO-type compounds. The transition temperatures, Tc, of some of the ternary Cr,O-type phases found in the survey are also discussed in this paper. EXPERIMENTAL PROCEDURE The alloys were prepared by arc-melting in an A-He atmosphere and were in the form of buttons, weighing approximately 3 g. The metals used in their preparation were received from commercial sources and were in all cases 99.9+ pct pure. The vanadium was spectrochemically pure except for 800 ppm of Si, 200 ppm of Fe, and 200 ppm of Al. For study of each pseudobinary system at least seven alloys were prepared having the general compositions case of alloys containing gallium, melting losses were small, usually less than 0.3 pct of the total weight of the metals charged for melting. Melting losses were somewhat greater for gallium alloys, ranging in general from 0.9 to 5.0 pct. Ternary V-Si-Ga alloys were chemically analyzed and the results are given in Table I. The analyses show most alloys to be within 2 at. pct of the intended composition. Throughout the balance of the paper the intended compositions are used in discussing the alloys. The cast buttons were annealed within evacuated quartz capsules for times ranging from 3 to 6 weeks at temperatures between 800" and 1000°C and water-quenched. The phases present in the annealed and quenched alloys were identified by optical and X-ray metallography. X-ray diffraction photograms were made with powders prepared from the bulk specimens, by the use of a 114.6-mm-diam powder camera with copper irradiation. The lattice parameter of CrsO-type phases and their standard errors

Jan 1, 1963

By M. A. Dayananda, R. E. Grace

Vapor-solid diffusion couples were employed in a study of ternary diffusion in the single-phase copper-rich corner of the Cu-Zn-Mn system Interdiffusion coefficients were measured at three different compositions corresponding to the intersections of diffusion paths determined at 850°C. Intrinsic diffusion coefficients were determined at two compositions by the Philibert-Guy method. New relations between the various intrinsic diffusion coefficients are also presented. MOST of the investigations of ternary diffusion have been made with "sandwich" couples formed by welding together two homogeneous solid solutions with or without markers placed at the interface. With such couples Kirkaldy and his coworkers1,2 have determined interdiffusion coefficients which describe ternary diffusion in the substitutional alloy systems, Cu-Mn-A1 and Zn-Cu-Al. Guy and Leroy3 have reported intrinsic diffusion coefficients for Co-Ni-Cr alloys as a function of composition using the analysis of Philibert and Guy.4 The use of semi-infinite diffusion couples, for example, vapor-solid couples, where the diffusing species are supplied from a vapor source in contact with a diffusion disk, has so far been confined to studies of binary diffusion. Balluffi et a1.5,6 have successfully determined the interdiffusion as well as the intrinsic diffusion coefficients in several binary metallic systems employing vapor-solid couples with markers initially placed at the vapor-solid interface. They have also discussed the general advantages and limitations of such couples. Vapor-solid diffusion couples can successfully be employed in ternary diffusion studies. The importance of such couples in ternary diffusion lies in the fact that inert markers lie close to the vapor-solid interface, and the composition of the marker plane is virtually known before experimentation begins. Another minor advantage is that the method of Philibert and Guy4 may be used with three alloys instead of four. Lastly, the diffusion paths need not be S-shaped, as pointed out by Meijering7 for "sandwich" couples. The Cu-Zn-Mn system is convenient for ternary diffusion studies because of the large single-phase field in the copper-rich corner. The vapor pressure of zinc above such alloys is known, and X-ray measurements indicate that the molar volume is nearly constant over a wide range of compositions.8,9 It is the purpose of this paper a) to extend the analysis of Balluffi and seigle5 to permit an evaluation of interdiffusion coefficients in ternary vapor-solid couples, b) to use the method of Philibert and Guy to determine intrinsic diffusion coefficients with such couples, and c) to present experimental data on both interdiffusion and intrinsic diffusion coefficients for the Cu-Zn-Mn system at 850°C. INTERDIFFUSION IN TERNARY VAPOR-SOLID COUPLES The analysis of Balluffi and seigle5 is extended in this section to a vapor-solid couple where two solutes diffuse from a vapor source into a ternary solid solution. To take into account the expansion of the diffusion disk caused by the absorption of mass from the vapor phase, the total fluxes are expressed on the basis of a coordinate system that moves with the velocity of the vapor-solid interface; thus. where D3ij are the intrinsic diffusion coefficients,4 v(x, t) is the local Kirkendall velocity of mass flow, and F(t) is the negative velocity of the vapor-solid interface (x = 0). Both these velocities are based on a fixed coordinate system far removed from the diffusion zone. Using the continuity relations and Darken&apos;s method1&apos; for a ternary system in which components 1 and 2 diffuse from the vapor phase into a solid of negligible vapor pressure, one obtains

Jan 1, 1965

By F. X. Spiegel, D. Bardos, Paul A. Beck

Ti6NileSi7)G is known to be cubic, with 116 atoms in the unit cell. In the present work four new G sili-cides were found with other transition elements and five G germanides. The titanium-group elements are very efficient in forming G phases. The vanadium -group elements are somewhat less efficient, and the chromium-group elements do not form any. Cobalt can replace nickel in most cases, but it is somewhat less efficient. Iron was not found to form any G phases. Germanium is somewhat less efficient than silicon. The (TiNiSi)E phase has been previously identified by Westbrook. However, its crystal structure has not been determined. In the present work the E phase was found to have an orthorhombic unit cell with lattice parameters a. = 7.02, bo = 5.18, co = 11.11A. A large number of other E phases were found at the same stoicheometric ratio in ternary systems of silicon and germanium with transition elements. Lattice parameters similar to those for TiNiSi were measured. The number of atoms per unit cell is estimated at 30 to 36. The vanadium-group elements are somewhat more efficient in forming the E phase than the titanium-group elements. Chromium-group elements were not found to form any E phases. Cobalt is somewhat less efficient than nickel, and iron is less efficient than cobalt. Germanium is not as efficient as silicon. ThE ternary silicide designated as G-phase was first identified by Beattie and versnyder1 and by Beattie and Hagel2 at the approximate composition TiNi2Si. Westbrook et al.&apos; gave the composition as Ti2Ni7Sis. This phase was found to be fcc with a. = 11.201A and 116 atoms per unit cell. Beattie4 and more recently independently Hladishevskii, Kripyakevich, Kuzma, and Teslyuk5 were able to show that this structure is isotypic with MgeSi7CuLe and Th6Mn23. The latter structures had been determined by Florio, Rundle, and snow,&apos; Nagorsen and Witte,7 and Bergman and Waugh.8 According to the crystal structure, Ti8Ni6Si7 may be considered as the ideal composition. The Russian investigators5 identified six additional ternary phases isotypic with Ti6Ni6Si7, in which titanium is replaced by Mn, V, Nb, Ta, Zr, and Hf, and one, namely MneNil6Ge7, in which germanium replaces silicon. They also found that Cr6Ni6Si7, which does not occur as a ternary phase, can be stabilized by replacing 1 at. pct of the Cr by V, Mn, or Ta. Another ternary silicide was found by Westbrook et al.3 at the composition TiNiSi, and it was designated by them as the E phase. The crystal structure of this phase has not been deter-

Jan 1, 1963

By D. I. Bardos

The occurrence of Laves phases (AB,) in various binary alloy systems has been reviewed in recent papers.13 Iron forms Laves phases with Sc-, Ti-, V- and Cr-group elements. However, two of the corresponding Laves phases with cobalt, namely MoCo, and WCo,, are missing. Nickel forms binary Laves phases with Sc and Y, but not with any Ti-, V-, or Cr-group element. These conditions are schematically illustrated in Table I. That electron concentration is a significant factor in determining which type of Laves phase is formed in a given case was recognized long ago by Laves and Witte,4 This concept has been more recently applied to Laves phases formed by transition ele-ments. It now appears that the average electron concentration (average number of electrons per atom outside of the closed shells of the component atoms) may be an important factor also in determining whether or not a Laves phase can occur at all in a given system. As shown in Table I, in the transition element systems considered Laves phases are absent at electron concentrations of 8, or larger (to the right and below the double line in Table I). These absences clearly cannot be accounted for on atomic size considerations, as evidenced by the many inconsistencies apparent in Table 11. While the ideal radius ratio for Laves phases is 1.222, in the systems which do have Laves phases the ratios of CN 12 radii range from 1.10 to 1.46. The consistent absence of Laves phases at RA/RB< 1.10 might be attributed to the low values of these radius ratios. However, Table II shows that absences also occur irregularly throughout the radius ratio range covered. In a previous paper7 it was concluded that, in transition element systems forming u-phases, silicon may act as an acceptor of electrons, thus stabilizing the u-phase at electron concentrations higher than

Jan 1, 1962

By C. G. Dunn, J. L. Walter

SILICON-iron .under certain conditions of processing and annealing will form a cube texture. The occurrence of this texture in thin tapes of silicon iron was first reported by Assmus, Detert, and Ibe.&apos; They concluded that the cube texture in their material formed by a secondary re crystallization process. In the course of present investigations on the development of the cube texture and the mechanisms involved in its formation by secondary re-crystallization it was found that under certain conditions the cube texture was replaced by a cube-on-edge or (110) [001] texture. We call this phenomenon tertiary re crystallization since it follows secondary recrystallization. The secondary recrystallization occurs after primary recrystallization and a substantial amount of normal grain growth. It is the purpose of this paper to describe briefly the occurrence of tertiary recrystallization and the structure and texture prior to and after tertiary recrystallization. The primary and secondary recrystallization structures and textures are the subject of another study and will be referred to briefly in the discussion of results. Evidence will be presented for the view that the dominant driving force for tertiary recrystallization is derived from the surface free energy which depends on the crystallographic planes presented to the gas-metal interface. In the absence of boundary restraining forces (dispersed phases, thermal grooves, and so forth) a sufficiently large difference in surface free energy across a grain boundary would produce boundary migration away from the center of curvature of the boundary intersection line in the surface of the sample according to risher.2 Therefore, particular attention was paid to observation of the direction of motion of these surface boundary lines. Furthermore, measurements of the actual boundary curvatures could be used to estimate the driving force supplied by surface free energy. EXPERIMENTAL PROCEDURE Samples were prepared by hot and cold-rolling vacuum-melted ingots of iron containing 3 pct Si to strips 0.012 and 0.006 in. thick. The ingots contained less than 0.005 wt pct impurities. The rolled strips were cut into samples approximately 1 in. wide by 3 in. long and were electropolished to remove surface imperfections. All samples were annealed in a vacuum at a pressure of approxi-

Jan 1, 1960

By A. J. Hatch

Recent disclosures concerning strengthening of anisotropic sheet materials under biaxial stresses are reviewed. This biaxial strengthening has been termed "texture hardening" by Backofen and is related to the normal anisotropy parameter, R. Measurements of R made on three annealed titanium alloys indicate a Ti-4Al alloy will have the greatest potential for texture hardening, followed in order by Ti-5Al-2.5Sn and Ti-6Al-4V. Relative to sheet texture, a review of the deformation mechanisms in titanium alloys would indicate that a preferred orientation for "texture hardening" is where the (0001) pole is normal to the sheet plane. This was confirmed by determination oj (0001) pole figures for the above three alloys. High R values are related to pole figures with (0001) normal to the sheet plane, whereas low R values are related to pole figures with (0001) split in the transverse direction from the sheet normal. PUBLISHED work by Backofen et al. 1 and Hosford and Backofen2 based on hill&apos;s&apos; mathematical theory of yielding in anisotropic metals has indicated that significant strength advantages under biaxial stresses are possible in hcp alloys, provided an ideal (0001) [1010] sheet texture exists. This tendency to strengthen under biaxial stresses has been termed "texture hardening".1 It can be quantita- tively characterized by a strain-ratio parameter, R, where R may be determined from the ratio of natural width strain to natural thickness strain as measured on a uniaxial test specimen cut parallel to the rolling direction. If R is greater than 1.0, then strengthening under biaxial tension would be predicted. Considering the a titanium alloys which have an hcp crystal structure, Williams and Eppelsheimer4 have reported that commercial-purity grades deviate from the ideal (0001) sheet texture mainly in an approximate 30-deg rotation of the basal plane, toward the transverse sheet direction. However, McHargue et a1.5 have shown that the angle of rotation of the basal plane is sensitive to alloy content. In particular, a titanium alloy with 3.8 pct Al was reported to have the ideal (0001) [1010] texture after cold rolling. Additionally, Backofen et a1.1 using plane-strain compression tests have reported R values greater than 5.0 for the commercial 0 alloy Ti-5A1-2.Sn (110-AT). In view of these findings, a program was initiated to further determine the potential of commercial Ti-5A1-2.5Sn and an experimental Ti-4A1 alloy with respect to "texture hardening". Also, because of its wide commercial use, Ti-6A1-4V was included in the investigation. BIAXIAL-STRESS THEORY FOR ANISOTROPIC MATERIALS Backofen ef al.,1 based on the mathematical theory of Hi11,3 has introduced the concept of "texture hardening". In this work, it was shown that possi-

Jan 1, 1965

By S. R. Goodman, Hsun Hu

The rolling texture of an 18-8 stainless steel (Type 304L or 304) has been found to change gradu -allv from the (110)[112] brass type to the (123)[412] copper type as the rolling temperature increases from 200° to 800°C. This texture transition is accompanied by a change in the stacking-fault frequencv or probability, similar to those found pre- In a series of previous publications, it was shown that the rolling texture of high-purity silver changes gradually from the brass type to the copper type with increasing temperature of deformati~n. A similar temperature dependence of rolling texture transition was found in electrolytic copper3—by lowering the temperature of deformation, the normally observed rolling texture of copper changed gradually viously in nonferrous fcc metals—the stacking-fault frequency decreases rapidly with the formation of the copper type texture. These results strongly indicate that the temperature dependence of copper type == brass type texture transition and its correlation with stacking-fault frequency is a general phenomenon in fcc lattices. to the brass type.* In both of these metals, the

Jan 1, 1964

By S. R. Goodman, Hsun Hu

The rolling texture transition in copper as a function of deformation temperature is found to be quite similar to that in high-purity silver. The ordinary copper type texture changes gradually to the brass type, accompanied by an increase in the stacking fault frequency, as the temperature of deformation decreases. The rolling textures obtained at low temperatures and the corresponding annealing textures compare remarkably well with the ordinary rolling and annealing textures of low-zinc brasses. Furthermore, it is also found that the observed stacking fault energies of copper specimens deformed at low temperatures correspond closely with those of low-zinc brasses, as deduced from measurements of the dislocation nodes. These results are consistent with the idea that texture transition may be significantly dependent on the stacking fault energy. It was shown in a series of previous publications1"3 that the rolling texture of high-purity silver changes gradually from the common silver type (or the brass type) to the copper type* with increasing temperature *The silver or brass type rolling textures are commonly described as (110)[112],4,5 whereas the copper type rolling textures are mainly (123) [a121 and (146)[211].5 Other ideal orientations have been used by various investigators for the copper type rolling textures, but the actual difference among these ideal orientations is rather small. of deformation, and that such texture transition can be correlated with the change in stacking fault frequency as a function of rolling temperature. Furthermore, it was shown that the formation of annealing textures in high-purity silver from the various rolling textures obtained during the course of texture transition depends entirely upon the deformation texture. From a common silver type or brass type rolling texture (such as produced by rolling at 0" or at 50° C) the recrystallization texture is mainly (120)[211], which is quite different from the recrystallization texture of brass.* However, for both ma- terials the orientation relationship between the rolling texture and the recrystallization texture can be expressed in terms of similar rotations around [Ill.] axes of approximately 30 or 35 deg. They differ in that the [Ill] rotational axes are different. In the case of high-purity silver, the axes of rotation are the two (111) poles near the plane normal of the deformed specimen with a (11l)[112] type rolling texture, whereas in brass, they are the (111) poles near the direction of rolling. For high-purity silver with a predominantly copper type rolling texture (such as produced by rolling at 150° or 200°C) the recrystallization texture is mainly (100)[001], or the cube texture, plus its twin orientations. It was shown previously6 that such reorientation observed in copper or aluminum can be described as [Ill] rotations of 30 to 40 deg. For intermediate cases, i.e., when the rolling textures are composed of a

Jan 1, 1963