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By C. Chiou, D. P. Seraphim

The effect of clustering on the superconducting properties of Al-Zn alloys has been studied by the ballistic induction techniques. The superconducting critical temperature, T,, changes when zinc-rich clusters are formed. The change in T, depends on the concentration and may be positive or negative, but always correlates with an effective decrease in the solute concentration of the alloy. Reversion after clustering effectively returns a 5 at. pct Zn alloy to the quenched state. These methods might be use&lly extended to the Al-Mg system, where X-ray data are difficult to obtain. In very dilute Al-Zn alloys, as in other tin,' indium, aluminum,2 or tantalums base solid solutions, there is a decrease in the superconducting critical temperature, Tc, roughly proportional to the temperature independent resistivity, i.e., -0.2°K per pS2cm. With increasing solute concentration, the T, generally reaches a minimum and then increases4 to a value well above that for the pure solvent. The change of T, with solute concentration follows an empirical formulation involving two constants which correlate with the scattering cross section of the impurities. As will be shown below, the Al-Zn alloys more concentrated than 1or2at. pct do not follow the empiricism, probably because they are unstable even when rapidly quenched.5 However, it seems possible to study at least the later stages of solute clustering by precise measurements of the 1', following various thermal treatments. A few experiments on quenched alloys, aged alloys, and tests on reversion are reported here. The heat treatments selected for obtaining the quenched and aged states were mainly deduced from the work of Dash and Fine.6 For the quenched state, specimens were homogenized at 480°C, quenched to -160°C, and immediately transferred and mounted in liquid nitrogen. Measurements on the superconducting properties and residual resistivities were then made at helium temperatures, after which the specimens were aged at room temperature for 3 to 5 hr, i.e., much longer than the time required for clustering to be essentially completed. The superconducting properties and residual resistivities were once again examined. To reach the reverted state of the A1-5at. pct Zn alloy, the aged specimen was heated for 5 min at 200°C and then tested as above.

Jan 1, 1963

By B. H. Kear

Clustering of slip bands in different systems has been observed in both ordered and disordered Cu3Au crystals compressed in the <110> orientation. Certain features of the slip band patterns are explained in terms of reactions between dislocations with Burgers vectors at 120 deg; in a few rare cases, the slip band evidence suggests that the Products of such reactions form diffuse tilt boundaries between the differently slipped regions of the crystal, similar to the known behavior of ionic crystals, but not nearly so well defined. WHEN crystals with orientations favorable for multiple slip undergo plastic deformation, slip may occur in different systems in localized regions, so that clusters of slip bands are developed. This phenomenon has been observed in fee1-4 metals and also in ionic crystals having the sodium chloride structure.5-8 In the ionic crystals, the boundaries formed between the differently slipped regions of the crystal assume characteristic orientations and are associated with macroscopic tilts that are plainly visible on the sides of the crystal. Recently, this interesting feature was explained in terms of the formation of diffuse low-angle boundaries as a result of interactions between families of glide dislocations in the different systems meeting in the boundaries.9 This paper describes an attempt to obtain similar evidence for boundary formation in an fee crystal. The alloy Cu3Au was chosen for the present study simply because this material is currently being used in a general study of differences in deformation behavior of the ordered and disordered phases of this alloy, which will be described elsewhere. EXPERIMENTAL PROCEDURE The two specially oriented compression specimens used in the present study were prepared from a melt-grown single crystal by careful cutting, grinding and polishing. The finished specimens were about 6 mm by 2 mm in cross section by 6 mm tall, with their ends parallel to (101) and sides parallel to (111) and (121); the broad faces of the specimens were made parallel to (711). Both specimens were annealed in an evacuated quartz tube at 850oC for several hours, slowly cooled to 500°c, and then quenched into water in order to obtain the disordered state for the alloy. One of the specimens was partially ordered (with S- 0.8) by reannealing at 350oC for 100 hr. After these annealing treatments, both of the specimens were electropolished. 10 Finally, the specimens were compressed in the [loll orientation and their stress-strain curves were recorded. The deformation was carried out in increments of a few percent strain. At each stage in the deformation the specimens were unloaded, removed from the compression jig, and examined optically for slip bands. Before reloading, the specimens were repolished, so that the observed slip band patterns were characteristic of a particular stage in the deformation. EXPERIMENTAL RESULTS Fig. 1 show!; the stress-strain curves for the ordered and disordered specimens and also indicates the stages 1, 2, and 3 where the deformation was interrupted for examination of slip bands. The curves show marked differences in flow stress and in work hardening rate in the two specimens, as observed by previous workers.11,12 Also, it was found, in agreemement with the observations of Taoka and Sakata,&apos;3 that the plastically deformed disordered alloy showed the characteristic coarse slip bands of alloys, whereas the ordered alloy exhibited a fine slip band structure similar to that of pure metals. Slip Band Observations. In what follows, the description of the slip band patterns, with accom-

Jan 1, 1962

By Shozo Saito, P. A. Beck

Mettrllographic and X-ray diffraction study of Cb-Co alloys in the Composition range of 7 to 33 nt. pct Cb, after annealing at 1175 °&apos;, showed that near 25 al. pct Cb on MgNi,-lype hexagonal Laves Phaseexists in a narrow composition range, and that an MgCl- type cubic Laves Phase occurs between approximately 27 and 32.6 ut. pd Cb. Lattice parameter and density measures indicated that, in both phases. the deviations from proper Loves sloichimetry result from the substitution of cobalt atoms for some 0f the columbium atoms. Tlle same phases occurant 1000&apos;C, together with an additional phase of unknown structure, which appears between the hexagonal Laves phase and the cobalt-bose terminal solid solutions KOSTER and Schmid&apos; identified a phase corresponding to the composition VCo,, and more recently the crystal structure of this phase has been determined.&apos; In the Ta-Co system Korchynsky and Fountain recently found two intermediate phases at the composition Taco,, one of them metastable. In contrast to the V-Co and the Ta-Co systems, in the Cb-Co system no intermediate phase has been found4 at the composition CbCo,. The present investigation was undertaken in order to reexamine the question of the existence of such a phase. EXPERIMENTAL PROCEDURE Twelve alloys were prepared by arc-melting in a water-cooled copper crucible under helium atmosphere. Electrolytic cobalt and 99.9 pct pure colum-bium powder have been used as starting materials. It was found that melting losses can be reduced by compressing the colunlbium powder in the form of thin pellets before melting. For the alloys used the melting losses were not higher than 1 pct. The intended compositions of all alloys and the chemical analyses of three of them are given in Table I. Specimens from all alloys were annealed in evac- uated fused silica tubes at 1175°C for 3 days and quenched in cold water. A second set of specimens of most alloys was annealed at 1000°C for 7 days and then quenched in cold water. The annealed and quenched specimens were examined metallographically, using the following etchant: 60 pct glycerine +20 pct H,NO, + 10 pct HF + 10 pct water. Powder specimens for X-ray diffraction were prepared by crushing annealed solid specimens in a mortar. Alloys containing 27.3, 28, 29.7, and 32.6 at. pct Cb annealed at either 1175" or 1000°C were very brittle and it was found unnecessary to reanneal the powders. However, alloys containing 24.8 at. pct Cb, or less, especially those annealed at 1000°C, were much less brittle and re-annealing was required to remove the strains present in the crushed powders. In each case reannealing was done at the same temperature at which the corresponding solid specimens were annealed and it, too, was followed by quenching in cold water. In many instances X-ray diffraction patterns were also taken of the polished and etched solid specimens and compared with the corresponding X-ray diffraction patterns obtained with powders. The X-ray diffraction patterns were taken with an asymmetrical focusing camera, using CrK radiation. For precision lattice parameter measurements some X-ray diffraction patterns were taken with a symmetrical focusing camera, again using CrK radiation. EXPERIMENTAL RESULTS A microscopic examination of the alloys annealed at 1175°C revealed that the alloy containing 25.5 at. pct Cb was composed of a single phase, but that the 24.8 at. pct Cb alloy did contain a small amount of a second phase, identified by means of X-ray diffraction as having a fcc structure, undoubtedly the terminal solid solution based on cobalt. Apart from the few very weak lines corresponding to this minor phase, the X-ray diffraction patterns of these alloys could be well interpreted in terms of a MgNi,-type hexagonal Laves phase structure. The indexing of the X-ray diffraction pattern, Table 11, yas based on the lattice parameter values a, = 4.740A and c,

Jan 1, 1961

By W. K. Hardy, H. H. Stadelmaier, L. J. Huetter

The ternary system Co-Sn-C was studied in the cobalt-rich region. The liquidus projection ulas established and two fonr -phase reactions were found. An isothermal section at 950°C was worked out. In the region studied the following phases were observed: cobalt-rich solid solution, graphite, binary compound y -cobalt-tin, and a ternary compound. As a result of the limited solubility of carbon in liquid cobalt-tin alloys, the investigation was limited to the cobalt-rich region. Tin contents up to 32 at. pct were studied. Cobalt of 99.99 pct purity, tin of 99.97 pct purity and synthetic graphite of 99.93 pct minimum purity prepare the alloys. Carbon determinations of all alloys were obtained by gas volumetric analysis. Samples were investigated microscopically and by X-ray diffraction using CoK, and FeK, radiation. Cooling curves were obtained with a number of alloys to determine the temperature of the four-phase reactions. The thermocouple was calibrated with the melting points of high-purity silver and copper. In the region studied the following phases were observed: Cobalt-rich solid solution (face-centered cubic) Graphite. Binary compound y (solidifies congruently at composition Cop~n)&apos;. Ternary compound k which is face-centered cubic with superlattice and nonferromagnetic.&apos; In Fig. 1 the fields representing the projections of the liquidus surfaces were determined by microscopic investigation and are labled by their primary phase. In the field labled C graphite is the primary phase. During solidification in this field much of the primary graphite was expelled through the surface of the melt. It could not always be established microscopically whether the remaining graphite was primary or purely eutectic. Therefore it was not possible to determine accurately the boundary curve which drops from the binary eutectic el at 1309 "~~ to the point of the four-phase reaction at F. The ter-nary carbide is the primary phase in the field labled k. To the right of this field the primary graphite is consistently nodular whereas flake graphite is observed at tin contents somewhat below F. The ternary carbide does not appear to solidify congruently. The graphite is attacked in a peritectic reaction along the boundary between k and C. Thus no graphite is observed in alloys in the k field. In alloys right of the boundary between k and C the graphite nodules are located in the center of the k crystallites, further documenting the peritectic nature of the reaction with the graphite. A four-phase reaction was found at F and 1120 &3"C. The most likely reaction is L+ CY-Co + C — k. Below the temperature of the binary eutectic ez at 1112" a second four-phase reaction was observed at E and 1100 lt3"C. It is the ternary eutectic L — CY-Co + k + y. Figs. 2(a ) to (e) show microstructures of the ternary alloys in the as-cast condition. Fig. 2(a) shows the eutectic a-Co + C (with some primary graphite) with the carbide k in the grain boundaries. Fig. 2(b) shows primary cobalt (dark areas) and the eutectic a-Co + k which is formed along F-E.

Jan 1, 1961

By R. B. Russell

The expansion coefficients of hexagonal (a) zirconium have been calculated from the lattice expansionparameters of both coefficientsofhexagonallow and high hafnium alloys in the range 0° to 600°C. It is found parametersofboththat the coefficients are straight-line functions of the temperature and that the effect of hafnium is to depress the mean coefficients by about 1 pct at room thattemperature and 9 pct at 600°C. The cubic (ß) form was stabilized so that its coefficient could be determined from 400&apos; to 900°C. COEFFICIENTS of thermal expansion for hexagonal-close-packed (a) zirconium were determined up to 600°C (1112°F), without any interference from the effects of preferred orientation, by measurement of the lattice parameters obtained from a symmetrical back-reflection focusing X-ray camera. This paper combines the results for a 1.2 atomic pct hafnium alloy from an earlier report&apos; with those for hafnium-free zirconium. The problem confronting the investigator is essentially that of getting sharp, high-angle diffraction lines which represent the material at a known constanl. temperature and chemical composition. In order to do this, it was necessary to use the proper X-ray emission wave length (Co K was chosen) and to have a fine-grained metal whose atomic spacings would be free from effects of cold work (the specimen was annealed) and to prevent a changing composition owing to gas absorption during the heating of the metal to temperature (the sample was heated in a vacuum). A summary of the effect of temperature on properties is listed in Table I which gives data for ordinary zirconium, as well as new values for the low hafnium metal. The coefficients of the linear thermal expansion that have been reported are given in Table 11. The discrepancies shown may be explained perhaps by the effect of preferred orientation, and McGeary and Lustman2 have done this successfully for sheet metal. The differences in texture cannot explain all the differences in expansion coefficient shown in Table II. liesults by the author agree most closely with those of Kroll.3 Chemical analysis of the high hafnium metal pre- viously reported1 showed that it was 98.22 atomic pct (97.40 wt pct) pure. The low hafnium metal was dehydrogenated Westinghouse crystal bar, designated SMZ-137A. Its analysis was as given in Table 111 (chemical analyses 5377, 5431, and 5432).

Jan 1, 1955

By M. G. Nicholas

The influence of deformation, relative movement and surface cleanliness on the strength of room -temperature "interparticle" bonds between copper surfaces has been studied. No bonding resulted when clean surfaces were merely pressed together, even under pressures of 50,000 psi. Bonding resulted, however, when clean surfaces were simultaneously pressed and moved relative to each other, some bonding occurring no matter how small the pressure or relative movement. Exposure of the clean surfaces to various chemicals inhibited bonding, none at all being produced by small relative movements. The size of the movements which first produce bonding between exposed surfaces depends on the nature of the chemical, being less than 0.002 in. for cleansing reagents and greater than 0.004 in. for lubricants, water and hydrogen sulfide. It is concluded from the experimental data that the prime requirement for bonding is the dispersal of surface films such as oxides and adsorbed contaminants. THE formation of interparticle bonds during the room-temperature compaction of powder metallurgy components greatly increases the ease of handling in subsequent production operations prior to sintering. The strength of unsintered compacts follows the same general trends as their density and hardness, but quantitative data are extremely Scarce. Quantitative correlation of the data on the overall strength of a compact with that of the individual interparticle bonds is virtually impossible due to the extremely complex geometry of powder compacts. Qualitatively, however, the strengths of unsintered compacts can be correlated with variations in parameters known to influence the extent of cold bonding, and hence interparticle bonding, such as surface cleanliness, relative movement, and deformation of the mating surfaces. Kuzmic and seelig4 found that contamination of powder particle surfaces with lubricants reduced the strength of unsintered iron and copper compacts. Interparticle shear has been shown to promote bonding by Train and Hersey7 while Gregory has shown that die rotation is beneficial. Eilender and schwalbe6 concluded that most bonding occurred during the early stages of pressing iron powder compacts when interparticle shear is one of the principal processes occurring. mile there can be little doubt that these and other factors, such as particle size and shape, affect the extent of interparticle bonding produced by compaction, evaluation of the influence of individual factors is not easy. Particle deformation and relative movement, for example, occur simultaneously during compaction, making any precise measurement of their relative influences difficult. Because of these and other difficulties encountered in studying interparticle bonding in actual powder compacts, the approach adopted in this study was to

Jan 1, 1963

By C. H. Li, R. J. Stokes, T. S. Liu

Poly crystalline silver chloride specimens having different micro structures were prepared by extruding monocrystals or precompressed powder at varying temperatures. Extrusion at high temperature (370°C) produced a Fully recrystallized equi-axed material which was brittle and fractured by simple cleavage below its transition temperature (60°C for notched specimens). Extrusion at low temperatures (down to -196°C) produced an extremely tough cold-worked material with grains elongated in the extrusion direction. This material fractured below its transition temperature (which varied from —20" to —200°C) by an intergranular mode. It is concluded that cold work increases the resistance to transgranular cleavage. Fracture is then de-flected to propagate intergranularly ouer surfaces making an angle with the tensile axis with a corresponding improvement in notch toughness. This deflection is greater for extruded powder than for extruded monocrystals because of the extremely elongated grain shape. PREVIOUS work1,&apos; on the fracture properties of fully recrystallized silver chloride under notch impact conditions has indicated that the fracture behavior is rather invariant. Thus polycrystalline materials originating from a number of sources ranging from relatively pure Harshaw single crystals to commercial-purity powder and prepared by hot extrusion at 370°C all fracture by simple cleavage with a ductile-brittle transition at 60°C, see Fig. 1. This behavior is also typical for fully annealed single crystals and appears to be fundamental to silver chloride for the particular loading conditions and notch configuration adopted. One way in which this characteristic fracture behavior can be modified is by the dispersion of alumina particles in the silver chloride matrix.&apos; Then, particularly when the alumina particles are in the shape of spheres, the duc tile-brittle transition is much less sharply defined and occurs over a range of temperature, the transition temperature is lowered considerably and the fracture mode changes from simple cleavage to a fibrous or shear mode. The reason for this change is believed to be partly due to the nucleation of cavities in the region of the dispersed alumina particles which influences the state of stress locally and the distribution of plastic constraint at the root of the notch. It is emphasized however that this is not the sole factor responsible for the increase in toughness. In extruded materials it is reasonable to expect that the resultant grain size, texture, and the residual cold work will be influenced by the presence of a dispersed phase. The present experiments were initially designed to determine the extent to which cold work influenced the fundamental ductile-brittle transition temperature and fracture mode of silver chloride. They were stimulated by the discovery that silver chloride could be extruded at temperatures as low as -25°C when commercial-purity powder was used and as low as -196°C when single crystals were used. This working temperature was far below that previously used (370°C) and it was found that such extruded material was exceedingly tough. The reasons for the enhanced toughness and its relation to the previous work on silver chloride-alumina compacts will be discussed. EXPERIMENTAL PROCEDURE Materials Used. Two forms of silver chloride were used in this study: i) silver chloride powder

Jan 1, 1964

By P. K. Koh, G. C. Dunn

Si-Fe single crystals in a number of selected orientations were cold rolled 70 pct and analyzed to obtain quantitative (110) pole figures. Stable end orientations were ddetermined, and the effect of orientation on deformation behavior was investigated. WHEN single crystals are cold rolled sufficiently, they rotate into end orientations that are stable or unchanged in position with additional deformation. The end orientations of heavily cold-rolled iron crystals, positioned in holes in copper to limit lateral flow and to simulate polycrystalline conditions during deformation, were determined by Barrett and Levenson.&apos; They found that the sum of the end orientations, a pole figure for all the crystals studied, reproduced the essential features of the cold-rolled texture of polycrystalline iron. In. an investigation of Fe-Si alloys, Barrett, Ansel, and Mehl&apos; observed little effect of silicon content on the textures of heavily cold-rolled samples. Since alloys with less than 4 pct Si are known to have the same slip systems as iron," it seems logical to expect that they would also have the same end orientations as those for iron. Whereas Barrett and Levenson in their study of single crystals were interested in the problem of the origin of the cold-rolled texture of polycrystalline iron, the present authors are more interested in the final cold-rolled textures of single crystals of Si-Fe from the point of view of their relationship to subsequent recrystallization phenomena which will be reported in another paper. Nevertheless, such points as the tendency of an end orientation not to rotate (stability), the tendency of a pole concentration to spread out, and the tendency of the crystal to widen during cold rolling become interesting separately when the initial orientations are selected already in end orientations. Also the initial orientations having a [I101 axis parallel with the transverse direction, which include the two end orientations (111) [Ti21 and (001) [110], comprise a series that needs further clarification regarding the way such crystals rotate during deformation. Finally the present use of quantitative pole figure methods* represents an advance over the X-ray photographic method employed by Barrett and Levenson in their work on end orientations. The present investigation provides information on the cold-rolling behavior of two groups of Si-Fe crystals: 1-—crystals initially oriented in recognized end orientations for iron crystals with a [I101 direction parallel with the rolling direction, and 2— crystals initially oriented with a [I101 direction 90" to the rolling direction and in the rolling plane. For the second group, there is a 90" range between the (001) [110] orientation and the (110) [001] orientation. Since the 35" portion of this range from (110) [OOl] to (111) Di2] was included in a previous investigation," only the remaining 55" portion from (111) fii2] to (001) [I101 needs further consideration. Also, since both (111) [xi21 and (001) [I101 are end orientations for iron crystals,&apos; there might be a critical orientation* or orientation range be- • W. W. Martin and C. G. Dunno noted that (112) [Till was such a critical orientation. tween them such that early in the deformation an orientation within this critical range splits into two orientations. Upon further deformation, these two orientations rotate to produce a final texture consisting of (001) [1101 and (111) [?i2] components. The end orientations obtained are discussed in connection with the work of Barrett and Levenson, while the observed tendency for lateral flow of crystals in these orientations is treated in terms of the theory of Hibbard and Yen.&apos; Lateral flow according to the theory increases with deviation of the plane containing the rolling direction and the slip direction from a position perpendicular to the plane of rolling. Experimental Procedure Single crystals in sheet form with preselected orientations were prepared by methods that have been described elsewhere." Three lots designated A, B,

Jan 1, 1956

By Y. C. Liu, R. H. Richman

From rolling textures of polycrystalline AgCl determined as a function of temperature, changes in the amounts of the individual textural components are correlated with temperature-dependent changes in glide geometry. The preferred orientations are influenced by the slip behavior as follows: 1) pencil glide, characterized by (110) slip directions and a greater participation in the deformation the higher the temperature, leads to a(012) fiber texture; 2) low-temperature deformation, primarily on (110) (110) systems, generates the {237}-(11, 9, 7) component; and 3) slip on {111}(110), comprising a significant portion of the room-temperature deformation, generates the {135}(211) and {110}(112) components. Typical re crystallization textures are also presented and briefly discussed. SPECIFIC relations between flow mechanisms in crystals and the individual orientations or components in deformation textures are necessary for a thorough understanding of the origin of such textures. Within each ordinary class of crystal struc- tures there is insufficient variety in the observed preferred orientations to permit direct and unambiguous correlation between flow mechanism and texture. A second approach to this problem is the correlation of textural changes with some known variation in glide geometry within one substance,1 In this paper, the latter approach is applied directly to a material in which there are distinct changes in the character of glide with changes in temperature. Silver chloride, although a nonmetal, is reported to exhibit slip,2,3 fiber textures,4 and rolling textures5,6 crystallographically similar to those of the well-studied fcc metals;7,8 in addition, there is a marked temperature sensitivity of the glide planes upon which flow occurs.3 Previous investigations of textures in AgCl unfortunately did not employ methods sufficiently sensitive to define possibly important minor components in the texture, nor did they utilize the temperature dependence of the glide systems. Thus, in order to clarify the effects of changes in slip mechanism, this paper reports a survey, using semiquantitative methods, of preferred orientations in rolled AgC1. Inasmuch as the influence of prior deformation texture upon preferred orientations after annealing is also of interest, some representative annealing textures are presented.

Jan 1, 1965

By N. K. Chen, R. Maddin

THE textures of straight-rolled and of cross-rolled molybdenum were first determined with the aid of pole figures by Custers and Riemersma.&apos; These authors have shown that for straight-rolling, the main texture was (100) [0111 with a considerable spread about the rolling direction. After cross-rolling a second texture was obtained in which there was a [111] axis perpendicular to the rolling plane, possessing a rotary symmetry. Recently Semchy-shen and Timmons2 in their study of preferred orientation of arc-cast molybdenum sheet obtained similar pole figures. In these two investigations, polycrystalline materials were used. In the main, the textures of molybdenum in these cases are qualitatively consistent with rolling textures of other body-centered cubic metals, such as iron and Fe-Si alloys." In the present investigation, single crystals of molybdenum were deformed by rolling with respect to specific crystallographic planes and directions. Since the deformation textures of polycrystalline aggregates have been related intimately to the behavior of individual grains,4 a study of possible correlations of final textures with respect to the initial orientations of the single crystals might conceivably lead to a better understanding of the textures. Annealing treatments were carried out after deformation in order to study the recrystallization characteristics as well as the recrystallization texture. Rolling of Single Crystals Single crystals of molybdenum (about 1/8 in. in diarn x 1-in. long), grown by the method previously described,5 were cold-rolled without intermediate anneals. Except for specimens MR-1 and MR-2, which were rolled at random in the preliminary test in order to find a proper rate of reduction, all other specimens were carefully documented as to their respective orientation before deformation by means of X-ray Laue back-reflection methods. A stereo-graphic projection was constructed from which it was possible to determine a particular plane and direction to be considered for rolling. Parallel flat surfaces were then polished metallographically with respect to a reference scratch on the specimen which was mounted in sealing wax. These flat surfaces contained the plane chosen for rolling. After etching, a Laue back-reflection photogram was again made of one of the flat surfaces to determine the exact plane of rolling. The initial orientations of all the specimens (specified by two elements for each crystal, direction of rolling and the pole of the rolling plane) are summarized in Fig. 1. The specimens were mounted in plasticene between two hardened Si-Fe plates. The sandwich was then rolled down, keeping the direction constant in the case of straight-rolling. In cross-rolling (MR-10, MR-11, and MR-12) the sandwich was rotated 90" after each pass so that every first and third pass contained the same rolling direction. Specimens MR-14, MR-15, and MR-16 were compression-rolled following the manner of Barrett." In these cases the direction of the sandwich through the rolls was changed by 3" to 5" for each pass until the specimens were rolled to the desired reduction in thickness. It was found that if extreme care was exercised in rolling, i. e., the amount of deformation per pass was extremely small, the specimens could be rolled to a large reduction in thickness. The rate of reduction used, with a single exception of MR-1. was generally 0.001 in. per 10 passes. Thus, it took about 800 passes to reduce a specimen to a final reduction in thickness of 87 to 98 pct. The results of rolling are shown in Table I. Cold-Rolling Textures The final orientations were determined by the conventional Laue transmission method," using molybdenum radiation, with or without a zirconium filter. Pole figures were plotted from (1101 and (200) reflections in which three degrees of intensity— heavy, medium, and light—were determined visually. Three main types of textures have been observed from the deformed single crystals. These are: 1—a

Jan 1, 1954

By J. A. Berge, R. S. Mateer, R. G. Aspden

IN a study of the cold-rolling characteristics of iron-base alloys with 3 to 12 pct Al, the influence of aluminum content on the cold-rolling textures was evaluated. This work was undertaken to determine if binary iron-aluminum alloys possessed cold-rolling textures similar to other body-centered-cubic alloys and to contribute additional information on the effect of solute concentration on the cold- rolling textures. Gradual changes in cold-rolling textures with an increase in solute content have been observed in copper1 and in titanium.2 Iron-base alloys containing 3.0, &apos;6.0, and 12.0 pct aluminum were vacuum-induction-melted in a magnesia crucible and permanent-mold cast. The materials used were high-purity electrolytic iron and 99.99 pct grade aluminum. Each ingot was 1 in. square in cross section by 12 in. in height. The bottom half of each ingot was soaked for 1 hr at 1000°C and rolled at 1000°C to a thickness of 0.055 in. with the long axis of the ingot parallel to the rolling direction. The strip was reheated between passes to maintain a rolling temperature of 1000°C.

Jan 1, 1960

By O. N. Carlson, H. A. Wilhelm, J. M. Dickinson

On the basis of microscopic studies, melting-point observations, and X-ray analyses, a phase diagram is proposed for the Cb-V system. A complete series of solid solutions is formed with a minimum in the solidus at 1810°C near 35 wt pct Cb. No compounds or intermediate phases were found in the system above 650°C. THERE is an ever increasing need for better structural metals and alloys for use in nuclear reactors. In addition to the normal properties of engineering structural materials, such as high temperature strength, resistance to corrosion, and fabric-abil~ty, the nuclear properties of the material must be considered. In a nuclear reactor it is important to conserve neutrons, so a material which removes these neutrons from the reaction excessively is considered to have unfavorable nuclear properties. In nuclear-reactor design the engineer must have nuclear as well as other data available on alloys in order to make a wise selection of materials. Due to the fact that many of the common structural materials have undesirable nuclear properties, it is vital that new alloys of metals having more favorable nuclear properties be investigated. Columbium and vanadium are both high melting metals, both exhibit resistance to chemical attack, and no great difficulty is encountered in fabricating them into desired shapes. With proper treatment both metals can be cold rolled extensively without failure. In addition they have desirable nuclear properties for certain types of reactors. Therefore, the alloys of columbium and vanadium should be of interest in the atomic energy program. Since an alloy-development program is enhanced by a knowledge of the phase equilibria of the components, this investigation was undertaken to establish the phase diagram for the Cb-V system. According to the Hume-Rothery rules for alloying,&apos; the chemical similarity, crystal structure, and atomic-size factor are favorable for a complete series of solid solutions for this system. Both elements are in the same family of group V of the periodic table and thus are quite similar chemically. The crystal structures of columbium and vanadium are compatible for extensive solid solubility, since both have body-centered-cubic structures. The atomic diameters of columbium and vanadium are 2.85 and 2.62Å, respectively. This difference of slightly more than 8 pct is well within the 15 pct maximum difference allowed for extensive solid solubility. Experimental Procedures Source of Materials: Columbium powder and sheet trimmings were obtained from the Fansteel Metallurgical Corp. According to the manufacturer the metal contains less than 1 pct impurity. An analysis of the metal showed approximately 1800 ppm C in the powder while the sheet trimmings contained less than 500 ppm C. Spectrographic analysis showed minor amounts of Ca, Cr, Fe, Mn, Si, Ti, V, and Zr in both forms of the columbium. No commercial source of vanadium having the ductility and purity desired was available to the authors at the beginning of this investigation. As a result, all of the vanadium used in this study was prepared by the bomb reduction of vanadium pen-toxide with calcium employing the method reported by Long.&apos; Yields of massive vanadium normally were about 80 pct. Chemical analysis of the vanadium prepared in this manner showed the presence of 200 to 500 ppm N and 800 to 1000 ppm C. Minor amounts of Ca, Fe, Mn, Si, Zr, Cr, and Cb were detected by spectrographic analysis. This vanadium metal was ductile and was cold rolled into 5 mil sheet. Annealing was not necessary during this rolling and the metal retained its cold-rolling characteristic after are-melting. Preparation of Alloys: The Cb-V alloys were prepared by melting pieces of vanadium sheet togethel-with columbium in the form of sheet or pellets of powder. The melting was carried out under argon in conventional arc-melting equipment employing a tungsten electrode and a water-cooled copper crucible. Each alloy was remelted three or four times, inverting the alloy after each melting in order to assure complete mixing. Alloys normally were obtained as round flat disks, weighed approximately 70 grams, and had roughly the shape of a disk 1 1/2 in. in diameter and 1/4 in. thick. Half of each alloy

Jan 1, 1955

By R. O. Williams

A recent article by Chernock, Singer, Mueller, and Beck 1 which supports the use of the integral of I sin $ d$ for comparing fiber texture data does not settle what happens to I sin ø as ø approaches zero, where I is the diffraction intensity for poles oriented 0 deg from the fiber axis. They indicate that they believe that it should approach some constant value (that is, I increases without bound) and the fact that in reality it does not was a result of the diffraction

Jan 1, 1960

By John K. Kennedy

ThE basic equations describing continuous zone-melting via the zone-void method were reported by Pfann. Using these equations, a theoretical study was undertaken in this laboratory on the effect of the variation of the significant parameters of the zone-void system. These parameters are the zone length, 1, the heater length, h, and the effective distribution coefficient, k. The object of this study was to determine the effect of these parameters on and (where a = Cp/Cf and ß = Cw/Cf, Cp, C,, and Cf being, respectively, the solute concentration in product, waste, and feed) so that, in future application of this technique to the ultrapurification of materials, the effect of 1 and h would be predictable regardless of the value of k. The basic assumptions made by pfannl were, for the purposes of this study, assumed to hold. For k < 1, the steady-state solute distribution in the stripping section and the shift in this distribution brought about by passage of a void is shown in Fig. 7 of Ref. 1. In this section the definition of k and ß precludes the possibility of ß assuming values less than unity when k < 1. In the course of this investigation it was found that computations based on Pfann&apos;s equations gave values of ß less than unity in the region 1 > k > l/h for almost all values of Ls, where Further computations showed that 6 assumes negative values when k > 1 for many values of L,, 1, and h. A check of the derivation of the stripping-section equations showed that an error had been

Jan 1, 1964

By L. B. Schmitt, N. H. Polakowski

Square bars of cold-dvawn copper, mild steel and a 2024 aluminum alloy were forced through an adjulstable gap between two hardened steel rollers by pulling or pushing. The rollers were either free to rotate or were locked in position so that conditions comparable with drawing, rolling, ad containerless extrusion could be created at will. Push or pull forces as well as roll separatirg forces were measured. Forming efficiency and coefficients of friction were determined. For- slow-rate deformation the values ot µ wrere between 0.09 and 0.14. Possibilities for further investigation are indicated. THIS work was undertaken with the aim of a) determining the principal parameters of cold drawing under simplified conditions directly, without recourse to speculative assumptions and b) comparing drawing with rolling and extrusion from the viewpoint of forces involved and deformation efficiency. Data of type a) relating forces in axially-symmet-rical drawing to die geometry. drafting schedules, material properties and lubrication conditions are abundant. Their intrinsic value is. however. limited as long as the radial die forces are unknown. On the other hand. the difficulties associated with measurement of these forces are serious as witnessed by the extreme paucity of such information."" Comparative investigations of type b) were also remarkably few and they all employed ring-shaped dies through which round bars were either drawn or pushed.2-4 Only in one case were the applied and the die splitting forces measured simultaneously.2 In order to simplify the problem, deformation under approximately biaxial conditions was adopted in the present work;. This enables both the drawing (rolling, extruding) efforts and the separating forces to be measured with reasonable ease by conventional methods. Since drawing, extrusion, and rolling can be carried out under virtually identical conditions only in a roller die. an adjustable tool of this type was constructed. The circular die profile involved a difficulty with interpreting the drawing and extrusion data. This was obviated by arbitrarily defining the angle between a contact chord and the drawing plane as the -.equivalent" die semiangle. The simplification was considered permissible since the effect of die angle was not specifically investigated. An additional conlplication was caused by the square cross section of the processed bars. This shape was selected as a compromise between the desire to have as massive as possible specimens and the capacity of the roller die. The metal was only partially restrained from spreading laterally. the flow conditions being therefore intermediate between plane stress and plane strain. On the Huber-Mises yield criterion the constrained and uni-axial yield strengths differ by up to 16 pct but the actual error due to the ignorance of the constraint factor was presumably much smaller than the maximum. In spite of these reservations it was felt that the method employed and results obtained have some merit. With certain further refinements to be indicated later, this technique can yield data perhaps more useful for checking metal forming theories than other, mostly indirect, methods. EXPERIMENTAL Materials—The metals used together with their principal mechanical properties are listed in Table I.

Jan 1, 1961

By J. Gurland

The effect of four different second-phase additions on the strength of composite alloys with a common matrix was investigated. The four compositions each consisted of 85 pct by volume of silver with additions of 15 pct of either tungsten, molybdenum, tungsten carbide, or nickel. These particular additives were chosen because of their low mutual solid solubilities with silver at room temperature, the alloys consisting essentially of a dispersion strengthened silver matrix in which the second phase is distributed randomly. Test specimens were prepared by a powder metallurgy procedure consisting of ball milling the mixed powders, compacting, and sintering at 920°C for 1 hr in hydrogen. The constituent powders were of CP quality and of particle size between 2 and 10 ?. The compacts were subsequently hot pressed at 1000 psi and 1400° in an Inconel die, coined at room temperature at 60,000 psi and finally annealed and outgassed in vacuum at 800°F. The microstructures of the alloys are shown in Fig. 1, and the microscopically observed porosity grade was A-2 or better as evaluated by ASTM Recommended Practice B 276-54. The properties of the sintered compacts are listed in Table I, and the load-deformation curves in bending are shown in Fig. 2. The samples, of rectangular cross section, were loaded at the center of 58 in. span, and the reported stress was calculated on the basis of linearly elastic behavior in order to allow the comparison of specimens of slightly different dimensions. It is apparent that the tungsten, molybdenum, and tungsten carbide alloys exhibit very similar deformation curves. The moduli of elasticity were determined from the slopes of the stress-deflection curves at small loads. The measured values agree quite well with those calculated by the strength of materials approximation proposed by Paul1 for the elastic constants of mixtures, the calculated values being 15.0, 15.0, 17.4, and 13.4 million psi, respectively, for the tungsten, molybdenum, tungsten carbide, and nickel combinations with silver. The strength of the Ag-Ni composition is appreciably lower than that of the other three alloys. The discrepancy between the two groups is accounted for by the following characteristics of deformation and failure which were observed microscopically during the tests: A) The deformation of the W, Mo, and WC alloys did not involve the second phase particles to any appreciable extent. Slip and fracture took place in the silver-rich matrix alone, although the fracture path generally skirted the particles by favoring the phase boundaries. The deformation proceeded according to the following sequence: 1) formation of small, discontinuous microcracks in areas of the matrix adjacent to a particle, 2) general slip in the matrix and extension of cracks, 3) failure of specimens by propagation of cracks through matrix. The near coincidence of the 3 load-deflection curves suggests that the nature of the second phase particles does not decisively influence the plastic behavior of composite alloys of this composition as long as the particles do not undergo any considerable deformation themselves. B) A markedly different behavior was exhibited by the Ag-Ni composition where the relatively low flow strength is associated with the low yield strength of the nickel grains. Slip in the nickel grains was first observed at a stress of approximately 10,000 psi, although failure, at a later stage, took place by the propagation of cracks through the matrix. The ductility of this aggregate, in which both components deform plastically, is appreciably greater than that of the elastic-plastic composites. In summary: 1) The measured values of the moduli of elasticity of the composite alloys under discussion are in good agreement with the "strength of materials" approximation proposed by Paul.&apos; 2) The overall plastic behavior of these particle-strengthened alloys is unaffected by the nature of the second phase particles, if the latter remain elastic. This work was supported by the Office of Naval Research. The experimental results are presented

Jan 1, 1962

By W. A. Wood, W. H. Reimann, K. R. Sargant

A study is made of the microstructural changes produced in armco and pure iron subjected to alternating torsion at amplitudes above and below the knee of the S/N curve. The aim was to identify the basic mechanisms of fatigue and to find why they should give rise to a sharp S/N knee in contrast to the type of S/N curve shown by fcc metals. It is found that the basic structural change above the knee in the iron is a pronounced cell formation in the grains, which determines irregular micro-cracks, and below the knee a dispersion of fine slip that determines a dispersed pore formation in the grains. Comparison of these changes with those in fcc metals similarly tested shows that they do not include an intermediate structural change, a concentrating of slip in narrow bands with associated slip-band cracking, that characterizes the fcc metals and accounts for their different type of S/N curve. RECENT conferences on basic mechanisms of fatigue emphasize that fatigue damage takes many forms: for example, cracking in slip bands (Thompson1), at grain boundaries and inclusions (Hempel&apos;), at precipitation zones in alloys (Han-stock, Forsythe), or at surface notch-peak effects, both hypothetical (Mott, cottrel16) and observed (Wood). Evidently it is not easy to see in this diversity a basic mechanism. At the same time the contributions suggest that some reconciliation may be possible if distinctions are drawn between three phases of the fatigue process: initiation of damage at nuclei, distribution of these nuclei, and the final macroscopic crack system. In initiation, the relevant features appear to be that nuclei may appear early, at least by 1/100th to 1/1000th of the life of a test specimen; that they may do so at many independent centers, appearing according to some observationsa as pores or voids; and that they are not necessarily self-propagating, damage spreading rather by their multiplication. In distribution, the relevant feature appears to be that they form preferentially in zones where the microstructure becomes abnormally strained (as defined more closely below). Now in these two phases the basic fatigue process should be systematic. But, at the same time, the "abnormal strains" in particular set the stage for considerable diversity in the final phase of macroscopic cracking, for where zones of abnormal strain develop will depend on what kind of fatigue deformation is applied and what kind of prior processing the test material has had. On this view it seems desirable to know what in

Jan 1, 1964

By F. E. Jaumot, R. L. Smith

Self-diffusion in zinc has been used as an instrument for comparison of the absorption and sectioning techniques as a means of studying diffusion. Single crystal as well as poly-crystal samples were used asanda the temperature range of diffusion extended from 200° to 415°C. For temperatures above 200°C, the data indicate that the results obtained from the absorption technique agree with those obtained from the sectioning technique. The effect on the values of the diffusion coefficient of electroplating vs evaporation as a means of applying the tracer was investigated and no significant difference observed. It was found that an excess or deficiency of tracer did not materially affect the results obtained from the sectioning technique, but invariably caused errors with the absorption technique. NEARLY all self-diffusion experiments and many alloy diffusion experiments have employed radioisotopes as the instrument for following the diffusion. Of the various techniques in which radioisotopes are applicable, the two most widely used are the sectioning and absorption techniques. In the former, the diffusion coefficient is obtained from a curve of log concentration vs the square of the mean penetration depth. In the latter, the diffusion coefficient is calculated from a knowledge of the decrease in activity at the surface after diffusion. In spite of the importance of knowing the reliability of the diffusion data and the widespread use of these two techniques, they had not been directly compared experimentally on the same samples. Thus, such a comparison was worthwhile. The specific system chosen for this comparison was the self-diffusion of zinc, since the metal was readily obtainable in high purity, single crystals were easily grown, and the anisotropic features of the diffusion added an additional factor of interest to the comparison. Self-diffusion in both single and polycrystals was investigated, since the possibility existed that grain boundaries and grain boundary diffusion might produce different effects for the different techniques. Also, the effects of electroplating vs evaporation as a means of applying the radioactive tracer, and the effects of applying an excess or deficiency of tracer material were studied. The general experimental procedures and the mathematical analyses of both techniques have been described previously in some detail, and only those points essential to an understanding of the present work will be discussed here. The mathematical analysis of the sectioning data requires only a knowledge of the solution of Fick&apos;s equation. where C,, is the initial concentration, C is the concentration at depth x, D is the diffusion coefficient, and p is the absorption coefficient of the solvent metal for the radiation from the solute. Having obtained F from the experiment, the value of z can be read from a plot of F vs In z, and D can then be calculated if is known. The necessity of knowing the absorption coefficient causes most of the criticism of the absorption technique. It is very difficult to measure accurately under the conditions existing in the diffusion experiment, but it is absolutely necessary, as will be indicated briefly later. Portions of this problem have been discussed in detail elsewhere." The difficulties involved in the measurement of the absorption coefficient arise from three principle factors: the physical characteristics of the absorber, the geometry, and complex radiation spectra. The first two are heavily dependent upon the third. If

Jan 1, 1957

By A. G. Rozner

TRANSVERSE rupture tests have been commonly used in mechanical investigations of brittle materials. The specimens are simple, easy to prepare, and loading presents no difficulty. Owing to the complexity of stress distribution in bending1 one can determine stresses rigorously only for the elastic range. The bend-test results, however, are at times converted into tensile-strength data, determined as maximum fiber stress. The assumption is made that for small plastic deformations the linear stress-strain relation applies. This problem was experimentally investigated by Bach3 who tested cast iron in bending and in tension and found that the maximum fiber strength obtained in bending exceeded the experimental tensile strength by a factor depending on the geometry of the tested piece. The factor varied from 1.57 for a rectangular cross section to 2.12 for a circular cross section of the tested beam. Since this work was carried out 60 years ago, the cast iron used could be expected to have been highly brittle, thus the extent of plastic deformation prior to fracture very small. We have therefore carried out a number of tests on magnesium-modified cast iron with an average elongation of -5 pct. It was hoped that the data obtained from the tests would serve as a guideline for comparing data obtained in tension and bending for brittle materials. Mechanical properties were measured at room temperature on an Instron tensile machine at a crosshead speed of 0.05 in. per min. The three-point bend test employed cylindrical bars 0.20 in. in diameter with a 1.0-in. span. For the tensile test, cylindrical tensile specimens with a gage length of 0.85 in. and a diameter of 0.160 in. were used. Results are summarized in Table I. From Table I one can see that fracture strength

Jan 1, 1965

By A. Cochardt, D. H. Feisel

HE success of the sintered aluminum powder material (SAP)&apos; in raising the maximum useful temperature of aluminum alloys has stimulated effort in investigating similar systems in which the metallic matrixes have higher melting points. Examples of such systems which have been studied are nickel-alumina,2, 3copper-silica,4 copper-alumina,4 titanium-cerium oxide,5 and 80/20 nickel-chromium,8 cobalt,= and molybdenum7 matrixes with various hardeners. Conventional precipitation-hardened high-temperature alloys are often thermally unstable in service. The original fine particles that are responsible for good strength properties grow in size at elevated operating temperatures, a process termed "over-aging." This growth of precipitated particles causes a gradual loss in strength properties. A previous paperb describes the factors that govern the process of overaging, and estimates the time required for overaging of a nickel-alloy matrix dispersion-hardened by various precipitates such as are found in current high-temperature alloys, and by alumina. These calculations indicate that nickel-base alloys hardened with Ni3Al, Tic, or TiN are quite unstable at 1340" F (727°F). Such alloys should overage within a few hours at that temperature. A nickel-base alloy hardened with alumina, on the other hand, should behave differently. It should not overage appreciably within a thousand years. It was further shown8 that the thermal stability of the oxide-hardened alloy is mainly due to the low solubility of oxygen in equilibrium with alumina in the nickel-base matrix. Since the oxygen solubility is small, little oxygen can diffuse, and structural changes require very long times. This paper illustrates the effects of time and temperature on the structural stability of nickel alloy matrixes hardened by dispersions of Ni3A1, Ni3Ti, Tic, and alumina, and of nickel-chromium matrixes hardened by alumina. Hardness measurements and metallography have been used as a measure of structural stability. EXPERIMENTAL PROCEDURE Using the best information from published phase diagrams,&apos;-l4 compositions of the Ni-Al, Ni-Ti, and Ni-Ti-C alloys were determined so as to have 5 vol pct of precipitate at 1350° F(732°C), the temperature selected to illustrate the effect of time on structural stability. Table I lists the chemical compositions of these alloys. Standard laboratory melting and alloying techniques were used in preparing ingots of the F-Series of alloys. The cast ingots, 1 3/4 in. in diam, were forged to 3/8-in. diam rods.

Jan 1, 1960