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By R. F. Hehemann, D. J. Cometto, G. L. Houze

The w transformation in the Zr-Nb system was studied using X-ray diffraction, dilatometric, re-sistornetric, hardness, and metallographic techniques. w forms in a diffusionless, completely reaersible manner on quenching and in a diffusion-controlled manner- on aging. The temperature at which w begins to form on quenching was determined as a Junction of composition and was found to decrease with increasing solute content. w formed bv aging establishes a metastable equilibrium with an enriched ß phase. The ß/w + ß transus has been determined for this metastable equilibrium and employed to rationalize retrogression and reversion phenomena observed in these alloys. The decomposition mechanism is discussed in terms of a gradual or continuous transformation from ß to the w state. BETA- stabilizing alloying elements lower the MS temperature of the martensitic bcc (ß) to hcp (a') transformation in zirconium and titanium alloys. In certain titanium alloys, a lower symmetry modification of the martensitic structure, termed (a"), also has been reported.1,2 These martensitic transformations are suppressed when the amount of the ß-stabilizing element exceeds a critical level. However, the high-temperature ß phase cannot be quenched to room temperature without change. At alloy contents just above the critical level, |ß trans-forms to the w phase when cooled rapidly.3-6 The amount of w in quenched alloys is reduced by increasing alloy content, and this phase virtually disappears above a critical composition.5 In addition to the transformation during cooling, w also can be formed by aging ß at temperatures below approximately 500°C, and significant alloy enrichment of retained 13 accompanies the isothermal transformation.1-8 The structure of LC is closely related to that of the ß from which it forms.9-15 The bcc (ß) lattice can be generated using a hexagonal cell with three atoms located at (000) and ±(2/3, 1/3, 1/3). This cell has an axial ratio of 0.612 and is oriented with respect to the cubic cell such that (0001)H (111)C and [1120]H [101]C. Consequently, there are four possible orientations of the hexagonal cell, associated with the four (111) planes of the bee lattice. Formally w can be obtained from 0 by allowing the two atoms at the ±(2/3, 1/3, 1/3) positions to approach the coordinates ±(2/3, 1/3, 1/2) and there are four equally probable orientations of w for each 0 grain. In titanium alloys w retains the cubic axial ratio of 0.612" and hence also can be indexed as triply cubic, but this is not the case for aged Zr-Nb alloys where w is clearly hexagonal with an axial ratio of 0.622." The lattice parameters and atomic positions of w depend on thermal history and alloy content. The atomic positions (000), ±(2/3, 1/3, 1/2) provide reasonable agreement between calculated and observed diffracted intensities for w in the fully aged condition.10,14 In the quenched condition, on the other hand, the atoms appear to be displaced from the central plane12,13,15 and assume positions ±(2/3, 1/3, Z = 0.42-0.48) rather than ±(2/3, 1/3, 1/2). This results in a structure with trigonal rather than hexagonal symmetry. The readily detectable parameter and axial-ratio changes of w in Zr-Nb alloys make this system especially attractive for studying the structural changes that occur in the formation and aging of w. Particular attention in this investigation has been directed to the relationship between the structure of w in the quenched and aged conditions, and to certain aspects of the reaction kinetics. MATERIALS AND PROCEDURE Zr-5, 12, 17.5, and 25 pct Nb alloys were prepared by a double arc-melting practice, encased in stainless-steel cans and hot-rolled to 1/8-in.-thick sheet.* Charged weights have been employed for the niobium contents and interstitial analyses are reported in Table I. Dilatometric, X-ray diffraction (filtered CuK, and MoK, radiation), metallographic, and hardness techniques have been employed to follow the transformations during isothermal and quench-aging heat treatments. In quenched specimens, the electrical resistivity also was studied. Betatizing was conducted for 1 hr at 900°C. With the exception of the isothermal dilatometric studies, samples were

Jan 1, 1965

By Paul A. Beck, M. N. Parthasarathi

The vecrystallization texture. formed by selective growth of random nuclei in an 80 pct volled 99.997 pct A1 crystal of initial orientation mear (123) (41.21 was found to consist of components related to the single-component rolling texture by 40 deg. rota-tiows around [111]axes. The relative-volume fractiovz of the various recrystallization-texture componezts may be accouzted for 012 the basis of the Lou) mobility of twist boundaries, if it is assumed that the mobility of the tilt boundaries also varies with their orientation. Similar results with a, aluminum-alloy cystal of the (110) [112] orientation, containing iron, silicon, and zinc, show that these solute contents do not impair the effectivezess of the oriented growth mechanism of re cry s-tallization- texture formation. In a previous investigation1,2 it was found that a sharp recrystallization texture is formed as a result of the oriented growth of random artificial nuclei in an aluminum crystal rolled to 80 pct RA, on the (110) plane in the [ 112] direction. Four recrystallization-texture components were found to occur, with orientations corresponding to 40-deg rotations around [111 1 axes of the deformation texture. Reorientations of this type were already found in earlier work,3,4 where the presence of random nuclei could be reasonably assumed, but was not rigorously documented. The same orientation relationship was found also by Lucke and Liebmann, who actually measured the orientation dependence of boundary-migration rates in the recrystallization of A1 crystals slightly deformed in tension.5,6 The occurrence in the recrystallization texture of only four of the eight crystallographically equivalent orientations of the type referred to above remained unexplained,1,2 although it was clearly not a result of the absence of suitable nuclei. Recently Burgers7 suggested that this effect may be due to the aniso-tropy of boundary mobility observed2'6; it was found that the recrystallized grains of the 40-deg [111]-rotated type grow much faster in directions perpendicular to the [ 111] they have in common with the matrix than parallel to this direction, i.e., the tilt boundaries have much higher mobility than the twist boundaries. The four orientations corresponding to

Jan 1, 1962

By J. J. Kramer, K. Foster

The orientation distribution of surface-energy -induced secondary recrystallized grains was determined. This work was conducted on thin sheets of a 3 pct Si-Fe alloy annealed under environmental conditions that furor grouth of grains with a (100) plane in the surface of the sheet. The texture was found to be extremely sharp and almost independent of sheet thickness. The distribution varied exponentially with the angular deviation from the {100} plane. It was possible to relate the distribution to the nu-cleation rate of the secondary rains as influenced by the surface-energy difference. THE role of surface energy in the secondary grain growth of cube-oriented grains (grains with a (100) plane in the plane of the sheet) in thin Si-Fe sheets has been previously discussed.1-4 In high-purity sheet material normal grain growth usually occurs until the grains have extended through the sheet. Further grain growth is inhibited by the thermal grooving of the boundaries at the sheet surface. However, additional growth of cube grains can occur by a secondary grain growth process under conditions where the (100) plane has a lower surface energy than other orientations. Apparently for these alloys, cusps exist in the polar plot5 of surface free energy with the lowest cusp energy occurring at the (100) orientations. This has been reported to be the result of preferential adsorption of sulfur on the (100) planes.6 As a result of this process, a distribution of orientations could arise from two possible mechanisms. First, when a cusp is present in the polar plot of surface free energy, there are orientations inside the cusp that have a lower surface energy than elsewhere on the polar plot. Also, at sufficiently high temperatures, flat surfaces whose orientations are inside or just outside the cusp (depending on its shape) can often thermally etch, yielding a microscopically stepped surface of even lower surface energy. As a result, grains oriented close to cube would also have a lower surface free energy, either because of the cusp shape or by thermal etching, and could possibly grow as secondary grains by the surface-energy phenomenon. One should thus observe a distribution in the surface orientation of the cube grains comprising the secondary structure. It is the purpose of this paper to investigate this orientation distribution experimentally and to discuss the factors involved in its formation. For this purpose, the surface orientations of a large number of secondary grains in various sheet thickness were determined by means of the Laue back-reflection X-ray technique. PROCEDURE A vacuum-melted 3 pct Si-Fe alloy containing a nominal impurity content of 0.005 wt pct was processed into strip. A single cold-rolling step of 90 pct reduction was used for each strip regardless of the final sheet thickness. Final strip thicknesses of 0.60, 0.30, 0.15, and 0.075 mm were used. Care was taken to insure that the final strip surface was smooth and flat. All strips of a given thickness were annealed together at 1200°C in dry hydrogen (dew point -70°C) to develop the desired secondary structure and to insure identical environmental annealing conditions. The annealing time was selected to develop a complete secondary structure in the thinner sheets but to permit the thicker sheets (0.60 mm) to have residual primary grains remaining. This was necessary to determine whether growth impingement could lead to one secondary grain consuming another at a greater angular deviation. For the X-ray determination of the surface orientation of the secondary grains, a special specimen holder was used. The camera and holder arrangement could be aligned by X-raying a grain in three positions rotated 180 deg to each other. Thus, with a small beam X-ray focus (1 mm), the surface orientation of any grain could be determined to within one-half a degree. The surface orientations of one hundred cube secondary grains were determined for each sheet thickness. The criteria of a secondary grain were its size relative to the sheet thickness and the number of sides of the grain observed in the sheet surface. (A primary recrystallized grain extending through a sheet will generally have six edges visible in the plane of the sheet, whereas a secondary grain will have many more when growing entirely into primary grains.) Grains were selected as randomly as possible by X-raying every secondary grain found along a line drawn on the strip. No attempt was made to determine the exact orientation of the planes of the surface, as many strips from randomly selected sheets were used. On1y the angular deviation of the surface plane from {100} was measured. In order to assess the volume distribution in the

Jan 1, 1965

By R. Maddin, C. H. Mathewson, W. R. Hibbard

According to conventional crystal mechanics, a face-centered cubic single crystal slips on the system {111} <110> which receives the highest shear stress in terms of the stated orientation of the stress axis of the crystal.l However, in brass, the slip process is never as simple as this conventional analysis indicates. There is always a second cooperating plane which contains the same slip direction as the primary plane.2&apos;3 There is still another cooperating slip system commonly called the conjugate system3 previously observed to function with considerable lag only after the rotation due to the action of the first plane had produced a symmetrical condition of equal shear stress on the two systems. In the 1943 Campbell Lecture,4 a slip mechanism was outlined in which the {111) planes moved step-wise in two adjacent <112> directions integrating into the <110> slip direction. "... the slipping process would be visualized as a succession of steps 30" to the right and left of the resultant, microscopically observed, slip direction; a first step setting up the twinned configuration in the single spacing involved and an ensuing slip returning the atoms to the untwinned configuration while completing a fully integrated <110> movement."4 This process is shown diagrammatically in Fig 1 which shows two close packed layers of atoms at a slip site. A movement of the full circles through the "valleys" into the adjacent sites would produce a twin configuration with respect to the open circles. The term, "twin fault," seems most appropriate to define this faulted condition of the lattice. Annealing twins in face-centered cubic metals may arise from preformed nucleii which could be described as twin faults produced by the first movement of this two stage slip process. However, it is clear that this kind of slip does not take place so as to embrace a measurable number of adjacent planes because careful X ray examination has not conclusively shown the existence of ponderable twin lamellae (as in the case of Neumann bands in the body-centered ferrite). In view of the imponderable character of these hypothetical twin faults, the search for evidence of their existence seems restricted to indirect methods. The most obvious experimental technique serving this purpose is the determination by micrographic and X ray methods of the alignment of the composition planes of annealing twins with the former slip planes in axially strained and recrystallized single crystals. In previous attempts of this sort not all of the annealing twins could be traced to slip planes known to be operative in the conventional sense and hence evidence appeared negative or inconclusive. With the discovery of additional slip systems2&apos;3 operating in conjunction with&apos; the major system of highest resolved shear stress in axially strained brass it became apparent that the simple technique outlined above might be re-employed with more gratifying results. If, on annealing single crystals of alpha brass previously strained in tension, twins are found with their composition planes parallel to all three operative slip planes and with no participation of the fourth plane there is good evidence in favor of the existence of twin faults and of the two stage slip process. Experimental Procedure and Resuits Single crystals of alpha brass of the 70-30 composition were grown by solidification from the molten state.

Jan 1, 1950

By R. D. Reiswig, J. M. Dickinson

The 0s-Ir equilibrium diagram was determined. The diagram is of the simple peritectic type, with a peritectic temperature of about 2660°C. The solid miscibility gap is narrower than previously reported, the nearly vertical solvus lines occurring at about 42 and 63 wt pct 0s. Widmanstaetten figures resembling martensite in some of the two-phase alloys account for a previously suspected transformation. In view of the long history of use of such alloys, it is surprising that the 0s-Ir equilibrium diagram has not yet been completely determined. The boundaries of the miscibility gap in this system have been estimated to be between 24 and 39 wt pct 0s and 79 and 100 wt pct OS. It has been predicted3 that alloys of ruthenium, osmium, and rhenium with the fcc platinum metals should solidify peritectically in their two-phase regions. The metallographic examination of natural osmiridium alloys has been said4 to indicate, by the presence of a martensitic structure, the existence of a transformation or a phase separation. On the basis of superconducting transition temperature measurements, a 70 at. pct 0s alloy was found5 to be two-phase prior to annealing. EXPERIMENTAL The elemental components used in this work were obtained from several suppliers as sponge. The purities, stated by each supplier to be in excess of 99.8 pct, were confirmed in each case by a semiquantitative spectrographic analysis. The pure components were premelted in an inert-gas tungsten arc furnace with the evolution of some fume, crushed, and pickled for use as alloy melting stock. Upon arc melting to the final compositions, no fume evolution was observed. Because weight losses of the order of 15 mg or less were observed in preparing alloy buttons weighing 15 g, "synthetic analyses" were used in determining the compositions of the alloys. Alloys used in this investigation contained 15, 20, 33, 40, 42, 43, 45, 50, 55, 60, 63, 64, 66, 70, 79, and 80 wt pct 0s. Heat treatments of the alloy buttons and pieces thereof were performed in a previously described6 tungsten tube furnace under pressures of 105 torr or less. It was found that homogenization treatments on the specimens were essential to remove the severe coring which was present, a 3-hr treatment above 2300°C removing all the metallographi-cally visible evidence of coring. Following the various heat treatments, the specimens were "vacuum-quenched" by melting a tungsten fuse wire and allowing them to drop onto a water-cooled copper plate. In a few instances, specimens were quenched into molten tin for comparison, but no difference was apparent. Vacuum-quenched specimens cooled from 2400°C to below a red heat in less than 60 sec. Typical equilibration times varied from 3 hr above 2300°C to 88 hr near 1100°C. The utilization of successively longer equilibration times and the approach to equilibrium from both directions indicated that the times used were sufficient to give a close approach to true equilibrium. Heat-treatment temperatures were measured with a calibrated Leeds and Northrup optical pyrometer through a 0.120-in. lateral sight hole in the tungsten heater tube. It was known that the assumption of black-body conditions inside the 1-in.-diam by 10-in.-long tube was unwarranted. The empirical results of work with thermocouples and melting

Jan 1, 1964

By E. W. Murbach

The oxidation of uranium carbide by oxygen at various pressures, and by air, has been investigated at temperatures up to 600°C. Arc-melted and cast uranium carbide displays oxidation behavior that appears to be dependent on its treatment after melting. ccReactive" uranium carbide, which results from exposure of freshly melted material to laboratory air, ignites in oxygen at about 275° to 350°C depending on pressure, and in air at about 350°C. The product of the initial oxidation appears to be UO2+.. Continued heating produces U30,. The initial oxidation rate is poportional, roughly, to the square root of the oxygen pressure. The activation energy determined for the initial oxidation reaction is 7.1 kcal, from 400° to 600°C in 10 mm of oxygen. AN investigation of low-decontamination methods for reprocessing uranium carbide reactor fuels is being carried out at Atomics International. One of the reprocessing cycles includes oxidation of the fuel at relatively low temperature (300° to 500°C) and reaction of the product oxide with carbon at elevated temperatures (above 1500°C) to produce uranium carbide. A number of small-scale experiments have been carried out in which uranium carbide was oxidized in air and oxygen at various pressures to measure reaction rates and to obtain other information concerning the oxidation process. In initial experiments with arc-melted UC, ignition was observed at about 275°C in oxygen or 350°C in air. However, similar experiments by Neymark&apos; at 426°C in air showed no ignition. It was soon determined that the difference in oxidation behavior was due to the treatment of the UC following arc-melting. Neymark&apos;s experiments were carried out with freshly melted UC while those reported in the

Jan 1, 1963

By W. C. Hagel

In 1952, Hauffe and pfeiffer1,2 reported that parabolic rates of nickel oxidation are decreased by factors of two and four at 1100° and 1000°C, respectively, when Li2O vapor is present. Hauffe &apos;s3 explanation was that monovalent lithium enters the crystal lattice of p-type NiO to decrease cation-vacancy concentration and, hence, diffusion of Ni2+ ions across the growing oxide. After first succeeding in adding lithium directly to molten nickel, the purpose of this investigation was to determine whether alloyed lithium is a more effective oxidation inhibitor than the external vapor. Seven 650-g heats were prepared by melting carbonyl nickel inductively under 1 atm of argon. Various amounts of lithium were added by dropping small chunks into the melt. Only about 10-pct recovery was achieved with the remainder lost by spattering and volatilization. Nominal and actual

Jan 1, 1965

By S. T. Wlodek

The scale md subscale reaction products were identified and their rates of formation were studied in air over the range 1600" to 2000°F (871 " to 1149°C) for periods of up to 400 hr and for hoth the solution-annealed and aged conditions. The effect of prior sltrface preparation on suhscale oxidation was also studied The general oxidation behavior of both Ni-Cr-Mo-Al-Ti type alloys was similar. A surface film of a, Al2O3, forms immediately on exposure Subsequent oxidation continued at a linear rate (QL = 55 * 5 kcal per mole) as colonies of Cr2O3 nucleated at the A12O3/gas interface Further oxidation proceeded at a paraholic rate zvhiclz could he fitted to two successive rate constants. During paraholic oxidation, and depending on temperature, the scale consisted of Cr2O3, NiCr2O4 , and TiO2 with traces of NiO. In the case of Rene 41, the activation energy of both paraholic processes was 66 * 3 kcal per mole suggesting that diffusion of cations tlzrozrgh Cr2O3 was the rate -determining process. An unusual decrease in the oxidation of Udimet 700 at 1900°F where a spinel of Ni(A1,Crh0, zuas the predominant reaction product prevented the accurate assignment of activation energies for this composition. In both alloys internal oxidation of Al2O3 commenced shortly after parabolic scaling was observed. Prolonged exposure prod7tced intemal oxidation of TiN, and in Udimet 700 a complex Mo-Ni nitride was also found. At 1900oF, the subscale reactions in Udimet 700 undergo an inversion which parallels the decrease in surface oxidation; internal oxidation ceases hut is replaced by the formation of "spherodized" 3.&apos; colonies. Surface-preparation techniqtles which introduce appreciable working, such as coarse surface grinding or grit blasting. increase the amount of alloy depletion and internal oxidation in Reni 41. The reverse is true of Udimet 700 for which electropolished or mechanically polished specimens show much more subscale oxidation than strongly worked stirfaces. The strongest commercial nickel-base alloys presently available are generic to the Ni-Cr-Mo-A1-Ti base which exploits the precipitation of Ni3(A1,Ti) as the main strengthening mechanism, while relying on solid-solution strengthening by molybdenum and chromium reinforced by the pre- cipitation of carbides to attain maximum properties. This study characterizes the oxidation behavior of Rene 41, the strongest alloy of the Ni-Cr-A1-Ti type commercially available in sheet form, and Udimet 700, whose higher aluminum and titanium content allows it to exhibit one of the more attractive combinations of high-temperature properties available in a wrought product. The scaling processes of complex, type nickel-base alloys have received relatively little attention. Malamand and vidal as well as Poulignier et al.2&apos;3 have determined the composition gradients across the metal/oxide interface produced by high-temperature oxidation and considered the effect of surface perature, Limited weight-gain data has also been published by Fere 5 for alloys of this type and Radavich6 has identified the reaction products on Udimet 500 and Inco 702 after oxidation at 1832°F. Reference can, of course, be made to the excellent reviews of Kubaschewski and Hopkins7 or Ignatov and Shamgunova8 for a summary of the data available on the oxidation of binary and ternary alloy systems which are related to the more complex alloys considered here. EXPERIMENTAL The analyses of the different commercial heats studied are given in Table I. Using the experimental procedures previously established,9 continuous weight-gain data were obtained on both heats of Rene 41 sheet (A and B) and 150-mil-thick slices of cast Udimet 700. Subscale oxidation reactions were followed by static exposure of cylindrical specimens obtained from swaged Rene 41 (Heat C) and Udimet 700 (Heats E and F). In brief, continuous weight-gain tests were performed on specimens with a surface area of 10 to 12 sq cm. These were abraded through 600 grit Sic paper, electropolished to 2p rms in an electrolyte of 10 pct H2So4 in ethanol, and lightly etched in 10 pct HCl in ethanol before final washing and rinsing in ethanol. All continuous weight-gain data were obtained in dried (-70°F dew point) flowing (1 liter per min) air to an accuracy of +0.1 mg. Subscale oxidation processes were followed by the metallographic examination of 0.5-in-diam by 1.0-in.-long specimens. After an initial center-less grinding, various additional surface treatments were employed to determine the effect of surface preparation on subscale oxidation processes. Before exposure in zirconia crucibles, all samples were lightly etched in 10 pct HCl-ethanol, washed in ethanol, and dried. The depth of internal oxidation was measured to ±0.00025 in. on unetched specimens mounted so as to provide a taper mag-

Jan 1, 1964

By G. S. Farnham, D. A. Scott

Partition of certain elements, particularly nickel, was determined for slowly cooled steels, the greater number containing from 0.30 to 0.35 pct C. Approximately 3 pct of the nickel, 18 pct of the manganese, 33 pct of the molybdenum, and 35 pct of the chromium occur in the carbide phase. Partition of one element is not much affected by the presence of another. THE object of this investigation was: 1—to determine the partition of nickel between cemen-tite and ferrite in a medium and high carbon series of annealed steels; 2—to evaluate the effect of chromium and/or molybdenum or both on the partition of nickel in medium carbon annealed steels; 3— to determine the partition of chromium between ferrite and carbide in a series of medium carbon annealed steels; and 4—to study the chemistry of the carbides formed in an AISI 2330 steel during both isothermal transformation, and quench and draw experiments at 1150°F (620°C). The partition of alloying elements between ce-mentite and ferrite in annealed steels has been the subject of many investigations, a summary of which has been given by Austin1. However, there is no satisfactory data available concerning the partition of nickel. It has generally been assumed that virtu- ally all of the nickel concentrates in the ferrite phase in annealed steels. Two investigators"." have reported the finding of nickel in extracted carbide residues. Recently, it has been shown4 that nickel substitutes up to 40 pct for iron in the lattice of cementite produced by the low temperature carburization of iron-nickel powders. In addition to the lack of data on nickel partition, the literature also fails to make reference to the effect of other elements on the nickel content of the carbide in nickel-bearing steels. The partition of molybdenum between ferrite and carbide has been reported"-&apos; for hypoeutectoid steels, isothermally transformed at 1300°F (705°C) and

Jan 1, 1954

By D. W. Rudd, D. W. Vose, J. B. Vetrano

In an earlier paper the permeability character of Mo-0.5 pct Ti to hydrogen was described.&apos; It was shown that this alloy is a more effective barrier to the passage of hydrogen than previously studied systems, i.e., copper,2 nickel,3 and several stainless steels.4 The permeability of Hastelloy B to hydrogen was investigated as a material having promising high-temperature structural properties which might have a low permeability to hydrogen. The general theory of permeation has been previously discussed,&apos; and here only a summary of the more important results are given. It is found that except for very thin membranes, at low temperature, permeation of a gas through a metal barrier is determined by the rate of diffusion of the gas through the crystal lattice of the specimen, and that for hydrogen and the common diatomic gases, this passage involves dissociated particles as the diffusing species. The rate of permeation for hydrogen is therefore proportional to the square root of the pressure. It has also been found that the rate of permeation is inversely proportional to the thickness of the barrier. The apparatus used in this investigation was fundamentally the same as that employed in previous studies,&apos; except that several simplifications result in the case of Hastelloy, since oxidation at the temperatures investigated is no longer a problem and welding of the specimen is relatively easy. A cross section of the membrane assembly is shown in Fig. 1. When finished, the assembly is about 5 in. long. The wall thickness is about 6 times the membrane thickness. The membrane (A) can be maintained at

Jan 1, 1963

By D. E. Etter, J. E. Selle

The Pu-Ce phase diagram was determined by differential thermal analysis, metallography, and elechm-nricroprobe analysis. The dingram is chararterized by a eutectic with extensive solid solubility in the cubic Phases of both cerium and plutonium. No intermediate phases were found. Cerium is soluble in e plutonium to 16.0 at. pet Ce at 626OC nnd in d plcltoniurn to 23.0 at. pet at 613 "C, decreasing to 17.5 at. pet at room temperature. No solubility in the a, ß, or ? phases was jound, and the mininum cerium required to stabilize 5 plutonium to room temperature is 3.2 at. pet. The maximum solubility of plutonium in 6 cerium is 19.5 at. pet at 721OC, and the maximum soluhility of plutonium in ? cerium is 34.5 at. pet. Invariant points in the system are: a) eutectic at 626oC, L16.5 = E Pu16.0 + y Ce65.5; b) peritectoid at613oC, e Pu13.5 -y Ce65.5- d Pu23.0. AS part of a program aimed at the determination of the phase equilibria in the Pu-Ce-Cu ternary system, preliminary investigations were made of the three binary systems. A survey of the Pu-Ce binary system by differential therma1 analysis indi- cated several areas of disagreement with the published diagram,&apos; primarily in the shape of the liquidus surface. Consequently, a thorough investigation of the system was undertaken, and this report presents the results of the study. EXPERIMENTAL PROCEDURE Three methods of investigation were used: differential thermal analysis (DTA), metallography, and analysis with electron-microprobe X-ray analyzer (EMX). The EMX analysis was particularly useful for the determination of solvus lines. The DTA method was most useful for the determination of the liquidus curves. Metallography complemented both of these techniques. Materials. The cerium metal used for most of this work had a freezing point of 805°C. This material showed 15°C premelting, and was found metallographically to contain some unidentified in~lusions.2 The temperature of premelting is defined as the point at which the DTA curve breaks away from the base line, analogous to a curve obtained from a solid-solution alloy. Several compositions were prepared near the cerium end of the diagram with higher-purity material which had a freezing point of 800°C. etallographic analysis showed this material to be relatively free of inclusions and only 8°C premelting was observed. Major impurities were found by spectrographic

Jan 1, 1964

By H. L. B. Gould, Jr. Wenny D. H.

TO date there has been but limited interest in alloys of 80 to 95 pct Co and Fe with or without other additions. In 1932, S. R. Williams&apos; reported practically zero magnetostriction for the 90 pct Co-10 pct Fe alloy. This same composition was used for pole pieces to control the magnetic field in magnetrons operating at high temperatures.&apos; In 1947, a 90 pct Co, 9.5 pct Fe, and 0.5 pct Mn composition was investigated3 for use in ballast lamps because of its high temperature coefficient of resistance. Recently, as part of a survey of ductile and malleable magnetic alloys capable of being processed to thin gages for possible applications in memory devices, the high Co-Fe alloys have been re-examined to determine their fabricating and magnetic characteristics. Alloys from 80 up to 95 pct Co and Fe with and without additions such as vanadium were prepared. In general these all had excellent ductility and malleability. Wire was drawn to thin gages without difficulty and fine hard-drawn wires flattened to 2) The phenomenon may have potential in estimating the carbon content of molybdenum. A given solution-treating temperature and cooling rate could perhaps be calibrated so that the area fraction of surface covered by carbide could be interpreted in terms of carbon concentration. By segregating to the surface, a big "magnification" of the carbon content is obtained, which should make the method cheap, direct, and sensitive and, therefore, perhaps of interest to the molybdenum industry. This work was conducted at Battelle Memorial Institute. Financial support by the Aeronautical Systems Division, United States Air Force, under Contract No. AF 33(657)-7376, is gratefully acknowledged, and the writer is thankful for the experimental assistance of Miss Marjorie Cantin. ribbon without splitting. Similarly, strip was cold-rolled as much as 99.75 pct reduction to a 0.25 mil thickness without intermediate anneals. Samples of round and flat wire and foil, in the severely cold-worked and annealed states, were checked for magnetic characteristics. The alloys exhibited widely divergent magnetic characteristics depending on the processing and heat-treating procedures used in preparing the samples. Direct-current hysteresis loops showing the different magnetic characteristics of 90 pct Co, 9.5 pct Fe, and 0.5 pct Mn alloys are shown in Figs. 1 to 8 inclusive, while Figs. 9 and 10 show the loops for an alloy of 82 pct Co, 11.75 pct Fe, 5.6 pct V, and 0.65 pct Mn. Figs. 1 and 2 show the high remanence and exceedingly square-loop characteristics of the hard-drawn wire and ribbon roll flattened therefrom. With the change in processing from roll flattening to flattening by drawing through flat diamond dies the resultant change in texture is accompanied by the drastic change in the shape of the loop as shown in Fig. 3. This skewed loop on annealing reverts to a more normal loop as shown in Fig. 8. The most unique loops are obtained with flat rolled strip. These are shown in Figs. 4 and 5. In the case of the hard-rolled material, Fig. 4, there is little change in induction until the magnetizing force reaches 50 oe. From 50 to 60 oe there is a sharp increase in induction to over 15,000 gauss, which for all practical purposes is saturation. The induction remains at a constant value while the magnetizing force decreases to 40 oe. At this point the

Jan 1, 1965

By C. G. Rhodes, D. Kramer

The kinetics of precipitation of transient a phase during the transformation in U-16 wt pct Mo between 400° and 550°C were studied using quantitative metallography and electrica1 resistiuity measurements. The volume percent of a phase goes through a maximum in time, corresponding to a critical degree of progress of y&apos; formation. prior to its re-solution after prolonged annealing. The fastest rate of a precipitation oc-cuvs at 500°C with a concomittant lowest peak of 9.5 vol pct a. At 450°C, 14 vol pct a, greatest fov any temperature, was observed after 24 hr witlz an associated slowest rate of formation. a precipitation at 400°C ceases after approximately 160 hr and the amount of a present remains unchanged after long times at temperature; no y&apos; was observed metallographically at this temperature. IN the U-Mo system at 16 wt pct Mo, an ordering reaction takes place below 600°C, see Fig. 1, whereupon the random bcc solid solution, y, transforms to the ordered bct phase, y&apos;."2 The new unit cell is constructed of three old unit cells somewhat distorted. While attempting to study the kinetics of this transformation it was observed that an unexpected phase was precipitating out early during the transformation anneal reaching appreciable amounts before disappearing after long times. The y&apos; appeared in the presence of this new phase and after long times constituted the whole sample. Owing to the high optical activity of this unexpected phase under polarized light, it was suspected to be the a phase. Subsequent X-ray powder pictures confirmed this belief. The kinetics of a phase formation and disappearance is the subject of this paper. EXPERIMENTAL PROCEDURE A) Alloy Preparation. Alloys for this study were prepared by vacuum induction melting normal reactor grade uranium and 99.5 wt pct Mo pellets in thoria crucibles. Two ingots were prepared, and sections of the top and bottom of each ingot were chemically analyzed for uranium and molybdenum content. The results of the analyses appear in Table I. B) Metallography. Samples of 16M were wrapped in tantalum foil, sealed in evacuated (5 X 10"6 mm of Hg) quartz capisules, and annealed at 1110"C prior to being transferred directly to the transformation temperature. They were examined metallographically after increasing times to determine the progress of the transformation at 400°, 450°, 500°, and 550°C. The samples were prepared for metallographic examination by mechanical polishing with diamond paste and electrolytic polishing with a solution of glacial acetic acid and chromic acid. Since the y&apos; phase is tetragonal, its presence can be detected by viewing the sample in polarized light. To increase the contrast between y&apos; plates, the samples were electro-etched in a 10 pct oxalic acid solution. To reveal a phase in bright field (it can be seen in polarized light in the "as-polished" state), the samples were electro-etched in a solution of citric acid and nitric acid. To determine the vol pct a phase, a variation of the one dimensional systematic point count method as described by Hilliard and cahn3 was used. The

Jan 1, 1962

By T. Vasilos, J. T. Smith

THE mechanism of pressure sintering, or hot pressing, for ceramic materials, has been investigated by several researchers.1-8 Plastic flow has been suggested as the rate-determining mech-anism1,2 while grain boundary sliding3,4 and fragmentation processess have been found controlling in other studies. Vasilos and Spriggs5,6 have reported that pressure sintering occurs by a diffusion-controlled mechanism. For the pressure sintering of metals, investigations with lead7 and with tin6 have not identified the operative mechanism during consolidation of these powders. This note reports the results of a pressure sintering study initiated to determine the mechanism controlling the densification of spherical copper powder. Temperatures of 300°, 400°, 500°, and 600°C at 2500, 5000, 7500, and 10,000 psi pressures were investigated. The copper powder for this study was obtained from Metals Disintegrating Co., Elizabeth, N.J. The minimum purity was reported to be 99.9 pct. The powder size was -325 mesh and a lineal analysis showed the average size to be less than 1 p. An alumina die body and punches were employed. Graphite spacers were placed between the copper powder and the punches before pressing. It was expected that the graphite would form carbon monoxide with any entrapped air and form a reducing atmosphere within the die cavity. This condition was generally realized during pressing although a minor amount of oxidation occurred at 500" and 600°C. The oxide particles were located randomly in the pressed samples and were not considered to influence the densification process. Calibration tests showed the temperature of the copper pellet to be about 20°C less than the die wall temperature, probably resulting from heat losses through the punches. This difference was compensated during the experimental pressing operations. The hot pressing equipment employed for this work has been described by Spriggs et a1.9 and was calibrated with a Baldwin-Lima-Hamilton load cell. Travel of the punches was monitored with a dial indicator showing increments of travel to 0.0005 in. Relative density was calculated at selected time intervals for a known copper weight, punch travel, and die diameter. The accuracy of this calculation was *1 pct. Average grain size after hot pressing was obtained by lineal analysis. Typical densification curves are presented in Fig. 1 and Fig. 2 where relative density is plotted vs log time as a function of applied pressure for the 300" and 600°C pressure sintering temperatures, respectively. The curves for 400" and 500°C were similar to those at 300°C and consequently are not shown. The curves at 300°C as well as 400" and 500°C do not show an end-point or limiting density for a particular applied pressure. At 600°C, a density in excess of 95 pct of theoretical has been reached such that further densification is extremely slow for the 5000, 7500, and 10,000 psi pressures. It is suggested that further densification at these pressure levels may be retarded or prevented by discontinuous grain growth, by entrapped gases, and by other microstructural limitations as described by Cob1e.10 The curve for 2500 psi hot pressing pressure is a straight line over the time interval although a slight change in slope may be occurring at the last time intervals. The plastic flow models,1,2 derived for hot pressing, predict the existence of an end-point (or limiting) density dependent on the yield point of the material at the temperature of interest. The yield stress of copper at a 0.005 strain level has been

Jan 1, 1965

By Philip N. Adler, Harold Margolin

Univariant equilibria between the Ibcc, IIhcp, and UIfee phases of thallium have been determined and indicate a triple point at 110°C and 37 kbars. Calculated II —I transformation curves are compared with the experimental. Transformation volumes and/or enthalpies are calculated and enthalpies are compared with available data. ThE triple point of solid thallium has been reported by Bridgman as 153°C and 38 kbar.&apos; More recently, Jayaraman et at.2 have reported the triple point as near 115°C and 39 kbars. Meyerhoff and Smith,3 assuming the elevated-pressure modification of thallium is fcc, as Kaufman suggested,4 and that it is continuous with the room-temperature fcc phase in the Tl-In system, have used pressure data up to 5 kbars to predict the triple point at 162°C and 35.8 kbars. Piermarini and weir5 have indeed found the elevated-pressure phase to be fcc. The effect of variations of pressure and temperature in the IIhCp — Ibcc transformation in thallium have also been reported by werner6 and more recently by Ponyatovskii.7 The work reported here is part of an investigation of the effects of temperature and pressure on the Tl-In system. I) EXPERIMENTAL PROCEDURE The onset and progress of transformation were determined by changes in resistance. A commercially available tetrahedral-type pressure apparatus was used. Specimens of 0.060 in. diam, representing 0.63 pct of the cell volume, were encased in AgCl and mounted in pyrophyllite, as shown schematically in Fig. 1. It can be seen that the specimen is in series with the graphite furnace. Resistance changes of 0.1 pct could be detected with a Kelvin bridge. Resistance of the sample was of the order of 1 milliohm. Other pertinent details in regard to calibration and cell-assembly design are given elsewhere.&apos; Temperature calibration of the cell was carried out in two ways as a function of input wattage. Through a hole drilled in the graphite furnace, a chromel-alumel thermocouple was inserted into the AgCl to a point 0.010 in. from the sample, with the

Jan 1, 1964

By K. E. Nelson

The effect of thorium and zinc variations on the strength and 100-hr creep characteristics of Mg-Th-Zn-Zr alloys was investigated. Optimum resistance to creep at 650° and 700°F are attainable within a certain range of thorium and zinc contents. This range does not conform to that which develops maximum tensile properties. RAPID advances have been made during the last few years in the development of magnesium alloys for elevated temperature applications demanding high resistance to creep. The beneficial effect of rare-earth metals on the creep resistance of magnesium alloys has been emphasized by a number of publications1-13 and such alloys are now in commercial production. The use of thorium as an alloying ingredient in magnesium was mentioned by McDonald and also in two Alien Property Custodian patent applications.10,17 The initial observation of Sauerwald that thorium contributes still higher creep resistance to magnesium than is attainable with rare-earth metals has recently been substantiated.&apos;" " In fact, it has been demonstrated that the useful temperature range of magnesium alloys is appreciably extended by the use of thorium. In all cases, it was observed that zirconium must be included in the alloys in order to render them fine grained and more readily castable. Several recent publications:2-27 indicate that a still further improvement in creep resistance and a further extension of the useful temperature range can be realized by the addition of zinc to alloys. The primary purpose of this paper is to present the results of a comprehensive study of the effect of zinc on the strength and creep characteristics of Mg-Th-Zr alloys. Compositions covering the range of thorium content from M to 6 pct and zinc content from 0 to 5 pct have been investigated. The creep characteristics at 650" and 700°F reported in this paper are based on results of tests of 100-hr duration. It is appreciated that creep tests of 100-hr duration might not yield adequate data for design purposes for parts with much longer expected life. However, for the purposes of the present discussion, it is felt that the combination of stresses and temperatures used in the 100-hr creep tests have yielded a clear representation of the compositional variation of creep resistance at the temperatures investigated. Creep tests of 1000-hr duration are now in progress on a few of the most promising alloys. Preparation and Testing of Alloys The alloys studied in this intensive investigation were prepared in 25 lb capacity mild steel crucibles. The thorium, zinc, and zirconium were alloyed and poured as described in earlier publications The thorium was introduced into the melt in the form of a Mg-Th hardener," the zinc added in the metallic form, and the zirconium alloyed in the form of the commercial hardener containing magnesium and 30 to 50 pct Zr.10, 26 Fluxing practices for melting and refining were the same as for magnesium-rare-earth metal-zirconium alloys. The melts were sampled for analytical determinations and poured into separately cast 1/2 in. diameter standard tensile bars. The test bars were given a precipitation treatment of 16 hr at 600°F in laboratory furnaces. It has been shown by other tests that a high temperature solution treatment followed by an aging treatment is unnecessary for the development of optimum properties in Mg-Th-Zn-Zr alloys. The selection of 600°F as the aging temperature was based on an attempt to achieve metallurgical stability without coalescence of the undissolved phases and the attendant loss in strength. The thorium, zinc, and zirconium contents of each melt were determined chemically. The zirconium contents are reported in two parts, "soluble" and "insoluble," referring, respectively, to the portions present in the alloy which are soluble and insoluble in dilute HCl acid. Distinction is being made between these two components of the zirconium content in the alloys because it has been found that only that portion of the zirconium content which is soluble in dilute inorganic acids affects the structure and properties of the alloys. The usual impurities consisting of copper, iron, manganese, and nickel were determined spectroscopically. The analysis for each melt is listed in Table I. A description of the methods of tension and creep testing has been detailed in earlier papers. The tests were performed with the cast skin re-

Jan 1, 1954

By J. J. Casey, R. S. Busk, D. L. Leman

The mechanical properties and the phase-temperature relationships of magnesium-lithium-aluminum, magnesium-lithium-zinc, and magnesium-lithium-aluminum-zinc alloys are presented. It is shown that the strengths of some of these alloys can be tripled by heat treating and aging but that the high strength is accompanied by brittleness and lack of stability. MAGNESIUM-LITHIUM alloys are interesting because the hexagonal structure of magnesium can be changed into the more ductile and workable body-centered cubic structure by the addition of only 10 pct lithium.1 After Loonam2 determined that the ductility of the magnesium-lithium alloys is destroyed by even trace quantities of sodium, it was possible to prepare ductile alioys and study their mechanical properties. Jackson3 and coworkers have ably described their results with both the binary and polynary magnesium-lithium alloys. The present paper is a report of the mechanical properties of wrought magnesium-lithium-aluminurn, magnesium-lithium-zinc, and magnesium-lithium-aluminum-zinc alloys. Experimental Procedure The alloys were made by melting 7 lb of magnesium in a steel crucible under a molten flux of 75 pct LiCl + 25 pct LiF3 previously dried at 1600°F. The metal was heated to 1300°F for alloying aluminum and zinc and then cooled to 1250°F for adding lithium (sodium content: 0.01 to 0.04 pct). After the lithium had been alloyed, dry LiF powder was sprinkled on the melt to thicken the excess flux. After a 15 min holding period, the flux was skimmed off. Paraffin vapors were used to protect the metal as it was poured into 3 in. diam by 11 in. long magnesium molds. Billets were extruded into 3/8-in. diam rod and 1/16 x 3/4-in. strip; sheet was prepared from 1/8x 7/8-in. extruded strip, first by hot rolling and then by cold rolling 2 pct per pass. Hot rollability, or the temperature range through which the alloys could be rolled, was determined by rolling the strip two passes at 30 pct reduction per pass. Cold rollability, or the percentage of cold reduction to cause detrimental cracking, was determined on 0.064-in. hot-rolled sheet. Standard ASTM procedures were followed to obtain the mechanical properties, using 0.2 pct offset for the yield strengths. The level of properties produced with commercial magnesium alloys using the procedure described is given in a previous publication." Results Magnesium - Lithium - Aluminum System: The good workability of magnesium-lithium alloys is illustrated by the data in table I, which shows approximately how the extrusion pressure decreases as the amount of cubic phase increases. Binary magnesium-lithium alloys are soft and weak, but aluminum raises their strengths con-

Jan 1, 1951

By R. P. O’Shea

AS part of a program to investigate phase relationships in ternary systems containing semiconducting elements and compounds, an X-ray survey of alloys in the system Pb-Te-Bi was carried out which revealed that the pair PbTe-Bi probably formed a quasibinary system. This portion of the ternary was examined fully by thermal analysis, X-ray, and microscopic procedures. More than twenty-five alloys were made by melting together the proper combinations of the elements in quartz tubes at a temperature of 1000o C. The compositions which were high in PbTe content were prepared in evacuated, sealed, quartz tubes, allowed to freeze in vacuum, and then transferred to Crucibles for thermal analysis. The bismuth-rich alloys were made directly in the crucibles under a protective gas cover. The starting elements were obtained from the American Smelting and Refining Co. and had a purity greater than 99.99 pct. Cooling curves were made with about 50 g of alloy in graphite crucibles which were placed in a covered quartz container in which a protective atmosphere of helium or nitrogen was maintained. The thermocouple protection tube was made of mul-lite and served as a stirring rod when connected to a reciprocating stirring motor. Temperatures were

Jan 1, 1962

By C. G. Dunn, W. W. Martin

Various X ray diffraction techniques have been developed for determining orientations of crystals. Transmission and back-reflection Laue methods1,2,3 in particular have been found to be very useful. In many applications, however, a need still exists for a simplification of the details of analysis so that orientations can be determined more rapidly, and operators with little or no training in crystallography can be easily instructed in the necessary procedures. Filling this need in part is a system of analysis that has been in use for several years in the Metals Section of the Laboratory, General Electric Co., Pittsfield, Mass. This method is an extension of the one of Majima and Togino,4 which uses fifty-five indexed standard transmission Lauemams to cover the entire range of crystal orientations. Greater rapidity in determining the orientation of a crystal is achieved by using a slightly larger set of standard Lauegrams and by systematizing the steps followed. This system of analysis in its present form constitutes the subject matter of the present paper. Although the method uses transmission X ray Laue patterns, many of the operations are applicable to the analysis of back-reflection Lauegrams. Very briefly the method is based on the fact that Bragg and azimuthal angles for two or more identified Laue diffraction spots completely determine the orientation of a crystal, and standard Laue photographs of the Majima and Togino type provide a ready means for identifying Laue spots. After a preliminary analysis, which consists of identifying the pattern in terms of one of the standard patterns, a correctly marked Lauegram is selected to carry out the final analysis. From measured Bragg and azimuthal angles, the normals (poles) of the reflecting planes for the selected Laue spots are located on a stereographic projection and the orient.at.ion of the crystal determined (in terms of the positions of cube poles for example) by a few additional operations on a \I1ulff net. The actual procedure followed will be illustrated in detail after a brief treatment or X ray problems connected with orientation determinations and a discussion of useful principles and techniques. X Ray and Plotting Problems RELATIVE ADVANTAGES OF BACK-REFLECTION AND TRANSMISSION LAUE DIFFRACTION METHODS In the back-reflection Laue method it is easy to make a direct measurement of the tilt of a zone of Laue spots. Specimen thickness also is generally no problem. In the case of thin specimens, however, the time of exposure is less for the transmission Laue method. For example, with silicon ferrite this advantage holds up to a thickness of 0.6 mm if tungsten radiation at 40 kv peak voltage is used. Furthermore, the transmission photograph represents the entire thickness of the crystal whereas the back-reflection picture represents only a surface layer (a thickness of less than 0.05 mm for silicon ferrite), the deeper layers supplying a negligible amount to the diffraction pattern. If a crystal is distorted, the transmission Laue pattern gives more information particularly with regard to the orientation of the crystal and the overall amount of spread in orientation. Finally a 3 1/4 X 4 in. film suffices for transmission photograms; a 4 X 5 in. or 5 X 7 in. film with a central hole is needed for back-reflection patterns Not only is the manipulation of film easier because of the size of film, but

Jan 1, 1950

By G. R. Speich, R. A. Oriani

The rate of coarsening (Ostwald ripening) of copper precipitate Particles in an a-iron matrix has been studied in three Fe-Cu alloys containing 2.3, 4.0, and 5.4 wt pct Cu in the temperature range 730° to 830°C. The kinetics of Ostwald ripening ohevs the law as deduced by Lifshitz and Wagner so that matrix diffusion controls the process. The equilibrium shape of the particles is that of rods with hemispherical caps, and is due to the BOTH Lifshitz and slyozov and wagner3 have recently solved the complex problem of determining the rate of coarsening of an assembly of precipitate particles. Both workers, although using somewhat different mathematical approaches, arrive at the same equation for the relation between the mean radius 7 of a distribution of spherical particles and the bulk solubility co, diffusivity D in the matrix, large anisotropy in the copper/a iron interfacial free energy (y,/y, = 4). A Lifshitz -Wagner equation, modified for both the nonsphericity of the diffusion field and the reformulation of the Gibbs-Thomson equation for the case of nonsphencal particles, is derived. Using the modified equation, the calculated values of the mean radius are smaller than those observed by a factor of about one -half. absolute temperature T, gas constant R, surface free energy y of the precipitate/matrix interface, and the time t of aging: where V, is the molar volume of the precipitate and v is a stoichiometric factor. Wagner&apos;s treatment is restricted to the case of spherical particles growing in a liquid matrix where strain energy is not a factor. Lifshitz and slyozov have discussed the effect of anisotropic particle shapes and internal strains on the rate of coarsening. Oriani4

Jan 1, 1965