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By Leonid V. Azároff, H. N. Murty

In the course of investigating the soft X-ray absorption spectra of Ni-A1 alloys, Das and ~zfiroff' observed an anomalous brittleness in solid solutions containing 3 to 5 at. pct Al. The brittle alloys also had anomalously larger lattice constants' and the fine structure at the nickel K edge of an alloy containing 3.96 at. pct A1 unexpectedly resembled that of NiA1 more closely than that of a solid solution containing 5.51 at. pct Al. This led them to suggest the possibility that short-range ordering was produced during the homogenization anneal to which each specimen was subjected, similar to that observed in dilute solutions of vanadium and of chromium in nickel: as the principal cause of embrittlement. It was decided to repeat the above measurements on several alloys in this composition range in order to test the above hypothesis and to establish the effect of heat treatment on the alloys' properties. Three solid solutions in this range were prepared containing, respectively, 3.16, 4.00, and 4.54 at. pct Al, in addition to alloys containing 6 and 9 pct. All samples were annealed in vacuum at 1800°F for 100 hr and then furnace-cooled to room temperature, as in the original study. The lattice constants of these alloys are shown by open circles in Fig. 1 and can be seen to fit quite closely the curve established by the previously measured values, indicated by solid circles. (The experimental errors are of the order of the size of the circles, +o.oo~R.) The three alloys containing 3 to 5 at. pct A1 were found to be quite brittle, as previously reported. In order to test the possibility of ordering, these three solid solutions were again heated to 1800°F and, following a few hours at that temperature, were quenched directly to room temperature. All three alloys lost their brittleness. In order to verify the proposed ordering hypothesis, it was decided to subject each sample to several cycles of quenching and annealing (disordering and ordering). After some experimentation, it was established that reproducible results could be obtained by anneal-

Jan 1, 1968

By D. Mills, G. B. Craig

ZIRCONIUM has been float-zone-refined in an electron-beam furnace and the redistribution of oxygen, iron, and tungsten has been measured. The iodide zirconium used in the present experiments initially contained both oxygen and iron in the range 100 to 200 ppm by weight, and tungsten in amounts less than 10 ppm. Float-zone refining of zirconium, using induction heating, has previously been attempted by Kneip and Betterton1 and Langeron.2 Kneip and Betterton were primarily interested in the removal of iron and nickel. They achieved some purification, with respect to both these elements, in material given up to six passes at 150 mm per hr under an argon atmosphere. Langeron reported purification for a large number of elements after four completed passes in a static vacuum system operating at pressures less than 10-6 mm of Hg and with rates of zone travel between 30 and 40 mm per hr. He did not report oxygen concentrations, but stated that there was inverse segregation of this element. westlake5 has also used a minimum of four zone passes to remove iron from zirconium. No rates or conditions were given. Belk6 has shown that tungsten contamination can occur during electron-beam melting. He reported an increase from 0.01 to 0.05 pct by weight tungsten in molybdenum after four complete zone passes. The vacuum system for the electron-beam unit used in the present investigation consisted of a single-stage rotary pump backing a liquid-air-trapped oil-diffusion pump. A pressure of less than 10-8 mm of Hg was obtained with Dow-Corning 705 fluid in the diffusion pump. In order to avoid contamination of the zirconium by evaporation from the tungsten filament, a special focusing cage4 was employed. Three rates of zone travel, 114, 38, and 4 mm per hr, were investigated, with the liquid zone moving from the bottom to the top of the bar. The bars used were 3 mm diam and the total melted length was 130 mm. Oxygen was analyzed by neutron-activation analysis using a neutron flux of 10' neutrons per sq cm per sec to form the isotope N16 by the 016 (n,p) N16 reaction. The standard deviation is quoted for each oxygen determination. The iron and tungsten analyses were performed spectrographically, and the precision is estimated to be ±6 pct. Analyses for tungsten were all below the detectable limit of 10 ppm, confirming the protection given by the focusing cage. No significant redistribution of oxygen was found at the two higher rates of zone travel. The redistribution of oxygen and iron obtained after five passes at a rate of zone travel of 4 mm per hr (1.1 x 10-4 cm per sec) is recorded in Tables I and 11. Burris, Stockman, and Dillon3 have estimated the distribution of solutes during multipass zone refining. Using their curves for an effective distribution coefficient of Keff = 1.5 and 2 for oxygen and Keff = 0.3 for iron, the expected concentrations were estimated for the present material. These are shown in Tables I and 11, along with the experimentally measured values. The calculated concentrations are based on a molten-zone length of 10 pct of the total melted length, whereas in the present experiments the molten zone was approximately 5 pct of the total melted length. The effect of zone length on solute redistribution is most pronounced after a large number of zone passes. Comparison with Pfann's published data8 for solute redistribution for various Keff's and zone lengths indicates only small differences at low numbers of completed zone passes. It is evident that the expected distribution has not been realized in the case of iron. Scrapings of material deposited on the inner surfaces of the electron-focusing cage were found to be magnetic. It is, therefore, concluded that redistribution of iron is masked by evaporation. Deposition of iron inside the focusing cage was observed at all rates of zone travel. The results of the investigation may be summarized as follows: 1) Segregation of oxygen is typical for a solute with a distribution coefficient, Keff > 1. 2) No redistribution of oxygen takes place at high rates of zone travel. 3) The distribution coefficient for oxygen lies between 1.5 and 2.0. 4) Purification with respect to iron occurs mainly by evaporation. 5) The focusing cage is effective in preventing tungsten contamination.

Jan 1, 1967

By J. W. Sandelin, L. S. Birks

The group five metals, vanadium, niobium, and tantalum all show cold-rolled textures of {100)(011) and {112)(0ll) with a preference for the (100) in vanadium and niobium and for the {112) in tantalum. Recrys-tallization texture of (100) and (111) was reported1 for vanadium after annealing (no temperature stated). For tantal~m ,~ annealing below 1200°C strengthened the (112) texture but the (111) texture increased and was predominant as the annealing temperature was raised to 2500°C. Annealing niobium up to 1100°C strengthened the ( 100) texture and developed a weak (111) texture.' In the present work niobium has been annealed up to 2000°C. Nearly 70 pct of the grains showed {111) texture with only a weak (112) component and an insignificant (100) component remaining. In addition, double-crystal rocking curves were used to follow the increase in crystal perfection of the grains with increased annealing time. Material and Treatment. Samples i by 4 in. were sawed from & -in.-thick cold-rolled niobium sheet obtained from Wah Chang. Chemical analysis of the material showed it to be 99.4 pct pure with tantalum as the major impurity. Individual pieces were annealed for 2, 15, 30, and 60 rnin in a vacuum of 10'* Torr using an rf generator. They were etched until the grain boundaries were well-defined, Fig. 1. Average grain size varied from 1 mm for 2 min annealing to 2 mm for 60 rnin annealing. Experimental Results. Preferred Orientation. The orientations of 97 of the 117 grains in the specimens of Fig. 1 were determined by back-reflect ion Laue patterns. A General Electric 10-cm-diam cylindrical camera was used rather than the conventional flat camera in order to record more of the low-index spots. The method proposed by Forest, Barton, and schieltz4 was used to transform the cylindrical film patterns to standard stereographic projections. Table I gives the compilation of data. Even 2 rnin anneal at 2000°C is sufficient to yield a strong (111) preferred orientation, and the percentage of grains with this orientation increases steadily with annealing time at the expense of all other orientations. Crystal Perfection. X-ray topographs were taken of each specimen in both the (111) and (112) orientation using the method described by Birks and Grant.' As the anneal time increases, the grains become larger and the individual reflections become slightly more uniform in intensity, indicating an increase in crystal perfection. For a more quantitative measure of crystal perfection, it was necessary to measure the double-crystal rocking curve. The spectrometer employed was a Gaertner instrument of Compton design. LiF was chosen for the first crystal because the (400) reflection with d spacing of 1.05 AU is a good match to the (222) niobium reflection with d spacing of 0.95 AU. CuKp radiation was used in order to further reduce the d spacing mixmatch correction.' A variable aperture was placed between the LiF and the niobium specimen so that a single grain or even a small portion of a grain could be examined. Fig. 2 shows typical rocking curves for the four annealing times. For the 2-min specimen, a single peak is observed with 10' of arc full-breadth at half maximum. As the annealing time is increased to 15 or 30 min, several sharper peaks of 40 to 80" breadth can be seen within a single grain. After 60 rnin anneal only a single peak of about 30" is observed. Table I1 shows the corrected rocking curve breadths and the dislocation densities, ND. The latter were calculated from the equation7 where W is the corrected rocking curve breadth at half maximum in radians and b is the effective Burger's vector. For (110)(111) slip in niobium, b is 2.86 AU. Discussion. The annealing of cold-rolled niobium at 2000°C causes rapid preferential growth of grains with {ill} orientation. After only 2 rnin annealing time, 35 pct of the grains have the (111) orientation, and after

Jan 1, 1968

By Seymour G. Epstein

Using a modified capillary-reservoir technique, electromigration of dilute solutions of silver in liquid bismuth has been measured with relatively good precision. The effects of experiment duration, temperature. current density, and silver concentration have been investigated. Chemical diffusion coefficients of silver in liquid bismuth were also measured at temperatures ranging from 300°to 600°C. The effective valence of silver in bismuth was calculated as a function of temperature from a comparison of electric mobilities and diffusion coefficients. In an alloy containing 1 wt pct Ag, the electric mobility of silver increased from 1.65 x X sq cm per v-sec at 300°C to 2.38 x 10-4at 600°C; the diffusion coefficient in-creased from 3.0 to 17.5 x 10-5 sq cm per sec over the same temperature range. Calculated values of effective valence decreased from 0.28 at 300°C to 0.10 at 600°c. ALTHOUGH there has been considerable progress in the theoretical treatment of electron-transport properties in liquid metals, the mechanisms of mass-transport phenomena in liquid metals are still not clearly understood. This is especially true for the case of electro migration, alternatively termed elec-trotransport or electrodiffusion, which is mass transport induced by an applied electric field. A thorough knowledge of the mechanisms governing electromigration will considerably aid in understanding the electronic state of an impurity atom in a liquid metal. Electromigration has been long observed and studied in liquid metals and alloys. However, no satisfactory theory to explain the observed effects has been developed. The disparity in the present theoretical treatments, summarized by Verhoeven,1 has been attributed to the difficulty in describing the momentum transfer between electrons and ions and to the lack of precise measurements in liquid metals, especially at elevated temperatures. Prior to the last decade most of the electromigration studies were qualitative, in that only the direction of migration of alloy constituents was determined with any degree of reliability. These results have been summarized by schwarz,2 who was able to measure the electric mobility of several elements in mercury. More recently electromigration in liquid metals has been measured by a variety of techniques, with a varying degree of success. Mangelsdorf3,4 made precise measurements on dilute amalgams at room temperature, using a technique relating changes in composition to changes in alloy resistivity, but was unable to make measurements at elevated temperatures. Using a modification of this technique, Verhoeven and Hucke5 measured electromigration and resistivities in liquid Bi-Sn alloys. However, the resistivity technique is not sensitive to small composition changes and is impractical for very dilute alloys. Verhoeven and Hucke6 also measured electromigration of several solutes in liquid bismuth by adapting the capillary-reservoir method for determining diffusion coefficients in liquids. The reported precision of these measurements was poor, however, primarily due to the large uncertainties in the chemical analyses to determine concentration changes. Belashchenko7 has summarized the recent Russian investigations of electromigration, many of which, however, are of questionable accuracy. The present study was initiated to develop an experimental technique which would be both versatile and sufficiently sensitive to yield precise measurements of electromigration of solutes in dilute metal alloys at elevated temperatures. Using this technique, a comprehensive study of silver in liquid bismuth was made to determine the effects of experimental variables on the apparent electric mobilities. Chemical diffusion coefficients of silver in liquid bismuth were also measured over the same temperature range for comparison with the electromigration data. EXPERIMENTAL TECHNIQUE Electromigration Measurements. Electromigration of silver in liquid bismuth was measured by the capillary-reservoir technique, similar to that used by Verhoeven and Hucke,6 but modified to yield more

Jan 1, 1967

By K. M. Rolls, J. Wulff, D. A. Colling

Solid-solution alloys of the Ta- Ti system containing up to 70 at. pct Ti were melted and fabricated into wire. Steady magnetic-field measurements of cold-worked wires at 4.2°K indicate that the resistive critical field increases smoothly with increasing titanium, reaching a maximum of 93 kG (J = 1 amp per sq cm) for a composition of about 50 at. pct Ti. The composition dependence of the resistive critical field indicates it is largely paramagnetically limited for alloys containing more than 35 at. pct Ti. For cold-worked wires the value of the critical current density, Jc , was found to be insensitive to composition changes for compositions in excess of 20 at. pct Ti. It was observed that Jc could be increased by final-size heat treatment at 400°C without altering the value of- the resistive critical field. Experimental data are used to calculate the electronic specific-heat coefficient and the "upper critical field" for the alloys studied. THE effect and importance of metallurgical structure on superconducting properties have been summarized by Livingston and Schadler.1 Specific work on solid-solution alloys such as Nb-zr2-7 and Nb-Ti5-7 have shown that critical current density (Jc) and superconducting transition temperature are structure-sensitive. Berlincourt et al.6,8,9 in their studies of transition-metal alloys find that Ta-Ti alloys remain superconducting at fields comparable to Nb-Ti alloys. 6,8 Transverse critical current tests of a Ta-10 at. pct Ti alloy10 and longitudinal critical current tests of a Ta-25 at. pct Ti alloy1' indicate that such alloys are capable of carrying high currents at intermediate fields (-40 kG). The work reported in the present paper was undertaken to explore the superconducting properties of further compositions in the Ta-Ti alloy system as well as the structure sensitivity of the superconducting properties. I) EXPERIMENTAL PROCEDURE Alloys were prepared by melting tantalum (Fansteel Metallurgical Corp., 99.96 pct purity) and iodide, crystal bar titanium (Foote Mineral Co., 99.92 pct purity) mixtures in an arc furnace on a water-cooled copper hearth using a tungsten electrode. Melting was accomplished in a purified, gettered argon atmosphere at pressures from 250 to 400 Torr. Each alloy was turned and remelted fourteen times to assure complete mixing of tantalum and titanium. Ex- cessive coring in the cast samples was, nevertheless, observed. This was somewhat reduced by annealing in vacuo at 1000°C for 13 hr. The ingots were then machined to 0.250 in. diam, jacketed in Type 304 stainless-steel tubing, and swaged to 0.140 in. diam. They were then drawn to 0.010-in.-diam wire. Alloys containing up to 70 at. pct Ti were readily cold-drawn to wire in this manner. Later experiments showed that oxygen additions up to -3000 ppm could also be cold-worked into wire. Test samples for superconductivity measurements in the 2-in, air-core Bitter solenoids of the M.I.T. National Magnet Laboratory were prepared from 10-in. lengths of 0.010-in.-diam wire, cleaned of residual wire-drawing lubricant. Current contacts were made by clamping clean copper tubing over both No. 18 current lead wire and the ultrasonically indium-tinned ends of the specimen wire.12 Potential leads of No. 32 copper wire were connected to the specimen by ultrasonic soldering in an indium bath. The potential lead separation was about 4 cm, corresponding to the circumference of the Bakelite rod to which the wire was affixed with half-mil Teflon tape. The final specimen configuration is shown in Fig. 1. Specimens were lowered into a liquid helium bath (4.2oK) in a double solenoid dewar and positioned so that the maximum magnetic field was oriented transverse to the wire axis of the gage section. Measurements were made by passing a current through the specimen at a preset solenoidal magnetic field. The current was increased until a 1-pv potential difference appeared across the gage section (1/4 µv per cm). At low magnetic fields, however, it was not possible to use this criterion because power-dissipation difficul-

Jan 1, 1967

By E. A. von Hahn, T. R. lngraham

The rates of cementation of pd11 on electropolislzed copper cylinders were studied in aqueous perchloric acid solutions at pdII concentrations from 0.02 to 0.1 mM. The cylinders were rotated at various high speeds to reduce the diffusion layer to a minimum thickness. In 0.1 M HClO4 solutions, the cemented palladium is dense, adherent, and shiny. The rate data indicate that there are two stages in the deposition. In the initial stage, the rate of cementation is first order with respect to the pd11 concentration. The second stage is much slower. The first stage is consistent with rate control by the diffusion of pd11 ions to the copper surface, and/or chemical reaction at the surface, and the second stage is consistent with rate control by the diffusion of copper ions from the cop-per surface through the deposit of cemented palladium, out to the main body of solution. In 0.001 M HClO4 solution, only the first stage is evident and the rate is more rapid than at higher acidities. This rate enhancement is attributed to PdOH+ ions that predominate at low acidity and aye more reactive than unhydrolyzed pdII. The deposit is porous and loose. All of the cemented deposits are Pd-Cu alloys rather than pure palladium. Activation energies are 9.5 kcal per mole in 0.1 M HClO4 and 7.4 kcal per mole in 0.001 M HClO4. CEMENTATION or displacement reactions occur between aqueous solutions of metal salts and immersed metals according to the general equation: where M1 is electrochemically the more noble metal. Although these reactions find considerable application in metals recovery processes (e.g., the cementation of copper192 or gold3), in electrorefining processes (e.g., solution purification before electrolysis4,5), and in plating and metal finishing processes,6 few studies have been made of the kinetics of such reactions.7-9 The rates of cementation reactions will depend on one or more of the following factors: a) chemical reactions at the metal-solution interface; b) ionic transport to or away from the interface;'-' and c) the character and adherence of the cemented deposit on the substrate metal,6a,10 because the deposit will inhibit the rate of transport of dissolving ions (M2m+) into the solution. In this investigation, the kinetics of the early stages of cementation were studied to obtain an understanding of the reaction mechanism under conditions in which the deposit was sufficiently thin that its inhibiting effect could be disregarded. To achieve this, very dilute solutions in Mn+ were used. Cementation rates are often fast in the initial stages and transport-contr01led.9 To find conditions under which transport control would shift to chemical control, very high stirring velocities were used to minimize the diffusion-layer thickness. Of the many possible cementation reaction systems, the palladous perchlorate-copper system was chosen for this investigation because it was believed to involve simple ion-for-ion exchange. In addition, there are no interfering side reactions, such as the reduction of hydrogen ions, and anion effects are usually absent in perchlorate solutions. EXPERIMENTAL Materials. Reagent-grade chemicals and redistilled water were used in all experiments. Palladous perchlorate stock solutions (0.02 M) were prepared by dissolving 99.5 pct Pd sponge (Johnson Matthey and Mallory Ltd.) in concentrated nitric acid, adding concentrated perchloric acid, evaporating twice slowly (to prevent PdO precipitation) to a small volume, and diluting in a volumetric flask. To prevent gradual hydrolysis of PdII, the stock solutions were made 0.4 M in HC104. Tests for pdIV (Ref. 11) and C1- were negative. Copper strips were cut accurately from 0.025-in. electrolytic tough pitch sheet, Sample 1 (Mines Branch stock or Canadian Copper Refineries Ltd., Montreal), or from oxygen-free, high-conductivity sheet, Sample 2 (American Metal Climax O.F.H.C. brand, sold by Utility Brass and Copper Corp., Brooklyn). The strips were annealed for 1 hr at 335°C in a stream of purified nitrogen (Union Carbide, Linde Division) which had

Jan 1, 1967

By K. M. Rolls, F. W. Reuter, J. Wulff

The superconducting behavior and microstructural characteristics of a nominal Nb-40 wt pct Ti-0.239 wt pct O alloy were studied as a function of ther mo -mechanical processing treatment. Critical current density us applied transverse magnetic field was obtained for 0.010-in.-diam wires at 4.2°Kin steady fields 14 to 110 kG. Both optical metallogvaphy and transmission electron microscopy were used to delineate the micros tructures of the same wires. It wan found that a 1-hr 500°C precipitation heat treatment after cold drawing to final size led to the highest critical current density. Heat treatment at 600°C also led to a high critical current density, but the precipitate differs in kind and form from that at 500°C. The resistire critical field was also found to be sensitive to precipitation heat treatment since the effective composition of the superconducting phase changes. This is discussed in terms of the oxygen in interstitial solid solution. Two types of high-field superconducting wire are at present used in the construction of high-field superconducting solenoids. These types are solid-solution alloy wire such as Nb-Zr and Nb-Ti and composites of the brittle inter metallic compound Nb3Sn. The latter generally have a high super cur rent-carry ing capacity which is difficult to vary if properly made. The supercur rent- carry ing capacity of the former can be varied drastically and often predictably by suitable thermomechanical processing treatments. In general, the critical current density Jc of the solid-solution type of alloy is increased by cold work and by additions of interstitial elements along with aging heat treatments. The imperfections which result are be-iieved to be responsible for the observed increase in Jc. In 1962 Kneip and coworkers1 found that the critical faurrent density of Nb-Zr alloys could be increased by proper heat treatment preceded and followed by cold work. Betterton and coworkers2 using a Nb-25 at. pct Zr alloy found that small additions of oxygen or carbon enhanced the effect of this heat treatment. They suggested that the interstitials present aided precipitation in the alloy, leading to a filamentary structure with superior properties. If the precipitation heat treatment was omitted, interstitial additions had a negligible effect on Jc. wong3 showed that higher heat-treatment temperatures lowered Jc. Walker and co-workers,4 who studied microstructure (by transmission electron microscopy) as well as superconductivity, found that the Jc anisotropy introduced by cold rolling was itself affected by heat treatment. They were unable to clarify the relation between microstructure and critical current density, although evidence of precipitation was indicated. More recent investigation of Nb-Zr alloys,5,6 besides showing that structural defects and fiber ing due to cold work and precipitation serve to raise Jc, also elucidate important optically observable microstructural changes which occur upon precipitation. In these reports, coarsening of the microstructural features was found to decrease Jc. Vetrano and Boom,7 who studied Ti-20.7 at. pct Nb, found that Jc was increased to a maximum by a 415°C, 3-hr heat treatment following quenching from 800°C and cold working. Heat treatments can also affect the resistive critical field Hr. Final-size heat treatments of Nb-Zr wire can lower Hr drastically if gross phase decomposition occurs5'* or moderately if the effects of cold work are eliminated without changing significantly the composition of the phase of interest.3,5,6,8 The percentage of oxygen which can be added to Nb-Zr alloys to enhance Jc is limited by the difficulty of subsequent cold drawing. Since Nb-Ti and Ta-Ti alloys in contrast can tolerate appreciably higher percentages of oxygen, it was decided to investigate the superconducting behavior of various alloys in these systems. The present paper describes the results of adding oxygen to a nominal 40 wt pct Nb alloy as a function of thermomechanical treatment. I) EXPERIMENTAL PROCEDURE A small alloy ingot was prepared from high-purity niobium, iodide, crystal-bar titanium, and Nb2O5 powder by arc melting on a water-cooled copper hearth in a gettered argon atmosphere. The ingot was turned and remelted fourteen times to insure homogeneity. After final melting and rapid cooling, it was machined round to 0.415 in. diam, jacketed in stainless steel, and cold-swaged to 0.117 in. diam. The jacket was removed and swaging continued to 0.051 in. diam followed by wire drawing in carbide dies to 0.010 in. diam. Although it was intended that about 1500 ppm O (by weight) be added, inert gas fusion analysis indicated a 2390 ppm 0 content, apparently due to additional oxygen pickup in the arc furnace. Even so, the alloy was sufficiently ductile to be cold-worked to greater than 99.9 pct reduction

Jan 1, 1967

By M. D. Banus, M. B. Bever, A. K. Jena

The heals of formation at 78", 195, and 273°K of the ternary compound h3SbTe2 based on the elements and based on the binary compounds In Sb and [inTe have been measured. The heats of formation at these temperatlcres of the binary compounds In Sb and In Te based on the elements have also been determined. Heal contents and free energies of the three compounds have been calculated from 0° lo 80I)°K. The free energies of joyrrzalion, heats of formations, and entropies of formation at 298°K have also been calculated. The results shown that the tertnary compound is metastable with vespecl to InSh and ln Te below 696 °K. bul is slable above that temperature. The weaker bonding of results in a positice entropy of formation which with incrensirzg temperature makes increasing negative conlvihtclions to the free energy and above 696°K renders the compound slable. THE ternary compound In3SbTez occurring in the quasi-binary system In Sb- In Te' forms on cooling at 829°K by a peritectic reaction.' Observations at 673" and 573 K have shown that this ternary compound decomposes slowly into the binary compounds InSb and1n~e.l'' It has not been possible to analyze the metastable behavior of the ternary compound because up to the present time data on its thermodynamic properties have been lacking. Some information on the binary compounds, however, is available. The heat of formation of InTe at 273°K and its free energy at 673°K are kn~wn.~'~ The heats of formation of InSb at 78", 273', 298", and 723°K have been measured5-' and its heat capacity between approximately 12" and 300"Kg9l0 is also known. In the investigation reported here the heats of formation at 78% 195% and 273°K of the ternary compound In3SbTez have been measured. The heats of formation of the binary compounds InSb and InTe at these temperatures have been obtained by combining new calorimetric results with previously published data. The heat contents and free energies of the three compounds at various temperatures from 0" to 800°K have been calculated. Against the background of this information, the metastability of the ternary compound will be discussed. 1) EXPERIMENTAL 1.l) Preparation of Specimens. The materials used consisted of the elements indium, antimony, and tellurium, the binary compounds InSb and InTe, and the ternary compound In3SbTez. The elements, obtained from American Smelting and Refining Co., had a nominal purity of 99.999+ pct. The compound InSb was Cominco semiconductor grade; the compound InTe was prepared from the elements by melting under a vacuum of 10-h m Hg followed by slow cooling. Three batches of specimens of the compound In3SbTez were used. One batch was prepared by melting appropriate amounts of the elements in an evacuated and sealed Vycor tube. The melt was held at approximately 100°K above the liquidus for about 8 hr, shaken repeatedly, and quenched into a mixture of ice and water. The specimen was annealed in vacuum at 760°K for 4 weeks. In preparing a second batch, a mixture of the component elements was melted and quenched in water. The resulting ingot was powdered. The powder was pressed into pellets, which were annealed in vacuum at 748" to 773°K for 4 weeks. A third batch was prepared in the same manner as the second, except that the starting materials were InSb and InTe rather than the elements. Metallographic examination of samples of the three batches and X-ray diffraction examination of a sample of the second batch did not reveal evidence of microsegregation or a second phase. The results obtained with the three batches showed no systematic differences. 1.2) Calorimetry. The calorimetric method has been described in detail." Samples of the compound In3SbTez, mechanical mixtures of InSb and InTe in the proportion of 1:2, and mechanical mixtures of indium, antimony, and tellurium in the proportion of 3:1:2 were added to a bismuth-rich bath at 623°K in a metal-solution calorimeter. These additions were made from 78°K (liquid nitrogen), 195°K (dry ice and acetone), and 273°K (ice and water). The heat effects of the additions were measured. The difference in the heat effects of the additions of a compound and the additions of the mixtures of its constituents, adjusted for differences in the concentration of the bath, is the heat of formation of the compound. In the concentration range not exceeding 1.5 at. pct solute, the heat effect of the additions was a linear function of concentration. The heat of formation refers to the temperature from which the additions were made, namely, 78", 195", or 273°K. Several calibrating additions were made in each calorimetric run. The calculated heat capacity of the calorimeter and hence the calculated heat effects of the additions of samples depend upon the values adopted for the heat contents of the calibrating substance. In this investigation bismuth at 273°K was used and a value of 4.96 kcal per g-atom was taken for (HGZ3"k . 2) RESULTS AND DISCUSSION 2.l) Heats of Formation. The heats of formation of the ternary compound In~SbTez from the component elements indium, antimony, and tellurium and from

Jan 1, 1968

By G. R. St. Pierre, M. J. Pool, R. Speiser

The activities of chromium in solid Cr-Ti solulions contaitning from 10 to 90 at, pct Cr were measured over the temperatutre range 1250" to 1380°C. The Knzudsen effusion technique with direct weighing of the Knudsen cell was used. This was possible because the vapor pressure 0f chromium is about 600 times greater than the capor pressure of titanium so that the amount 0f titanium effusing could be neg-Lected. the activities of titanium were determined by g-raphical integration of the Gibbs-Duhertz equatiotz. The solutio~ls sllcdied exhibited positive deviations from Raoult's law, but they approached ideality as the temperature was raised from 125 to 138t C. The integral free energy of mixing was calclated from the activity values. AF has a maximum value of - 1.66 kcal a1 1390°C and a composition oj 55 at. pcl Cv. The free energy of formation of 'TiCr, is —0.76' lzcal a1 1250°C. The heats and entropies of mixing usere determined using the second law and thus tile mlues could be in error by a considerable amount. The Cv-Ti system is not a regular solution. ThE increasing emphasis on application of various models to alloy formation and solution behavior has led to an intensified interest in the thermochemistry of alloys. In a majority of the cases where different models have been applied the alloy system used was a simple one. However, the amount of data available on more complicated systems is small and in order to fully understand the behavior of metallic solutions more experimental work has to be carried out on these systems. The Cr-Ti system is reasonably complex and even though this system was being investigated as early as 1937 by Kroll,' no previous work has been done on the thermodynamic properties of either the liquid or solid state. This research was undertaken to determine the thermodynamic properties of solid Cr-Ti alloys. EXPERIMENTAL PROCEDURE The experiments on the Cr-Ti system were carried out in a water-cooled metal vacuum system. This system is composed of a brass vaporization chamber wound with copper tubing for cooling purposes, a brass liquid-nitrogen trap, an oil diffusion pump, and a Duo-Seal forepump. The pressure in the system is measured with an ionization gage. The interior of the vaporization chamber is shown schematically in Fig. 1. The Knudsen cell is supported by the molybdenum tripod, and a magnetically operated shutter is provided to protect the Pyrex sighting glass against deposition of metallic vapors in the intervals between temperature measurements. Heating is accomplished by means of a 20-kw 450-kc Westinghouse oscillator. The susceptor in this case is 10-mil tantalum seamless tubing which also forms the outside of the sample container. The remainder of the sample container consists Of an im~ervims re-crystallized alumina crucible and a 10-mil-thick tantalum lid. Two holes are drilled in the lid. The diameters of these holes were measured with the calibrated stage of a "Tukon" microhardness tester. The thicknesses were measured with a micrometer. A Leeds and Northrup disappear ing-filament optical pyrometer was used to measure the temperature. This is sighted through the smaller of the two holes in the lid and corrections for the emissivity of the sample were not necessary since the ratio bf cell length to hole diameter was large enough to insure a hohlraum. The pyrometer was calibrated at the ~~t~~~l Bureau of Standards, and the correction for the Pyrex sighting window was determined in the laboratory. The titanium used in this investigation was obtained from Mallory-Sharon and the chromium from The Union Carbide Metals Co. The reported purities were

Jan 1, 1968

By R. H. Maes, R. E. de Strycker, J. J. Jacobs

A simple method, based upow density measurements, has been perfected in order to determine the critical temperature of liquid miscibility gaps. Applied to the Pb-Cu system, it yielded a value of 980"C for the critical tempevature of the miscibility gap in this system. In the Pb-Cu-As system, the immiscibility between lead and coppev evolves towards a miscibility gap be-tween lead and copper arsenides; the liquation temperatures were determined joy different avsenic contents and values up to 1065°C were found. Combining these results with those given by the clzemical analysis of liquid samples taken at dqj.event temperatures, it has been possible to make a cowzplete determination of the liquid miscibility gap in this ternary system. In the Pb-Cu-Fe-As system, a ternary miscibility gap has been discovered; for some compositions, the molten and slowly cooled alloys separated into three layers: an Fe-As alloy, a Cu-As alloy, and a lead-rich layer. ThE Pb-Cu-Fe-As system is of interest for the study of the behavior of speiss or arsenical alloys in the lead blast furnace. The speisses produced in lead metallurgy generally have complex compositions, but they may be considered, in a simplified way, as a mixture of copper, iron, nickel, and cobalt arsenides; there is often an excess of metallic constituants, and lead is also present. Speisses liquate from lead at various temperatures according to their composition; according to industrial experience, iron speisses are already distinctly separated from lead bullion at temperatures of the order of 1100°C, whereas copper speisses only begin to liquate from lead at much lower temperatures, near their solidification points, which were never accurately determined until now. Besides these considerations of practical interest, the present study is of importance for the knowledge of liquid miscibility gaps. Particularly the data of the literature about critical temperatures of liquation are incomplete and present sometimes considerable disparities; a method has been perfected in the present work to determine these temperatures with accuracy. MISCIBILITY GAP OF THE LEAD-COPPER SYSTEM Review of Literature. The generally admitted outline of the Pb-Cu equilibrium diagram is shown by Fig. 1. This figure, with its numerical indications, is taken from the review of Hansen and ~nderko;' compositions are given in weight percent, as will be the case in the remainder of this paper. Between 36 and 87 pct Pb, Pb-Cu alloys separate into two layers, in the liquid state. The critical temperature of liquation has been much disputed in earlier studies, where values up to 1500°C were claimed; but a reliable study of ~ohnen,' based upon differential thermal analysis, gave a value of 990°C; this rather low value was verified by a recent study of Schiirmann and Kaune? based upon thermodynamical properties of the Pb-Cu alloys, which predicted a value of 1002°C. Liquid miscibility gaps were also studied by means of chemical analyses. Satisfactory results were generally obtained with the method consisting in taking samples of the two liquid layers in equilibrium and submitting them to chemical analysis; however, this method becomes delicate near the critical temperature. With this technique, Friedrich and waehlerts ob-

Jan 1, 1968

By A. P. Woods, J. A. Elias, A. J. Heckler

A method is described by whichX-ray pole figures and axis density figures may be plotted automatically. The X-ray data are obtained using a Siemens texture gonionleter and the Schulz reflection method. Composite sampling methods are used to obtain a full quadrant of the pole figure. The contour lines of the pole figure and axis density figures are drawn using a computer-controlled x-y plotter. The computer program used to draw the contours was obtained from ZBM. POLE figures used to describe crystallographic orientations of poly crystalline sheet materials are usually plotted manually from X-ray diffraction output in the form of a strip chart recording of the intensity distribution of a given crystallographic pole in physical space. Today's computers and associated plotting equipment now make it possible to plot pole figures and axis density figures automatically. Inverse pole figures have been determined by mathematically "inverting" X-ray pole figure data.'-' In general, these methods require complicated, time-consuming numerical calculations. A more economical but less accurate method is to determine the axis densities of low-index crystallographic poles using an X-ray diffractometer. When these axis densities are plotted in a unit stereographic triangle contour lines can be drawn to represent the axis density distribution. This can now be done automatically. Computers were used to plot pole figures by a method developed by Holland, Engler, and powers.' In their method intensity ratios based on a random standard sample were obtained from a strip chart recording, and the stereographic (x, y) coordinates of these ratios were determined by a computer. Symbols characteristic of different intensity ratio levels were then plotted using an automatic x, y plotter and intensity contours were drawn manually. The present paper describes a method by which pole figures and axis density figures have been automatically plotted from digital data obtained from a Siemens X-ray unit. Plotting was accomplished using an IBM 1620 computer and a Calcomp 563 digital plotter. COMPUTER PROGRAM The computer program was obtained from IBM Systems Reference Library, No. 1620-CS-05X, and was designed to plot contour maps from randomly distributed data. It consists of a set of twelve separate subprograms, four of which were used in this work. The program takes a given distribution of data points and by a numerical method determines the values at mesh points of a square grid system superimposed over the input data. The output data from this program is in a form suitable for input to the contour plotting program. A smoothing program is also available which calculates new mesh point values from a continuous function fitted to neighboring mesh points. The degree of smoothing depends on the number of points around a given grid point which are used to adjust or smooth the grid value. The contour plotting program is designed to allow the user to specify the interval, scale, and annotation of the contours to be plotted. EXPERIMENTAL PROCEDURE Pole Figure Data Points. Diffracted X-ray intensities were determined using a Siemens texture goniometer and the schulzs reflection method. Due to geometric

Jan 1, 1968

By R. D. Carnahan, G. A. Stone, R. J. Arsenault

An experimental technique has been developed for carrying out a dynamic tensile stress-strain test in which plastic strain is measured continuously throughout the microstrain region extending through the macroflow region to total deformations of 5 pet. The tests were carried out on niobium, samples having interstitial impurity levels of 160 and -800 ppm at temperatures ranging from room temperature down to 100°K. A band spectrum of activation energies was obtained from calculations based on the measured activation volumes and temperature dependence of flow stress. The inability of a single rate-controlling process to predict such a phenomenon has led to the proposal that several sequential rate-controlling dislocation mechanisms are operative in the preyield micro-strain region. These are thought to be: the motion of geometrical kinks, the formation of double kinks in edges, and finally in the macrostrain region the formation of double kinks in screws. In less pure niobium the effect of interstitial impurities is shown to be dominant in the microstrain region, suggesting that a fourth probable mechanism is overcoming interstitial barriers instead of geometrical kink motion and the formation of double kinks in edge dislocations. THE principal fundamental aim of most mechanical properties studies of crystalline solids is to determine the dislocation mechanisms responsible for the plastic behavior. The extension of such studies to the region of preyield microstrain behavior has received renewed attention in the past few years as the details of macroscopic flow behavior have become better understood. For the most part, microstrain studies have been carried out using one of two variations of a single experimental technique.1 The first method involves loading the sample at a constant extension rate to a predetermined stress level (less than the macroscopic yield stress), instantaneously unloading, and then determining the extent of plastic deformation to a sensitivity of 10"6. The microscopic stress-strain curve is then constructed from the results of a repeated series of such individual measurements, each of which involves loading to an incrementally higher stress level, by relating the accumulated plastic microstrain for each step to the maximum applied stress in that step. The second variation, also carried out by step-wise increases in stress, measures the total strain to a sensitivity of ~10"6 and relates the maximum stress to the area of closed hysteresis loops that are generated under certain conditions (namely, prestrained samples). The foregoing methods have yielded interesting information about the nature of preyield deformation and plastic response to stresses below the yield stress, but they offer certain disadvantages for studying the dynamic behavior of dislocations. In the first method it is physically impossible to specify a plastic strain rate inasmuch as the test sample is reexposed, as it is loaded to its maximum stress in a particular step, to lower stresses at which flow will take place. The strain corresponding to a given stress level on the stress-microstrain curve thus generated will then always be greater than that at the corresponding stress level of a truly dynamic test carried out at the same extension rate. The second method has been used to obtain correlations with dynamic macroscopic flow parameters by applying dislocation damping theory to the hysteresis loop areas.2 The approach has had some success but suffers in that only samples that have been prestrained a few percent can be studied. Although the stresses

Jan 1, 1968

By Arnulf Muan, Klaus Schwerdtfeger

Equilibrium ratios C02/C0 of a gas phase coexisting with selected phase assemblages of the system Fe-Mn-0 have been determined in the temperature range 1000" to 1300°C. The oxygen pressure for the "hfnO" +hfn30, equilibrium and for the "(Fe,hTn)O" + (Fe,Mnh 0* equilibrium at high manganese contents has been determined by electromotive force measurements using stabilized zirconia as a solid electrolyte. The notstoichometry 01' "hTnO" and of "(Fe, iM1z)O" solid solutions has been determined by ther-mog-/avi?netry and by wet-chemical analysis. The data obtained are used to derive activity-composition relations in "(Fe,hfn)O" and (Fe,Mn),O4 solid solutions. WUSTITE "FeO" and manganosite "MnO" form a continuous series of solid solution at high temperatures,' and so do magnetite Fe304 and the high-temperature, cubic modification of Mn304 (Ref. 2) (high hausmannite, -1170). The oxides "FeO" and "MnO" are cation-deficient phases.495 The nonstoi-chiometry of "(Fe,Mn)O" solid solutions has been studied by Engell and ~ohl' at two selected C02/C0 ratios at 1250°C. The two oxide end members of the spinel solid solution, FesO4 and Mn,04, however, are known to be close to stoichiometric under the experimental conditions used in the present investigation.''' The oxygen pressures of "(Fe,Mn)07' solid solutions in equilibrium with iron have been determined by Schenck and coworkers,8 by Foster and welch," and by ~n~e1l.l' The two former groups equilibrated the condensed phases in C02-CO atmospheres of lmown compositions, whereas Engell" used a galvanic cell with stabilized zirconia as a solid electrolyte. The results of these investigators are not in good agreement. Activities of FeO in manganowiistite as calculated from the results of Foster and Welch show ideal behavior, those of Engell yield a pronounced positive deviation, and those of Schenck et 01. show a moderate positive deviation from ideality. In the present work oxygen pressures for the iron + manganowiistite and manganowustite + spinel equilibria and the nonstoichiometry of manganowiistites have been measured. The data were used to calculate activities in the manganowiistite and spinel solid solutions. EXPERIMENTAL METHODS The COz/CO ratios at which manganowustite and iron are in equilibrium were determined by thermo-gravimetric and quenching methods. Experimental details are described in a previous publication.'2 In the thermogravimetric technique, incipient reduction of manganowiistite pellets to metallic iron was observed as a break in the weight vs log COZ/CO curve. In the quenching technique, manganowiistite samples were partially reduced to metallic iron, or the metallic iron of manganowustite + metallic iron mixtures was partially oxidized to manganowustite, in atmospheres of constant C02/CO ratios. After quenching the composition of the oxide phase was determined by X-ray lattice parameter measurements and comparison with a standard curve obtained from oxide solid solutions of known compositions. The nonstoichiometry of "MnO" and "(Fe,Mn)07' solid solutions was determined by chemical analysis of samples equilibrated in C02-CO atmospheres and quenched to room temperature, as well as thermo-gravimetrically by reducing (Fe,Mn),04 or Mn304 to manganowiistite or manganosite. The equilibrium between manganowiistite and (Fe,Mn),04 was measured thermogravimetrically by reducing (Fe,Mn),04 solid solutions having composition in the range of %„ l(NFe +NM) from 0 to 0.63. No experiments could be performed with this technique at higher manganese contents, because the equilibrium C02/C0 ratios are too large for accurate control. An additional difficulty arises at the higher manganese contents due to the strong increase in oxygen content of the manganowustite phase with increasing log Py near the manganowiistite-spinel boundary. Consequently a sharp break in the weight loss vs log C02/CO curve cannot be observed at the phase boundary. At high manganese contents of the manganowiistite, e.g., (NMn/(NF~ + NMn) > 0.9, electromotive force measurements with stabilized zirconia as a solid electrolyte were made to determine the equilibrium oxygen partial pressure. Experimental details are described in a previous paper.* Mixtures of "(Fe,Mn)O" and (Fe,Mn),04 were pressed to pellets, and the oxygen pressure of the equilibrated samples was compared to that of Ni + NiO mixtures in the cell The composition of the manganowiistite in the equilibrated two-phase mixture was determined by lattice parameter measurements and comparison with known standards. The oxygen pressure for the Ni + NiO equilibrium was taken from available data.l3~l4 No reliable results were obtained with the electromotive force technique on iron-rich oxides. The electromotive force drifted strongly with time in this composition range. An additional difficulty arises from the partial de-

Jan 1, 1968

By S. Jana, C. M. Wayman

The reverse transformation of bcc martensite to the fcc phase was studied in an Fe-33.95 wl pct Ni alloy by nzeans oj dilatometry, melallography, and electron microscopy. Upon "slozc" heating (-1°C per min) length cJmnge us temperature plots showed u gradual contracLion over the temperature range 200" to 280"C ,followed by a more abrupt contraction beginning a1 -280°C. Howet,ev, zchen the heating rate was increased -4°C per tnin, no gradual contraction was observed and only the abrupt contraction starting at -2BO"C was found. Thus on slower heating- the AS "temperature" for the subject alloy, unlike the MS temperature, is better defined as a range of temperatures. Both optical and transmissiorl electron microscope observations showed that some of the martensite plates exizibited a partial loss of transformation twins during reversal. The midvib region of the martensite plates disappeaved relatively early duirng the reversal. Metallographic observations slowed that the earliest detectable stage of the rezlerse tvansforrvration begins (axd Moues inulardly) at The Martensens i te - parent interface. At higher temperatirres, the. formation of martensitically reversed jcc plates within the bcc martensite plales was observed. It is concluded that the reverse transformation consists of a diffusion less process (martensitic); but this is ps-obably aided by a prior or simultaneous dijjusiorz-comltvolled process, at leasl in the case of slower heat-ing' experiments. ALTHOUGH numerous investigations have dealt with the parent-to-martensite ("forward") transformation (fcc — bcc) in Fe-Ni alloys, comparatively little is reported on the ("reverse7') martensite-to-parent transformation.'-4 Even though such reverse transformations have been studied in detail in some nonferrous systems, one of the difficulties of studying the reverse transformation in most ferrous mar-tensites is that the martensite decomposes by tempering during heating. However, carbonless Fe-Ni alloys do not exhibit this difficulty since the transformation in these alloys is completely reversible. The present investigation represents an attempt to shed more light on the nature and mechanism of the martensite-to-parent transformation. 1) EXPERIMENTAL PROCEDURE 1.1) Alloy Prepatation. Fe-Ni alloys of compositions near 34 wt pct Ni were prepared from zone-refined iron (99.994 wt pct Fe) and high-purity nickel (99.999 wt pct Ni) by induction melting in recrystallized alumina crucibles in an argon atmosphere, with prior vacuum evacuation to 10"3 mm Hg. The alloys were homogenized by induction stirring in the molten state for 5 min. After solidification, the alloys were further homogenized in evacuated quartz capsules for 96 hr at 1230°C. 1.2) Dilatometry. Slices of the ingot were hot-forged (750°C in air) into approximate rod form and these specimens were then hot-swaged (750°C in air) into long cylindrical rods 0.55 mm diam. From the rods, specimens about 1 in. long were cut. These were then vacuum-annealed for 24 hr at 1200°C, cooled to room temperature, and subsequently transformed to martensite in liquid nitrogen (whereby about 40 pct transformation was obtained). Dilatation measurements were made by observing length changes in a vacuum dilatometer with an externally mounted LVDT sensing element. 1. 3) Preparation of Electron Microscope Specimens. Slices of the ingots were cold-rolled (with intermediate vacuum anneals) to -0.020 in. Out of these rolled sheets, specimens (about 1 by 1 in.) were cut. These were then vacuum-annealed, transformed to martensite by cooling in liquid nitrogen, and subsequently heated from room temperature to various temperatures to effect either partial or complete reverse transformation. These specimens were then chemically polished to 0.002 in. in l:l HsOz (30 pct) and &PO4 (85 pct) solution, and thinned to electron transparency in an electrolyte consisting of 150 g CraOs, 750 ml glacial acetic acid, and 30 ml ~~0.~ Observations were made with a 100-kv Hitachi HU-11 electron microscope equipped with an HK-2A tilting device. 1.4) Optical Microscopy. Metallographic observations were made with a Leitz MM5 metallograph on the same 0.020-in. sheet specimens as were used for electron microscopy and on bulk specimens which were 0.2 in. or more on a side. The chemical thinning solution when cooled below 20°C also served as an etchant for this alloy. Observations of surface relief were made with a Zeiss interference microscope employing a Thallium light source of wavelength 0.54 p. Specimens for interference studies were prepared by two-stage polishing on Buehler vibromet polishers using 0.3 and 0.05 p alumina abrasives. 2) EXPERIMENTAL RESULTS 2.1) Comparison of the MS,AS, and Af Tempera-tures wTth Previous Re sults. The AS aLd Af tempera -tures of several Fe-Ni alloys were determined dila-tometrically. The MS temperatures of the same alloys were determined by continuously lowering the temperature using a mixture of isopentane and liquid nitrogen and observing the highest temperature at which a prepolished specimen showed surface upheavals. For the present the As temperature is defined as the temperature at which an abrupt decrease in length occurs in the dilatation plot. The Ms,As7 and A determinations in the present investigation and those of Kaufman

Jan 1, 1968

By B. F. Oliver, C. W. Haworth

Pure tin and Sn-0.5pct Pb ingots have been frozen unidirectionally from the base. For quiescent melts that were initially undercooled, a transition from lower eqlciaxed structure to an upper columnar structure is found in the alloy ingots. Columnar to equi-axed back to columnar transitions are observed in superheated alloy ingots, but no such equiaxed band is observed impure tin. The reproducible equiaxed band is associated with a thermal undercooling followed by a recalescence. This undercooling is <5"C, whereas the critical (maximum obtainable) under-cooling for both the pure tin and the alloys used is -20°C. A similar undercooling is observed at the same position in the pure tin ingots, although in this case no clear transition in structure can be seen. The structure of the pure tin ingots is either entirely columnar or mixed columnar-equiaxed. A consideration of the detailed thermal history of the ingots indicates that the ingot macrostructures are determined by the occurrence of a local therlnal undercooling in conjunction with nuclei multiplication and transport mechanisrris. GENERALLY it is found that a pure metal ingot solidifies so as to produce an entirely columnar structure. Frequently an alloy ingot is found to have a columnar outer zone and an equiaxed central portion. Early systematic work to examine the factors controlling the formation of the equiaxed structure was reported by Northcott&apos; who showed that, for copper alloys frozen unidirectionally with a given ingot practice, the alloying element influenced the length of columnar crystals and the extent of the equiaxed structure. Northcott showed that alloys with a wider freezing range more readily produced the equiaxed structure. The nucleation process can be important in producing equiaxed structures; frequently an alloy which readily solidifies with an entirely columnar structure will produce an entirely equiaxed structure when a nucleating agent is added to the melt.&apos; The formation of the equiaxed structure was attributed by Winegard and chalmers3 to the presence of constitutional supercooling; that is, a region of liquid in front of the growing solid could have a temperature below its equilibrium liquidus temperature. Thus, with a small enough temperature gradient in the liquid, it was suggested that the presence of constitutional supercooling may be sufficient to bring about the nuclea-tion necessary for the formation of an equiaxed structure. Although this explanation is plausible, and may be relevant in many ingots, Walker has described an experiment&apos; for which constitutional supercooling seems to be an unlikely cause of nucleation. A Ni-20 pct Cu alloy, repeatedly undercooled more than 50"C, was crystallized and found to show the typical colum-nar-equiaxed structure. The separation between the liquidus and the solidus for the alloy is 40°C. Thus, in this experiment the nucleation required for the formation of the equiaxed structure must have come about in some other way than by the nucleation catalysis constitutional supercooling hypothesis. Chalmers has suggested more recently5 that nuclei (in a typical ingot) are present immediately after pouring and are prevented from redissolving by the constitutional supercooling effect. More recently Uhlman, Seward, Jackson, and ~unt&apos; have shown direct evidence using ice and organic materials that freeze dendritically that the "remelt mechanism" may be an extremely effective crystal multiplication process during the freezing of ingots under conditions involving dendritic growth. JSlia" experimentally demonstrated the detachment of dendrite arms. chernov14 has analyzed the dendrite arm detachment process as a coarsening phenomena driven by the minimization of interphase area. Katta-mis and ~lemings" working with undercooled steel melts give evidence supporting this mechanism. Mechanisms of dendrite arm detachment such as those assisted by convection are believed to be the origin of the macrostructures obtained in this study. This study makes no attempt to distinguish the relative contributions of these mechanisms. The object of the present work was to obtain accurate temperature measurements during the solidification of an ingot and to correlate these measurements with the formation of equiaxed grains in the resulting ingot structures. Similar previous work is very limited. The measurements carried out by Northcott are neither sufficiently extensive nor sufficiently accurate for any interpretation. Plaskett and winegard7 carried out experiments on A1-Mg alloys in which they observed values of the temperature gradient, G, in the liquid and rate of freezing, R (for a given alloy solute content Co), at the transition from a columnar to an equiaxed structure. They reported that equiaxed crystals were produced at values of G/G approximately proportional to the solidus composition. Similar experiments using Pb-Sn alloys carried out by £111011" showed a linear relation between G/R and the solidus composition. However, the thermocouples were in the mold wall rather than in the melt and, in one case, ingot surfaces were examined. There is ambiguity in the meaning of the values of G and R measured in all these experiments. APPARATUS AND EXPERIMENTAL PROCEDURE Alloys were prepared by induction melting 99.999 pct Sn and 99.999 pct Pb to form a Sn-0.5 wt pct Pb alloy in air in a graphite crucible and casting into a cylindrical graphite mold 6 in. long, 1 in. in diarn , and with a & in. wall thickness. This mold was mounted on a copper base through which cooling water could be

Jan 1, 1968

By B. H. Kear, B. J. Piearcey

The orientation and temperature dependence of the tensile and creep propevties oj Mav-M200 crystals halle been determined. Crystals oriented for single slip exhibit ,maximum ductility, minimum work hardening, and least creep resistance, whereas crystals in multiple slip orientations show nzinimu~n ductility, strongest work hardening, and greatest creep resistance. At 1400" and 1600°F there is a marked improzlenzent in creep strength for orientations approaching [001] and [111], indicating that interactions between dislocations gliding in intersecting octahedral slip systems play an important role in creep resistance. At 1800°F the creep strength is much less dependent on orientation, which is rationalized in terms of slip in cube planes as well as octahedral A frequent mode of failure in high-strength cast nickel-base superalloys is by intergranular fracture, particularly along those grain boundaries oriented transverse to the major stress axis. VerSnyder and ~uard&apos; demonstrated that this problem could be overcome by eliminating transverse grain boundaries through unidirectional solidification. Piearcey and versnyder2 made use of this principle in the development of unidirectionally solidified gas turbine blades and vanes, where the growth direction of the columnar grain structure coincides with the axis of principal stress under operating conditions. The present investigation was undertaken to determine if further improvement in properties may be obtained by eliminating grain boundaries altogether, so as to take advantage of the well-known dependence of mechanical properties of single crystals upon the orientation. 1) EXPERIMENTAL PROCEDURE Single crystals of Mar-M200* were grown from the melt under vacuum by a modified Bridgman method. The melt was poured into a preheated alumina mold, and crystal growth was promoted from one end by appropriate gradient cooling. Tensile and creep specimens (2 in. diam by $ in. gage length) were prepared by a series of operations involving electrical discharge machining, precision grinding, and electropolishing. The orientations of the specimens were determined by the Laue X-ray back-reflection method. Tensile tests were carried out in aWiedemann machine with furnace attachment, using a strain rate planes. In all orientations imzpvouement in the strength charactevistics of the material can be induced by heat treatment. Creep ad stress rupture data for (001) oriented crystals are compared with similar data obtailted previously joy random polycvystalline material, and also columnar grained material having a prejevved (001) orientation. The single-crystal )material exhibits both longer rupture life and lower minimum creep rate at all temperatuves, and the rupture elongation is comparable with that in the columnar grained material. From these results it is concluded that single crystals should and useful application in gas-turbine blades and tanes. The optimum orientation for a blade is considered to be with its axis of principal stvess parallel to (001) ou (111). of 0.0001 sec-&apos;. Load and extension were recorded directly on an X-Y recorder. The strain measuring device consisted of extension arms attached to the grips at one end and leading out of the furnace to an LVDT (linear variable differential transformer) at the other. Creep tests were performed using a standard constant load creep frame. 2) DISCUSSION OF RESULTS 2.1) Structure of Alloy. The structure and segregation in as-cast and heat-treated Mar-M200 has been described in detail else where.~ The main features are as follows: the as-cast material is heavily cored, due to pronounced dendritic segregation during solidification; the dendrites are rich in tungsten and cobalt whereas the interdendritic regions are rich in chromium, titanium, nickel, and carbon. The structure consists of -60 vol pct coherent precipitate of y&apos;, basically Ni3(A1,Ti), in a matrix of y (nickel-base solid solution), interspersed in the interdendritic regions with minor MC carbides and -y&apos; eutectic. The y&apos; particles, and y-y&apos; eutectic, contain more titanium, aluminum, and nickel, and less tungsten, cobalt, and chromium than the y matrix. Heat treatment partially removes segregation, eliminates the eutectic, refines the y&apos; dispersion in y, and gives additional partially coherent M23Cs carbide. An as-cast crystal, therefore, is composed primarily of two oriented phases (y + y&apos;), whereas the normally heat-treated crystal consists of three oriented phases (y + y&apos; + MZ3Cs). Typical electron transmission micrographs of the y + y&apos; structure are shown in Fig. 1. Crystals grown in the (100) orientation develop a simple unidirectional dendritic structure, Fig. 2, since (100) happens to be the preferred direction of growth for dendrites in this material. Crystals grown in the {110) and (111) orientations, however, tend to promote equal growth, generally in separate colonies, in the two and three geometrically favored (100) growth directions, respectively. In other orientations of growth, a single (100) dendrite direction generally pre-

Jan 1, 1968

By B. E. Terkelsen, B. J. Piearcey

A study has been made of the effect of unidirectional solidification on the creep behavior, stress-rupture properties, and thermal shock resistance of four nickel-base superalloys. The alloys Mar-MZOO, B-1900, IN 100, and TRW 1900 show improved rupture ductility and thermal shock resistance when tested with the columnar grains parallel to the major stress axis during test. The relative improvement in creep-rupture properties depends on the intrinsic strength of the alloy, a property which depends on composition, heat-treatnient, and crystallographic orientation. The data clarifies some of the factors affecting the properlies of the cast nickel-base superalloy. COMPONENTS designed for high-temperature use have, in recent years, been fabricated by precision casting techniques using nickel-base alloys developed specifically for use in the conventionally cast condition. This development was a result of both the increasing complexity of the component and the recognition that high-temperature strength was incompatible with workability. The use of castings can be economically favorable, but, more important, the recent complex designs of air-cooled gas turbine blades and vanes in alloys which possess the necessary high-temperature strength cannot be forged. Common modes of failure of high-temperature components are excessive creep, creep rupture, and thermal fatigue. If rupture occurs then the mechanism is usually by inter crystalline cracking along those grain boundaries oriented transverse to the major stress axis. In the stronger alloys, rapid propagation of intercrystalline cracks result in apparent premature failure demonstrated by the absence of third stage creep in a creep-rupture test and low rupture ductility. Tensile ductility shows a similar trend, that is, decreasing ductility with increase in strength. It is not surprising, therefore, that increases in creep strength have only been obtained with a loss in resistance to thermal shock, a property which shows a dependence on tensile ductility. VerSnyder and ~uard&apos; showed that the application of unidirectional solidification to a brittle Ni-Cr-A1 alloy both improved the rupture life of the alloy and increased rupture ductility. Since this casting method results in columnar grains, it appeared that the solution to the lack of ductility in nickel-base alloy components was the elimination of the source of failure, namely the transverse grain boundaries. This concept was recently developed2 to produce cast-to-size gas-turbine blades and vanes consisting entirely of columnar grains oriented parallel to the major stress axis of the component. Not only did the process impart increased rupture life and ductility to the components, but it also increased their resistance to thermal shock. During this development, the effect of unidirectional solidification on the properties of several alloys was investigated. The resulting data allows certain conclusions to be drawn regarding the factors affecting the creep and stress-rupture properties of the cast nickel-base superalloy and also its resistance to thermal shock. Detailed information was obtained on the alloys Mar-MZOO, B-1900, In 100, and TRW 1900. The composition of the alloys studied are shown in Table I. MATERIAL PREPARATION AND TESTING The conventionally cast alloys are cast as +-in.-diam bars in innoculated shell molds under conditions designed to maintain control of grain size. Unidirec-tionally solidified casting were produced in the form of 3-in.-diam ingots. Specimens from the latter were machined such that the axes of the columnar grains were parallel with the axis of the test bar, except in certain cases where the transverse properties of the material were being evaluated. The specimens were tested in an axial loading creep machine.&apos; Temperatures were measured with thermocouples, attached just outside each end of the gage length, to assure uniform temperature along the entire specimen. Chromel/alumel thermocouples were used up to 1800"F and Pt/Pt-10 pct Rh thermocouples above 1800"F. The specimen temperature was recorded and maintained within i2"F throughout the test. Creep extensions were measured automatically with extenso-meters attached to ridges on the specimen and recorded continuously using linear variable differential transformers and multipoint recording equipment. MICROSTRUCTURE The microstructures of transverse sections of the four alloys studied are shown in Figs. 1 and 2. Fig. 1 shows optical micrographs of the alloys, both in the conventionally cast and unidirectionally solidified conditions. Each alloy displays a cored dendritic structure, a distribution of MC carbide, and the presence of the y — y&apos; eutectic constituent.~ In the case of IN 100 and B-1900, Fig. 1 indicates that in these specimens the eutectic is degenerate. That is, it consists only of y&apos; pools indicating that the y constituent has been solutioned during the cooling period after solidification. The major effect of unidirectional solidification on the microstructures of the alloys is a tendency to promote a more dendritic form of MC carbide and an increase in dendrite arm spacing, both effects being a result of the slower solidification rate in the unidirectional solidification technique. V, lec tron micrographs of transverse sections of the four alloys in the two cast conditions are shown in

Jan 1, 1968

By M. J. Blackburn

The phasal equilibria in TI&apos;:Al alloys has been studied Ry transmission electron microscopy, electron diffraction, and X-ray diffraction. It is shown that three-phase fields exist below the transformation, viz.: a disordered hexagonal solid solution a; an ordered hexagonal phase based on Ti3Al, aa; and a two-phase region. (a + (YZ) , occurring between the single-plzuse fields. The size. disiribution, and morphology of- the a, phase in the two-phase region is shown to be dependent upon the alloy composition and heat treatement. The ordering process in the a, phase leads to the formation of anliphase domain boundary networks; the contrast and stability of these networks is shown to be dependent on alloy composition. A new and relatively simple phase diagralil for the systemm is constructed. The phasal equilibrium at the titanium-rich end of the Ti:Al system has long been the subject of considerable controversy. crossleyl has recently reviewed the phase diagrams that have been advanced for the system and in his turn proposed yet another one. There would seem to be general agreement that a phase exists at a composition at or near Ti& with an ordered hexagonal DO,,-type structure. However, several other phases have been proposed; for instance, Ence and ~ar~olin&apos;s&apos; diagram includes two phases, Ti2A1 and Ti3A1, in the region up to 25 at. pct Al. The position of the proposed phase boundaries varies over a considerable compositional and temperature range, a point graphically illustrated in Fig. 1 of Crossley&apos;s paper.&apos; It is perhaps germane to briefly examine the techniques used by other workers in the establishment of the phase structure. These fall into three categories: 1) X-ray diffraction, the only limitation of which is the rather low intensity of superlattice reflections in this system.3 Thus, superlattice reflections become increasingly difficult to detect in the more dilute aluminum alloys. 2) A considerable amount of optical microscopy has been performed"2&apos;4 and it will be shown in a later section that many of the obscure mi-crostructures that have been observed can be attributed to hydrogen contamination. It is concluded, therefore, that such observations should be interpreted with considerable care. 3) The variation of some physical property with aluminum content has been measured.5"7 These methods are a useful indication of the existence of a transformation, but are usually insufficient in themselves to identify the transformation. The results using these techniques will be considered in more detail in the discussion. In this paper, the phase structure of alloys containing up to 25 at. pct A1 has been studied by transmission electron microscopy and X-ray techniques, and the results are used to generate a new and relatively simple diagram for the titanium-rich end of the Ti:Al system. EXPERIMENTAL A series of alloys containing 5.05, 10.16, 13.39, 13.99, 14.40, 20.83, and 25.05 at. pct A1 were obtained in the form of hot-rolled sheet 0.02 in. thick. These alloys contained 0.10 to 0.15 pct 0. Thirty-gram arc-melted buttons were prepared from iodide titanium and 99.99 A1 which contained 16.1, 16.9, 18.2, 19.0, 21.6, 22.2, 23.0, and 23.9 at. pct A1 and had an average oxygen content of 0.03 pct 0. The differences in oxygen content did not appear to influence the nature of the phase transformations which were found. Oxygen content may affect the temperature of these transformations, but insufficient samples of the high-purity alloys were examined to answer this question. The heat treatment of alloys was performed in dynamic atmosphere and vacuum furnaces or in helium -filled silica capsules. A special vacuum furnace was constructed to provide a high quenching rate in which rates of 4000°C per sec were obtained using water as the quenching medium. As pointed out by Crossley,&apos; it is important to avoid heat treatment in the a - [} phase field as the resultant segregation of the alloy into the two phases containing different amounts of aluminum can persist through subsequent treatments leading to ambiguous results. It will be demonstrated later that step cooling of specimens from the a phase region can also lead to anomalous effects. Thus, all solution treatments were carried out in either the a or 3 phase fields followed by water quenching and subsequent isothermal annealing at various temperatures. X-ray analyses were performed on a diffractometer on fine-grained bulk specimens. The value of long-range order parameter, S, wa_s calculated from the integrated intensity of the (1211) reflection. In general, the intensity of the (1011) superlattice reflection was compared with that of the (2021) fundamental reflection to correct for experimental differences between measurements. Thin foils for electron microscopical examination were produced by electropolishing using the technique described in an earlier paper.8 Dark-field techniques were employed to a large degree in this investigation using the gun-tilt method. Finally it should be noted that the lattice parameters of the D01,-type superlattice are approximately related to the disordered hexagonal phase parameters by Crf - c, and 2ad - a, where the subscripts s and d refer to the ordered and disordered unit cells, respectively.

Jan 1, 1968

By E. Miller, R. Geffken, K. L. Komarek

The thermodynamzc properties oJ liquid Cd-Sb alloys were investigated using the cell arrangement measurments were obtained every 2°C at a heating and cooling rate of 12°C per hr and at equilibrium every 2O0C frorn 500°C down through the stable liquidus. The S-shaped asCd US composition curve was used in the cotnposition regzon near Cd,Sb, to calculate a tempeerture-dependent inleraction coefficient from quasichemical theory. Rapid changes in a scd were observed at a transition temperature varying from 400" to 465°C depending on con/kosition. It could not be determined if the changes in aScd were discontinuous, but tlze composition dependeke of the magnitude of the change is indicative of a second-order phase transformation in the liquid. The values of the experimental changes in ASCd are in agreement with calculations from the slope of the transition temperature, using the concept that a second-ovdev phase transition occurs in liquid Cd-Sb alloys. II is suggested that the transformalion is associated with the formation of Cd4Sb3 molecules in the liquid. ThE structure of liquid alloys is the subject of many investigations. X-ray, resistivity, and thermody-namic data have been interpreted as indicating varying degrees of short-range order in the liquid in alloy systems forming inter metallic compounds. In general, the melting process is not a transition from an ordered to a completely disordered state, but some degree of order is retained in the liquid. Maximum ordering in the liquid state occurs close to the melting temperature of the compound and the arrangement of atoms becomes more random at higher temperatures. Of special interest in this respect is the Cd-Sb system. It is one of the few metallic systems which form both stable and metastable compounds when liquid alloys are cooled at normal rates. The stable system exhibits an intermetallic compound, CdSb, melting at 459"c.l A second compound, CdrSbs, has also been reported,&apos; melting close to this temperature. The metastable system has one compound, CdsSbz, melting at 420"c.I Resistivity measurements on liquid Cd-Sb alloys close to the liquidus temperatures have been interpreted in terms of a complex ordering behavior which changes rapidly with increasing temperature.3 The resistivity-composition curve is characterized by two maxima corresponding in composition to CdSb and CdsSbz. The resistivity-temperature plots show sharp breaks for alloys in the composition range of 45 to 70 at. pct Cd on cooling through a transition temperature close to the stable liquidus. Fisher and phillips4 investigated the influence of temperature and composition on the viscosity of liquid CdSb alloys. The viscosity of some alloys increases sharply on supercooling below the stable liquidus. A maximum in the viscosity-composition curve occurs at the composition CdSb. The thermodynamic properties of liquid Cd-Sb alloys have been investigated by Seltz and ~e~itt&apos; and Elliott and chipmane by the electromotive-force method and their results are in good agreement. However, these investigations were carried out at temperatures well above the liquidus temperatures of the alloys, and the temperature coefficients of the electromotive force, dE/dT, were obtained from experimental points for each alloy at a few temperatures considerably above the liquidus temperature. Scheil and ~aach&apos; investigated the thermodynamic properties of this system by the dew point method in the temperature range from about 100°C above the stable liquidus down into the supercooled liquid region. They reported several anomalies, i.e., the activity of a melt on heating differed from that on cooling, and the activity increased sharply in the limited temperature interval immediately above the liquidus temperature of the stable alloy, followed by a sudden decrease below the liquidus. Values obtained on heating and cooling were not in agreement. A reinvesti-gation of a few alloys by Scheil and Kalkuhl&apos; by the electromotive-force method failed to confirm these observations. The authors concluded that the anomalies were due to inhomogeneities in the starting alloys and they discarded their previous results. The present investigation was undertaken in order to obtain thermodynamic data close to the liquidus temperature and in the supercooled region where the anomalies were originally reported, employing the electro motive-force method. This method is quite precise and will most easily permit observations of small changes in activity and partial molar entropy with temperature. Measurements were taken every few degrees so that the dE/dT values could be calculated over the entire temperature range and small changes in the thermodynamic properties close to the liquidus temperature could be observed. I) EXPERIMENTAL PROCEDURE Specimens were prepared from 99.999+ pct Cd and Sb (Cominco). Surface oxide was removed by scraping and then melting the metals under vacuum and filtering through Pyrex wool. Appropriate amounts of the metals were weighed on an analytical balance to k0.1 mg, sealed in double Pyrex capsules under vacuum,

Jan 1, 1968

By H. A. Wriedt

Zinc-vapor pressures in equilibrium with bcc solutions of zinc and iron at 703.5", 757", 793", and 900°C were measured by an isopiestic method. Thermody-na?nic properties of zinc in these solutions were derived directly and those of iron with a Gibbs-Duhem integration. Zirzc and iron exhibit large, positive, and moderately temperature-sensitive deviations from ideality. Excess enthalpies and Gibbs energies are positive in the ranges studied. It has been establzshed in relating published phase equilibria to the Gibbs energies of forming solid solutions that positive departures from ideality also exist in the fcc solutions. Tlze stability, as defined by Darken, of the bcc solutions is small but positive at 703.5"C. Its decrease with decreasing temperature indicates that a miscibility gap may exist in the stable or metastable bcc solutions at lower temperatures. The changes of thertnodynanlic properties relatable to changes in magnetic state are minor compared to other changes due to alloying. The accuracy with which the relation between lattice parameter and composition of the bcc solutions is known has been improved. Zinc increases the lattice parameter of a iron by O.ZOZA per unit atom action zinc to at least NZn = 0.25. TO our knowledge, only Ciganl has previously investigated thermodynamic properties of the bcc solutions of the Fe-Zn system. He measured zinc-vapor pressures in the temperature range from about 320" to 360°C by an effusion technique. We have measured the vapor pressures of zinc in equilibrium with bcc solutions of zinc and iron at 703.5", 757", 793", and 900°C by an isopiestic technique. This investigation and the pursuant thermodynamic computations are part of a program for studying several properties of these solid solutions. In the appendix we report a remeasurement of the relation at room temperature between the lattice parameter and zinc content. Related work not reported herein embraces measuring the diffusivity2 of zinc in ferrite at 700" to 900°C and the Curie point temperature3 as functions of composition. EXPERIMENTAL METHOD AND APPARATUS Purified iron at a series of constant temperatures was saturated with zinc at a series of constant zinc-vapor pressures, the range of conditions being so limited that the equilibrated solid solutions were ferri-tic. The equilibration was conducted in long capsules, Fig. 1, that had been evacuated, gettered with titanium, and sealed. The purified iron, part of a vacuum-melted ingot of electrolytic-iron stock (Glidden Co. 1044-grade), had the impurity contents listed in Table I. It was used in the form of millings about 0.001 in. thick. A zinc of 99.999 pct purity (United Mineral and Chemical Corp.) and later one of 99.9999 pct purity (Cominco Products, Inc.) were used, with no observed difference in results. Sufficient zinc was placed in the capsule (separate from the iron and kept at a different temperature) to supply the vapor phase and the iron alloy, while leaving a pool of pure liquid zinc throughout the run. The imposed temperature of the residual zinc pool fixed the vapor pressure of zinc throughout the capsule; condensation was avoided at other locations by making this the lowest temperature in the capsule. Two common complications of nonisothermal systems were, we believe, effectively absent from ours: first, molecular species other than monatomic have no appreciable concentration in zinc vapor at 700" to 900~~: second, the development of zinc-pressure differentials in the capsules was suppressed by use of relatively large-diameter tubing and the absence of diluents in the zinc vapor other than desorbed impurities. The specimens and zinc pool were located at two isothermal zones (each ranging within 1°C over several inches) of a wire-wound, 37-in.-long vertical-tube furnace. The required temperature distribution was achieved by external adjustment of the currents in three electrically isolated windings. The temperatures of the isothermal zones were controlled independently by ther-

Jan 1, 1968