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By R. E. Miner, J. K. Jackson, T. L. Wilson

A study has been made of the transformation kintetics and morphology in gas-quenched Fe-5 pct Ni alloys. Massive ferrite or nlurtensite was found to form at temperatures in agreement with earlier results. Surface Preparation appeared to have the greatest effect in determining the type of transformation Product formed. Grain size change, subgrain morphology, and dislocation densities of massive ferrite were essen-tinlly invariant functions of cooling rate. The mar-tens~tic structures were typical of transformation in austenite grains of the order of the sample thickness. BY introducing the idea of a massive transformation in dilute iron-base alloys, Gilbert and Owen in 1962 served to stimulate a closer inspection of phase transformations in ferrous systems.' A wealth of literature has followed concerned with the morphology,2-5 transformation mode,2'3'6-8 and thermodynamics of transformations.2,5,9 interesting controversies have resulted from the experimental observations; closer attention has been paid to the definition of massive transformations and still to be resolved is the Ms temperature of pure iron. The purpose of this note is to report some observations made in a detailed kinetic and metallographic examination of the transformations that occur in gas quenched Fe-5 pct Ni. The results are pertinent to the Ms extrapolations that have been made for pure iron and the observed substructure of massively transformed alloys. PROCEDURE The material used in this investigation was cast as a 5 lb vacuum induction melted, carbon deoxidized heat. The ingot was hot worked to a strip 0.10 in. thick, decarburized, and cold rolled to final thickness. Specimens were punched to sample shape and decarburized in wet hydrogen at 600°C. The chemistry of test specimens is indicated in Table I. Samples 1 by $ in., by thickness were resistance heated and rapidly cooled in an apparatus as illustrated in Fig. 1. Similar helium quenching apparatuses have been described previously;2,6'10 the distinguishing feature of this apparatus is the nozzle geometry. Twin nozzles with three convergent-divergent, sonic ports were mounted perpendicular to the specimen for quenching purposes. The specimen temperature was monitored with an intrinsic 3 mil chromel-alumel thermocouple spot welded to the specimen and con- nected to an oscilloscope. A single switch stopped the current flow through the specimen, opened a gas flow solenoid valve, and initiated an oscilloscope sweep. Cooling rates were read directly from a photograph of the oscilloscope trace and the quenching rates quoted later were made at a point immediately prior to transformation on the trace. All quenching studies were conducted either under a vacuum of 10"-5 torr or a helium atmosphere. The sample geometry was chosen large enough to permit optical and electron microscopy to be performed on samples. The samples were electropolished for both types of examination in a perchloric-acetic solution. Electron microscopy was conducted on a HU 11A microscope operated at 100 kv. Dislocation densities were determined at the same general operating reflection (110) on foils thick enough to produce Kikuchi patterns. RESULTS AND DISCUSSION Transformation Temperature Observations and OD-tical Microscopy. The quenching results of this study are illustrated in Fig. -2(a). Samples which were treated in the as-decarburized condition transformed in a manner similar to that reported previously—a decrease in transformation temperature occurred as the quenching rate increased until a plateau region was reached. Within the scatter of the results obtained here, there was no break in the transformation temperature-quenching rate plot. This suggests no change in transformation mode. Further metallographic examination confirmed this observation. A martensitic transformation was obtained for similarly treated samples except that the surface was first electropolished. The Ms temperature was approximately 85°C below the plateau determined for the unpolished samples and cooling rates as slow as 2000°C per sec were adequate for producing marten-site. The transformation temperatures are in gen- / STRIP SAMPLE / COPPER I SAX/-/}' / CONNECTOR-----A f, SON|C poRTS AC LEADS 3j ]j"I rofe[ \ Mf/yinK/iKj''^ THERMOCOUPLE X-(^—iie"s*c? ' LEADS 1 INCH Fig. 1-—Schematic scction of gas quenching apparatus

Jan 1, 1970

By E. W. Dewing

The solubility of A14C3 in cryolite melts has been determined by holding the melt with aluminum in a graphite crucible for 4 hr at temperature, quenching the crucible in water, and analyzing the frozen salt by CH4 evolution. The solubility is given by log(wt % Al4C3) = 1.298 -1.l74 (wt ratio NaF/AlF3) independent of temperature in the range 985" to 1080°C. Additions ofA1203, .NaCl, CaF2, or LiF up to 10 wt pct do not affect the solubility, provided that LiF and CaF2 are replaced by equivalent NaFin calculating the ratio. The variation with ratio can be accounted for quantitatively on the hypothesis that the Mechanism of solution is Al4C3 + 2AIF3 = 3Al2CF2 since the solubility varies as (aAIF3) 2/3; Mechanisms involving C", C2- , AlC- , A12C2+, and Al2CF+are re-jected. At ratios above 1.5 truces of c;- are also present. It was noticed some years ago in this laboratory that carbon specimens made cathodic in a cryolite bath decrease in diameter over a period of time.' It was also found that cryolite which had been melted with aluminum in a graphite crucible was sometimes yellow, and that A14C3 could be seen intimately embedded in the cryolite crystals. CH, could be recovered by acid attack. There was thus reason to believe that A14C3 had a genuine solubility in molten cryolite, and the object of this work was to measure it. EXPERIMENTAL AND RESULTS A graphite crucible, 2 in. ID and 2.5 in. deep with a close-fitting lid, contained about 150 g of cryolite (plus additives) and about 20 g of 99.99 pct pure Al. It was placed in a vertical Inconel furnace tube, covered with refractory brick to provide thermal insulation, and the top of the tube was roughly sealed with disks of asbestos paper. An alumina tube led argon down to the level of the top of the crucible to expel air, and an alumina sheath with a Pt-10 pct Rh ther-mocouple inside was positioned between the crucible and the wall of the Inconel tube. The tube was heated in an electrical resistance furnace for the desired length of time; the refractory bricks were then removed, and the crucible was lifted out with tongs and quenched in water. After such a crucible was broken open it could be seen that its inner surface was yellow owing to the presence of A1,C3, and sometimes needle-like crystals could be observed (at X75 magnification) growing out into the nielt. It was concluded that Al,C, had been present as a separate phase when the crucible was at tenlperature and that the melt had been saturated: in

Jan 1, 1970

By A. S. Venkatadri, H. B. Bell

The equilibrium sulfur distribution between molten iron and Ca0-Mg0-Al203 slags containing iron oxide was investigated at 1550°C. The results were used to derive the iron oxide activities at low iron oxide concentrations in the slag by combining the sulfide capacity data obtained from gas-slag work with the free energies of both the sulfur solution in iron and the iron oxide formation in slag. The derived ferrous oxide activities were compared with values based on Tem-kin's kin's and Flood's ionic models. One difficulty in using these models is that the nature of the aluminate ion in slag is uncertain. Nevertheless, such indirect methods, in particular, those described in the present paper, are of value because of the difficulty of measuring small amounts of oxygen in liquid iron in equilibrium with slag. It is shown that these methods confirm the consistency of thermodynamics data on liquid iron and slags. It is well established that decreasing the iron oxide activity in the slag increases the desulfurization of molten iron at constant slag basicity. This effect is most pronounced at the very low iron oxide activities, characteristic of blast furnace slags. Yet a precise quantitative determination of the significance of low iron oxide contents in slag in blast furnace desulfuri-zation is not possible for the following reasons: a) difficulty of separation of iron "shots" from the slag, and b) errors in chemical analysis of small amounts of iron oxide in slags. In view of these obstacles, one must resort to indirect methods of calculating iron oxide activities. EXPERIMENTAL TECHNIQUE The apparatus for providing the sulfur equilibrium data has been described previously1 and was similar to that used by ell' in connection with the study of slag-metal manganese equilibrium. The procedure consisted of: a) melting about 50 g of Armco iron in a magnesia crucible in a platinum furnace, b) adding a mixture of about 15 g of lime-alumina slag and varying amounts of Fe2O3 and CaS, and c) maintaining the temperature at 1550°C for more than an hour in an atmosphere of argon to enable the sulfur equilibrium to be attained. Several melts were made using lime-alumina slags with basic composition 55, 50, and 45 pct lime. During the experiment the temperature was controlled manually using a Pt/10 pet Rh-Pt thermocouple. After the experiment, the Power was shut off and the flow rate of argon was increased to freeze the melt as quickly as possible. The analysis of sulfur in the metal was carried out by the oxygen combustion method3 using uniform drillings from the top and bottom of the metal button. After crushing and grinding and removal of any iron particles with the aid of a hand magnet, the slag was analyzed for sulfur by the CO2 combustion method.4 The E.D.T.A. method was employed for the analysis of lime5,6 and magnesia,= the ceric sulfate method7 for the analysis of slag iron oxide, and the perchloric acid dehydration method5 for the analysis of silica. The remaining amount was taken to be Al2O3 precipitation with ammonium hydroxide in several preliminary melts had confirmed the propriety of using this simple procedure. RESULTS The activity of iron oxide in binary, ternary, and more complex slags has been the object of numerous investigations, and the two experimental methods for its determination are: 1) Equilibrating the metal with the slag in question and measuring the oxygen content of the metal. The ferrous oxide activity is then given by aFeO L%OJSat where [%0]sat is the oxygen content of the metal in equilibrium with pure iron oxide slag. This method was used by Chipman et al.8,9 2) Equilibrating the slag in iron crucibles with known partial pressures of H2/H2O or CO/CO2 mix-tures.10-12 This method is limited to temperatures between 1265" and 1500°C. The very low oxygen content of the melts in this investigation made it impossible to derive the ferrous oxide activity by the first of these methods. Therefore, the iron oxide activities were computed by means of: Sulfide capacity data from the gas-slag work" Temkin's concept14 Flood's approach15 a FeO from Sulfide Capacity. The method of calculating the aFeO involves the sulfide capacity of the slag (c,), the sulfur distribution coefficient (Ls), the free energy of dissolution of sulfur in iron, and the free energy of formation of iron oxide in the slag. Bell and Kalyanram13 have investigated the sulfur absorption characteristics of lime-alumina slags containing magnesia by the Carter-Macfarlane method16 (based on comparing the sulfide capacity of the slag in question with that of a standard slag of unit lime activity) and have derived lime activity values. The relation between sulfide capacity and their lime activity a'CaO is given by: Cs= 3—: Xa'CaO at 1500°C

Jan 1, 1970

By William D. Klopp, Joseph R. Stephens

ThE phenomenon of high ductility associated with a phase transformation, termed super plasticity, has been reported for a number of alloy systems. Abnormally high ductility has been observed during tensile testing of cast eutectic alloys of Pb-Sn and Mg-cu.1 Brasses of Cu-37 at. pct Zn and Cu-40 at. pct Zn elongate up to 180 pct when tested near their transformation temperature.' Enhanced ductility was also observed in Cu-47 at. pct Zn tested near the order-disorder transformation temperature. The most remarkable super-plastic effect reported to date is for the eutectoid alloy A1-60 at. pct Zn, where elongations up to 1000 pct have been reported.3 underwood4 has recently summarized the conditions that normally accompany the superplastic effect which are: 1) a phase transformation or allotropic change or order-disorder transformation; 2) an externally applied stress; 3) presence of an interface between two phases; and 4) a metastable structural condition. In addition, higher ductility is observed when the specimen is tested at slow deformation rates rather than under impact. The purpose of this note is to present evidence of the superplastic effect in a Cr-30 at. pct Co alloy. In addition, low-temperature bend ductility data are given for the alloy in the quenched, metastable, single-phased condition and in the slow-cooled, two-phased, near-equilibrium condition. Sixty-gram alloy ingots were prepared by noncon-sumable arc melting of iodide chromium and electron-beam-melted cobalt. These were cast into rectangular ingots and fabricated to 0.035-in. sheet by rolling at 2300°F. Attempts to roll at slightly lower temperatures invariably resulted in severe cracking, apparently due to presence of the brittle a phase below the transformation temperature of 2230°F. Tensile specimens with a 0.25 by 1.0 in. reduced section were cut from the sheet by electrical discharge machining. The specimens were heated at 2400°F for 1/2 hr to obtain a recrystallized, bcc, single-phased structure. Temperature of the specimens was subsequently raised or lowered to the desired test temperature and held for 10 min prior to testing at a crosshead speed of 0.01 in. per min. Fig. 1 shows the chromium-rich side of the Cr-Co phase diagram.6 The alloy composition investigated is near the point of maximum solubility of cobalt in chromium, and the alloy is two-phased when in equilibrium at temperatures below 2230°F. The results of tensile tests conducted at temperatures indicated by X's in Fig. 1 are summarized in Figs. 2, 3, and 4. Tensile elongations shown in Fig. 2 increase from 11 pct at 1900°F to a maximum of 160 pct at 2250°F. The elongation then decreases to 44 pct as the test temperature is further increased to 2600°F. The proximity of the temperature of highest elongation and the solvus temperature suggests that a superplastic effect is present in Cr + 30 at. pct Co similar to that observed for other alloys that show anomalous ductility. The variation of upper yield stress with temperature is shown in Fig. 3. A small peak in yield stress occurs at 2250"F, coincident with the maximum in

Jan 1, 1967

By N. S. Stoloff, R. G. Davies

The annealing behavior of a heavily cold-worked FeCo-V alloy has been studied at temperatures both above and below Tc, the critical temperature for ordering. It was found that re crystallization and grain growth can take place below Tc in the ordered state and that, in agreement with results for fcc alloys, order retards grain growth. The activation energy for pain growth is -55 kcal per mole in the disordered condition; it was not possible to assign a value to the activation energy for grain growth below Tc because the degree of order is changing rapidly over the temperature interval investigated. Precipitation of an fcc y phase was observed upon annealing the cold-worked alloy at temperatures of 675°C and below. SINCE it is known that diffusion constants are greatly affected by long-range order,' it is not unreasonable to anticipate that order will influence the diffusion-controlled phenomena of recrystallization and grain growth. In the course of a study of strain aging of ordering alloys around the composition Ni3Fe Vidoz, Lazarevic, and cahn2 noted that a disordered cold-worked alloy annealed at a temperature T1 just below Tc will order but not recrystallize; Tc is the critical temperature for order. The same alloy annealed at a temperature T2 just above Tc, or an alloy of different (nonordering) composition annealed at either T1 or T2, will recrystallize after equivalent amounts of cold work. Initially Selisskii and coworkers,3,4 who examined a series of cold-worked and annealed bcc Fe- Co alloys, reported that alloys forming the B2 FeCo super lattice could not be recrystallized below Tc. However in a later paper5 it was shown that FeCo ordered alloys could be recrystallized below Tc. Thus the first question to be answered by the present investigation was whether or not a heavily cold-worked Fe-Co-V alloy could be recrystallized below Tc. When it was found that the alloy would indeed re-crystallize below Tc, the influence of long-range order upon subsequent grain growth was studied. EXPERIMENTAL PROCEDURE The grain-growth measurements were made upon FeCo-2 pct V wire of 0.031 in. diam; the alloy was the same as that used in an earlier study of the mechanical properties of FeCO-v.6 The wire was produced by cold swageing and drawing a rod 0.125 in. diam that had been quenched into iced brine from 850°C to disorder the alloy; the initial grain size was 0.05 mm. Rates of grain browth were measured after isother-mally annealing the wire in the temperature range 675° to 825°C; T, is -720°C. Grain diameters were measured on the transverse cross section of the wires. The line-intercept method of determining the grain diameter was used; conventional light photomicrographs were taken when the grain size was greater than 10 . Carbon replicas were produced and examined in the electron microscope if the grain size was less than Dislocation configurations after recrystallization anneals were studied by electron transmission microscopy of foils produced by cold rolling a 0.125-in.-thick strip of the quenched material to 0.012 in. The foils, after heat treatments similar to the wire samples, were thinned by electropolishing in a chromic-

Jan 1, 1967

By Bengt Helmerson

A new tankhouse was started in January 1959 at the Copper Refinery of the Boliden Co., Skelleftehamn, Sweden. In early 1961 a semiautomatic machine for looping and straightening of starting sheets was introduced to substitute previous manual operations. The machine as well as consequent, racking and handling is described. Finally, economic benefits achieved and new constructions for the future are discussed. HISTORY Tankhouse operations at the Boliden Co. 's Smelter in Skelleftehamn started in 1932 on a scale that since then has been gradually increasing. In order to rationalize the operations, a new tankhouse was erected. It went into production in January 1959 and was designed for a yearly production of 35,000 metric tons at a current density of 16 amps per sq ft. Already from the beginning, however, the tankhouse has been run with a current density of 21.5 amps per sq ft, corresponding to a maximum capacity of 50,000 tons a year. During the time of planning of the new tankhouse, due consideration was given to the possibilities of mechanizing the operations. New ideas could be and were actually tested on full or pilot plant scale in the old tankhouse. In that way we developed a system for the cathode handling including washing, weighing, and transportation. A system using forklift trucks in combination with trailers and overhead cranes has been developed for transportation and racking of anodes and also for handling of anode scrap. With respect to the electrolyte, circulation and temperature control are partly automatic. Special attention was paid to the production of starting sheets, which operation normally requires a great number of man hours, and it was made the object of a particular survey as to the best ways of straightening starting sheets, attaching loops to the sheets, and so forth. As a result, a machine was designed during 1957-58 and built in our own shops. After a period of tests, the new equipment was put in operation commercially on the 10th of April, 1961. Since this date the machine has produced about 2 million starting sheets without interruption and has been the only source of looped starting sheets for the tankhouse. GENERAL DESCRIPTION Fig. 1 shows all the major parts of the system. These parts are: slitting machine with feeding table (I), standard roller leveler, type Ungerer (2), tilting roller table (3), conveyors for sheets and suspension bars (4, 7), storage of suspension bars (51, punching and stitching machine (6), two lift-turn-tilt tables (8) for discharging the starting sheets, two rack conveyors (9) for sheet assemblies. Included are a hy-daulic pump station and panels for pneumatic and electric controls. SLITTING AND STRAIGHTENING The machine is fed by hand by a man sitting in front of the slitting-machine table, Fig. 2. This operator inserts one sheet at a time into the slitter. After being trimmed, the sheet is automatically conveyed to the following step of handling, straightening. The slitting machine cuts away about 8 mm of the upper and lower parts of the sheet (in relation to its original position in the stripper section) for two reasons. The first one is to give the machine sheets of uniform dimensions; the second one is a matter of quality. Normally the starting sheets as stripped have a thick edge at the bottom end. This part would collect anode slimes and might also cause short circuits when inserted in the cells without trimming. The straightener used is of German construction

Jan 1, 1967

By E. D. Levine

A study was made to delineate the dislocation mechanistns controlling prismatic and basal slip at low temperatu.ves in titanium containing approximately 100 ppm interstitial impurities. Mechanical tests were performed on single crystals prepared from electron-beam zone-refined crystal bar by a thermal cycling technique. Measurements were made of critical shear stress and activation volume as functions of temperature. The study revealed that, for prismatic slip, two overlapping thermally activated processes aye rate-controlling below 300°K. From absolute zero to 220°K, reasonable agreement is obtained between experiment and the mechanism which involves overcoming 01. the Peierl's stress by the formation of double kinks in dislocation lines lying in close-packed directions. Between 220°and 300°K prismatic slip is controlled by a difierent thermally activated process which is characterized by a stronger temperature dependence of critical shear stress. Above 300OK, prismatic slip is controlled by an athermal process. Basal slip is controlled by a thermally activated process between 195' and 400°K. An approximate analysis suggests that it is the Peierl's mechanism that is rate-controlling in this range. Below 195°K, the occurrence of mechanical twinning obscures the features of basal slip. Above '400 X, basal slip is controlled by an athermal process. TITANIUM, an hcp metal with a less-than-ideal axial ratio of 1.587, deforms primarily by {1010}( 1210) slip at low temperatures.'-3 The critical shear stress is strongly affected by the presence of interstitial impurities. For example, churchman3 reported that a decrease in the combined nitrogen plus oxygen content from 0.1 to 0.01 pct decreased the critical shear stress for prismatic slip from 9.19 to 1.4 kg per sq mm at room temperature. Basal slip, however, is less strongly affected by interstitials; Churchman's data for {0001}( 1210) slip show a smaller decrease in critical shear stress from 10.90 to 6.3 kg per sq mm at room temperature, for the same decrease in interstitial content. In the present study, an attempt has been made to delineate the dislocation mechanisms controlling prismatic and basal slip in titanium. In view of the strong effects that impurities may have on at least some of these mechanisms, experiments were performed on material of reasonably high purity, so that comparisons could be made with investigations on material of both lower and higher impurity contents. EXPERIMENTAL PROCEDURE Single crystals were prepared from 1/4-in.-diam electron-beam zone-refined crystal bar, furnished by Materials Research Corp. The analysis of the zone-refined material is given in Table I, indicating that the interstitial content is of the same order as the purer of the two materials investigated by Churchman. The Vickers diamond pyramid hardness number of this material was 85.6. This value agrees quite well with the correlation between hardness and interstitial content proposed by Ogden and Jaffee.4 On the basis of hardness, the material investigated in this study is, therefore, intermediate in purity between that studied by Churchman (Vpn = 97) and that studied by Spangler and Herman (Vpn = 62.4).5 Single crystals were grown by thermal cycling through the a-ß transformation. Two-in.-long by 1/2-in.-diam bars were heat-treated in vacuum for 4 hr at 1160°C. The temperature of the furnace was then lowered to 865°C, and the specimens were maintained at this temperature for 7 days. This procedure was repeated several times, resulting in the growth of individual crystals up to 14 in. in length. Specimens for testing in direct shear along either {l010}(1210) or {0001}(1210) were obtained by electrical-discharge machining. The specimens were in the form of cylinders 1/4 in. in diam by approximately 1/2 in. in length, with a reduced section 1/8 in. in diam and a in. in length. All specimens were electropol-ished subsequent to electrical-discharge machining. The ends of the test specimens were cemented in stainless-steel collars. Shear tests were conducted in an Instron Tensile Tester at various temperatures from 4.2° to 508°K. Measurements were made of critical shear stress and activation volume. The latter measurements were performed by measuring the change in flow stress upon cycling the strain rate between ?1 = 2.62 x 10-3 per sec and ?2 = 2.62 x 10-4 per sec. Activation volume, I)*, is defined as u* =ßkT [11

Jan 1, 1967

By M. C. Flemings, T. Z. Kattamis

Dendrite morphology and grain size were studied in bulk samples of iron and nickel base alloys under-cooled up to 300°C. In the alloys studied, dendrite morphology gradually changes with increasing undercooling from the usual dendritic structure to a structure composed of cylindrical dendrite arms. At a critical undercooling of approximately 170°C in these alloys there is an abrupt transition to a fine-grained structure of spherical morphology. Final dendrite-arnz spacing of melts nucleated at less than the critical undercooling decreases with increasing undercooling and with decreasing distance from a chilled surface. Grain size of melts nucleated at greater than the critical undercooling decreases similarly with these factors. It is concluded that structure coarsening, with reduction of surface area as driving force, is the principal mechanism determining final dendrite-arm spacing in melts nucleated at small undercoolings and grain site in melts nucleated at large undercoolings. IN solidification of highly undercooled melts, nuclea-tion and growth of the solid is accompanied by a rapid release of latent heat of fusion. The heat release is sufficiently rapid in bulk specimens that it takes place essentially adiabatically with respect to the surroundings; hence the specimen heats up ("recalesces") to some maximum recalescence temperature. For bulk liquid metals and alloys, a simple thermal balance shows the undercoolings obtainable in bulk samples are not sufficient to lower the maximum recalescence temperature below the equilibrium solidus temperature. Hence, some liquid always remains at the end of recalescence; this liquid must solidify by rejection of heat to the surroundings. Metals, including iron and nickel base alloys, normally freeze with only a few degrees of undercooling. In general, nucleation begins on impurity particles present in the metal, at a "heterogeneous nucleation temperature" that is quite close to the equilibrium liquidus temperature. Very small droplets of metal can be undercooled considerably before nucleation occurs,1-4 presumably because subdivision of the liquid reduces the probability of solid impurity in a given droplet. Following early work on metal droplets,' a series of studies has shown that it is possible, without great difficulty, to undercool bulk samples of many metals and alloys. Undercoolings obtainable are approximately two tenths of the melting point for nickel, iron, copper, and other metals and alloys.5-9 Specimen sizes have generally been in the range of 50 to 500 g. Studies on undercooling to date have been concerned primarily with techniques of undercooling and with nucleation and growth kinetics. Only limited consideration has been given to dendrite morphology in under- cooled alloys. One such study, in bulk Cu-Ni alloy,8 showed that at small undercoolings a typical dendritic structure was obtained. At undercoolings the order of 175°C the structure of a Cu-40 pct Ni alloy became "ray or starlike", and a dense network of subbound-aries appeared. At still larger degrees of undercooling, of the order of 250°C, the structure became more random and included an ensemble of globular forms. Comparable results were obtained by walker10 on Cu-Ni alloys, and by Kamenetskaya et al.11 in research on Cu-Ni, Sb-Bi, Fe-Ni, and Fe-Cr alloys. In this work, study was made of the processing parameters influencing dendrite morphology and grain size in undercooled alloy melts. Experiments were specifically designed to show the effects of coarsening ("ripening") during solidification on final casting morphology. EXPERIMENTAL PROCEDURE Apparatus used was patterned after that of Walker,10 Fig. l(a). Melting was by induction under reduced pressure of helium. Two different types of crucible materials were used; there was no difficulty in obtaining substantial undercooling in both types of materials. Alumina crucibles were used for the bulk of the runs made, with a glass coating separating the metal from the crucible walls. Other runs were made using fused silica crucibles with or without slag coating. A graphite susceptor was employed around the crucible to aid in melting and to reduce melt stirring. The major portion of the experimental work was on

Jan 1, 1967

By R. E. Miner, J. K. Jackson, D. F. Gibbons

Internal-friction measurements were made during isothermal aging treatments of 18 pct Ni maraging-steel wires between 420° and 612°C. A sharp initial drop in the internal-friction level was interpreted in terms of a rapid dislocation recovery reaction. The rate of the recovery was found to be independent of temperature. Truncation of the recovery reaction was apparent in those systems that exhibited the rapid age-hardening characteristics of the maraging steels. It is proposed that this initial hardening response is the result of a low activation energy preprecipitation reaction involving molybdenum. A gradual rise in the internal-friction level at longer aging times was associated with the precipitation of a molybdenum-containing precipitate. CONSIDERABLE research effort has recently been devoted to the mechanisms responsible for the high strength and toughness of 18 pct Ni maraging steels. These bcc martensitic alloys are found to harden during aging at temperatures between 400° and 600°C. The aging response is very rapid with more than half the total hardness increase occurring within the first few minutes of aging, Fig. 1. Precipitate particles have been observed in the later stages of aging by a number of investigators.'-' While the identity of the precipitates remains the subject of some controversy, one or more of the intermetallic compounds, Ni3Mo, Ni3Ti, and Fe2Mo, are almost certainly involved. Heterogeneous nucleation of precipitates at dislocations has been noted8 and cobalt is believed to enhance the formation of fine, homogeneously distributed particles by lowering the solubility of the alloying elements in the martensitic matrix.3,7 Transmission electron microscopy studies of the initial hardening behavior have not been fruitful. Strengthening mechanisms based on atomic ordering have not been generally supported by X-ray and neutron-diffraction experiments.9 In the present investigation, the aging reaction was studied using isothermal-aging internal-friction measurements. The marked sensitivity of internal friction to short-range diffusion and precipitation processes has provided a convenient tool for the elucidation of the strengthening mechanisms operative in the maraging steels. EXPERIMENTAL PROCEDURE An upright torsion pendulum was used to measure the internal friction of several maraging steels during isothermal aging. The pendulum was placed in an atmosphere of helium at a pressure slightly less than atmospheric. To permit measurement of internal friction during the first minutes of aging, the specimens were rapidly heated by passing a current through them in a manner similar to that described by Ang and Wert.10 Thermal gradients in the specimen were largely eliminated by using specially designed specimen grips11 which produced temperature variations of less than 1 pct over 90 pct of the specimen length and less than 3 pct over the entire length at aging temperatures up to 620°C. A proportioning strip-chart program controller was employed to maintain constant specimen temperature during treatments of up to 36 hr duration. A number of vacuum-melted laboratory heats having compositions near that of a typical commercial 18Ni-7Co-5Mo-0.4Ti alloy were investigated. The amount of cobalt was varied from 0 to 7 wt pct in five of the heats. Either molybdenum or titanium was omitted in the remaining two heats. The alloy desig-

Jan 1, 1967

By Horace Pops

Martensitic transformation has been studied during cooling and heating in ß-phase Cu-Zn alloys to which small additions have been made of Ni, Ag, Au, Cd, Ga, In, Si, Ge, Sn, and Sb. The start and finish of the martensitic reaction and the variation of transformation temperature with third-element content were determined by electrical-resistivity measurements from alloys which had either constant zinc contents or constant values of electron concentration. All of the third elements, except nickel, lowered the transformation temperature if the results were plotted along the lines of constant zinc contents in each ternary system. A significant difference in the rate of lowering of the transformation temperature per atomic percent of the third element was observed for elements which had the same nominal valence. No systematic variation of transformation temperature with the valence of the third element was observed. It is suggested that the observed increase in transformation temperature for nickel-bearing alloys is due to the transfer of electrons from the conduction band of the alloy to virtual bound states. However, electron concentration is not the most important factor controlling the instability of the 0 phase. The transformation temperatures of the ternary alloys can be predicted from the following approxilnate expression: Ms (°K) = +3280 - 80 Zn + 8 Ni - 30 Ag - 12 Au - 140 Cd -90 Ga- 145 In - 80 Ge -175 Sn - 120 Si - 150 Sb MOST binary ß -phase alloys based upon the noble metals copper, silver, and gold are unstable at low temperatures and transform spontaneously by a martensitic reaction. This transformation has been studied recently in the ordered bcc ß'-phase Cu-Zn binary al1oys1,2 where the transformation temperature is below the room temperature and decreases with an increase in zinc content. It has been reported that the transformation temperatures can be raised above room temperature by small additions of a third element such as silicon3,5 or gallium,4,5 but no quantitative study has been made. The transformation temperature of different binary alloys can be altered by third-element additions. For example, it was shown that nickel and copper may have a large effect on the Ms temperature of CU-Al6 and Au-cd7 alloys. The present investigation was made to determine systematically the influence of various third elements on the martensite-transformation temperature of Cu-Zn ß-phase alloys. Since these alloys have an electronic origin,' alloy compositions were chosen so that the transformation temperatures could be determined at constant zinc contents or at constant values of electron concentration. I) EXPERIMENTAL PROCEDURE Ten ternary alloy systems were obtained by adding nickel, the noble metals silver and gold, and some B-subgroup elements to a Cu-Zn matrix. These are arranged according to their rows and columns in the Period Table as follows: Each ternary alloy was prepared by melting and casting weighed quantities of high-purity metals (99.99+ pct) in sealed quartz tubes under a partial pressure of helium to make a 4-g ingot. The molten alloys were shaken vigorously and then quenched in water. Since the weight loss was negligible, the compositions of the ingots after casting and annealing were assumed to be the same as the nominal compositions. The ingots were homogenized after casting in helium-filled Vycor tubes for 24 hr at temperatures between 750" and 810 C and quenched into brine. Metallographic examination revealed that all alloys were homogeneous, poly crystalline ß-phase alloys, and that the grain size was in the range 1 to 5 mm. Electrical-resistivity measurements were made to determine transformation temperatures of the ternary alloys during continuous cooling or heating. Transformation temperatures of the ternary alloys can be determined by electrical-resistivity measurements since the resistance of the martensitic phase is much higher than that of the 0' phase. The technique has been described previously in connection with a study of Au-Zn alloys.9 The reproducibility of transformation temperature was approximately ±6°C. II)RESULTS A hysteresis was always observed in electrical-resistivity curves and was usually less than about 12°C.

Jan 1, 1967

By R. F. Hehemann, S. T. Zegler

The w phase in zirconium alloys containing more than G pct Nb can form in a difjUsionless manner during quenching or with composition change during aging at temperatures below 550°C. The latter treatment establishes a metastable equilibrium between an w phase containing 5 pct Nb and a ß phase containing 45 pct Nb. The superconducting transition temperat~ire of quenched Z?-10 pct Nb and Zr-15 pct Nb alloys is increased by aging at 375°C and both alloys exhibit the same Tc after prolonged aging. The superconducting transition both in the quenched and in the aged alloys is relatively sharp (occurring ovev a temperature range of <0.5°K), and is interpreted in terms of the microstructure that characterizes the ß-w) mixture and proximity effects postulated for thin-film superconductors. At elevated temperatures alloys of the transition elements from groups 4 to 6 in the periodic table exhibit wide ranges of mutual solubility in the bcc structure. Except for compositions that are high in a group 4 element? the bcc structure can be retained at room temperature by quenching and several systematic studies of the influence of composition on superconductivity in bcc alloys have been reported.&apos;&apos;2 The high-temperature bcc form (ß) of the group 4 elements—titanium, zirconium, hafnium—is not retained upon quenching but instead transforms marten-sitically to the hcp structure (a). The Ms temperature can be lowered by alloying the group 4 metals with a solute from a higher group and with a critical amount of solute the martensitic transformation can be suppressed. The 3 phase, however, is not retained without modification. Rather, a metastable phase, w (in particular in titanium- and zirconium-base alloys), occurs over a fairly wide range of compositions above that required for suppression of the martensitic transformation.3,4 As in the formation of a, the transformation of ß to w occurs in a diffusionless manner during quenching and the temperature for its initiation (ws) is lowered by increasing solute content. The ws temperature can be well below room temperature and reversibility of the transformation allows w that forms on cooling to low temperatures to disappear again when the sample is returned to room temperature.5 In Zr-Nb alloys the w phase can also be formed during aging at temperatures below 550°c.5 In this treatment, a metastable equilibrium is established between a niobium-rich ß phase and a zirconium-rich w phase. The present study of Zr-Nb alloys was undertaken to examine the influence of microstructure: as controlled by composition and heat treatment, on the superconducting transition temperature, Tc. MATERIALS AND PROCEDURE The alloys were prepared from zirconium crystal bar that was spectrochemically pure except for 30 ppm of Al, 40 ppm of Cu, 100 ppm of Fe, and 1 ppm of each B and Pb. The niobium used was 99.5+ pct pure containing according to the suppliers&apos; (Stauffer Chemical Co.) analysis a maximum of 50 ppmO, 60 ppm of N, 30 ppm of C, 5 ppm of H, 200 ppm of Ta, 50 ppm of Mo, and 15 ppm each of Fe, Co, and Ni. The alloys were prepared in the form of buttons weighing approximately 4 g by arc melting on a water-cooled copper hearth in an inert-gas atmosphere. All alloys were heat-treated in the ß-phase region for 3 days at 1000°C and water-quenched. For heat treatment, the specimens, 0.5 cm in diam by 1.0 cm long, were wrapped in zirconium foil and sealed in an inert-gas atmosphere inside quartz capsules. Quenching was done by breaking the capsules under water. Subsequent aging treatments of Zr-10 pct Nb and Zr-15 pct Nb alloys were done by placing the specimens in a lead bath at 375°C and quenching in chilled brine. The phases present in the quenched and aged alloys were identified by optical and X-ray metallography. X-ray powder patterns were taken at room temperature with a 114.6-mm-diam Debye-Scherrer camera and CuK-a radiation. Superconducting transition temperatures were determined from magnetic-permeability measurements6 in a 10 oe field. The superconducting transition temperature was determined from a plot of galvanometric deflection as a function of temperature by extrapolation of the linear portion of the curve to zero deflection. In Figs. 1 and 3 the arrows indicate the temperature range of the transitions. RESULTS AND DISCUSSION Superconductivity in Quenched Alloys. In Fig. 1 Tc for alloys quenched from 1000°C is plotted against nominal composition. Indicated also in Fig. 1 are results of metallographic studies. Alloys containing less than 5 pct Nb transform martensitically to a during quenching from 1000°C and Fig. 1 demonstrates that Tc increases significantly as the niobium concentration in a is increased. On the basis of the Bardeen-Cooper-Schreiffer theory of superconductivity7 it seems most likely that this increase in Tc results from an increase in the density of electronic states as zirconium is alloyed with niobium. This increased density of states has been revealed clearly in low-temperature specific-heat measurements.8 The martensitic transformation to a is suppressed when the alloy content exceeds 5 pct Nb. In the range from approximately 6 to 30 pct Nb, quenched alloys consist of a mixture of B and w phases and the amount

Jan 1, 1967

By M. Nicholas, D. M. Poole

The diameters of sessile drops have been found to increase linearly with time in five reacting binary metal systems. The spreading rates of the drops are markedly dependent on temperature and on prior alloying of the solid with the lower melting point metal, hut are independent of the drop volume, wetting atruosphere , solid-surface roughness, and prior alloying of the drop with the substrate metal. A mechanism has been suggested that relates the linear-spreading rate to lateral diffusion of the liquid-metal atoms into the solid at the drop edge. An Arrhenius- type equation has been derived that describes the temperature dependence 0) the spreading rate, and although the agreement between the actual and the predicted pre-exponen-tial terms is poor that between the activation energies is excellent and the variation in the spreading rate of copper on Ni-Cu alloys produced by different extents of alloying can be predicted with considerable accuracy. CHEMICAL interactions frequently change the wetting behavior of solid-liquid systems causing, for example, "secondary spreading1 of sessile drops beyond the size defined by the surface and interfacial tensions of the unreacted components. The kinetics of the contact-angle decreases associated with this spreading are similar for many systems, but few studies have been made with the objective of determining whether the similarities are a reflection of a common mechanism. Some workers2,3 have assumed the secondary spreading is controlled by changes in the liquid surface and liquid-solid interfacial tensions and hence by the composition of the liquid, and contact-angle changes measured by the vertical-plate technique have been used to follow the course of liquid-solid chemical reactions.4 Other processes that have been invoked to explain these time-dependent changes in specific systems include the removal of adsorbed gas from the liquid-solid interface.5 penetration of containment layers on the solid Surface,6 interdiffusion,1,7 reori-entation of the solid surface into a wettable configuration: vapor-phase transport of the liquid onto the solid in advance of the drop,9 and, from vertical-plate studies. capillary flow between oxide layers and the solid surface.10 One of the reasons for the profuseness of these suggestions may be the complexity of the contact-angle change kinetics. However, in an analysis of secondary spreading gold and copper on UC,11 it was found that the diameter of the contact area between the sessile drop and the solid surface showed a simple linear increase with time although contact-angle changes were more complex. To check whether the linearity was merely fortuitous! additional exploratory work was conducted with four reacting metal-metal systems: Au on Ni. Cu on Ni, Cu on Fe, and Ag on Au. Linear spreading was observed in every case even though the kinetics of the contact-angle changes were complex. A further detailed study of the kinetics of linear spreading of five reacting metal-metal systems has been made with the object of determining the mechanism involved. The influence of variables such as temperature, drop volume. and the initial composition of the drop on the linear-spreading rate has been measured and compared with those predicted by a number of possible mechanisms. The systems employed in this study (Cu and Au on Ni and Pt, and Ag on Au) were selected because of the availability of potentially relevant chemical and physical property data. the simplicity of their phase diagrams at the wetting temperatures, and the ease of experimentation. EXPERIMENTAL TECHNIQUES The purities of the metals used in the study were: copper, 99.9 pct; gold. 99.96 pct; nickel, 99.2 pct; platinum 99.99 pct; and silver, 99.999 pct. The wetting tests were performed in a split tantalum tube vacuum resistance furnace of a conventional design. The furnace element was held vertically and was 1 $ in. in diam and 6 in, long. Viewing ports were provided in the water-cooled chamber to enable the specimens to be observed in both the horizontal and vertical planes. The temperature in the hot zone of the furnace could be held at 1500" i 5°C for an indefinite time. The surfaces of the solid-plaque metals were ground flat on Microcut paper and both the sessile drop and substrate metals were ultrasonically cleaned in methyl alcohol prior to their insertion in the furnace. After loading, the furnace was pumped down to a pressure of 2 x 10-5 mm of mercury and degassed for 30 min at 900° to 950°C. The temperature was then increased at more than 100°C per min to that used in the wetting test. The vacuum at the wetting temperature was better than 5 x 10-5 mm of mercury. Dewetting and retraction of the drop on cooling did not occur and the contact-area diameters, therefore, were measured after solidification with a vernier traveling microscope. The diameters quoted later are arithmetic means of ten measurements. The standard error of the mean never exceeded 3 pct and was often less than 1 pct.

Jan 1, 1967

By P. H. Turnock, R. D. Pehlke

The effects of mutual interactions between elements on the solubility of nitrogen in liquid iron alloys containing chrormium, columbium, molybdenurn, nickel, or silicon have been determined. The equilibrium nitrogen solubilities in the liquid iron alloys were measured by the Sieverts method. The second-order effects caused by the presence of chrormium, columbium , or silicon in the melt were found to be significant. The solubility of nitrogen in liquid iron alloys containing several alloying elements has also been measured as a function of melt temperature and nitrogen pressure. The heat of solution of nitrogen in microcomponent iron alloys has been found to be a function of the logarithnz of the activity coefficient of nitrogen, irrvespectilje of the composition of the melt. The heat of solution of nitrogen in pure liquid iron was determined to be 1265 ± 450 cal per g-atom of N over the temperatxue range 1600° to 1800°. Sieverts &apos; law was obeyed for all melt compositions studied in the pressure interval 0.4 to 1.0 atm. The solubility of nitrogen in liquid pure iron and in many dilute binary liquid-iron alloys has been the subject of many investigations. Pehlke and Elliott&apos; made a comprehensive investigation on this subject in which they summarized and compared previous researches. However, the practical application of these solubility studies requires prediction of the nitrogen solubility in the multicomponent alloys encountered in operating practice. The interaction parameter as defined by wagner2 has been used to predict the behavior of a solute in a complex alloy, and a modification of this approach has met with reasonable success in several instances.&apos;1"5 The necessity for extending the nitrogen solubility data on binary alloys is caused by the lack of direct measurements of the solubility of nitrogen in molten alloy steels. A direct comparison between measured solubilities and those predicted for multicomponent alloys based on data in binary systems has been made in only a few instances, and most of these were reviewed by Langenberg,6 who developed a graphical technique for predicting nitrogen solubilities in multicomponent iron alloys. The solubility of nitrogen has also been studied over the entire Fe-Cr-Ni ternary system by Humbert and Elliott.? Both studies compared the solubilities computed from binary iron alloy data with those measured in ternary alloys in the liquid state at 1600°C. Most of the nitrogen solubility measurements referred to above were for iron alloys at 1600°C; thus, the ability to predict nitrogen solubilities in molten iron alloys has been restricted primarily to alloys at 1600°C. Recently, Nelsson8 devised an approximate method for calculating the change in nitrogen solubility in molten iron alloys as a function of temperature, but an assumption of regular solution behavior limits the applicability of the method. Chipman and Corrigan9,10 have developed an empirical correlation. based on the activity coefficient of nitrogen in the iron alloy at 1600°C, which facilitates the estimation of nitrogen solubilities at higher temperatures. This method also has limitations, due to the necessity of calculating the activity coefficient for nitrogen in a multicomponent iron alloy at 1600°C from the nitrogen solubility data for binary iron alloys. The shortcoming of such a straightforward approach to the problem of predicting nitrogen solubilities is the disregard of mutual interactions between elements in a complex iron melt. This investigation was conducted in an attempt to provide a more reliable method of predicting nitrogen solubilities in multicomponent iron alloys. The effects of the mutual interactions between chromium, nickel, silicon, molybdenum, or columbium in iron were studied, using the Sieverts&apos; technique1,7 to measure the effects of alloy composition, temperature, and pressure on the solubility of nitrogen in ternary liquid iron alloys.

Jan 1, 1967

By V. O. Struebing, K. A. Johnson, F. H. Ellinger, C. C. Land

The system Pu-Ce-(Pu,Ce)Co2 has been investigated by thermal, micrographic, and X-ray diffraction methods. The principal characteristics of this system were found to be: 1) a "eutectic" valley extending across the phase diagram from the 6 Pu/Pu3 Co eutectic at 408°C to the 7 Ce/Ce24Co11 eutectic at 42Z°C, the maximum temperature in the valley being 443°C; 2) three ternary inter mediate phases, each having an appreciable mutual solid solubility for plutonium and cerium; 3) eleven four-phase equilibria comprising one eutectic, four peritectics, and six Class II reactions; 4) a pseudobinary section between the ternary phase A and the y Ce solid solution, and another such section between the 6 Pu solid solution and the 77 phase (Pu3 CO) containing cerium in solid solution. ALLOYS of plutonium that are completely molten at temperatures in the vicinity of 500°C are of interest for use in the liquid state as fast reactor fuels. Since exploratory work1,2 on Pu-Ce-Co alloys had indicated that a "eutectic" valley extended across the ternary system between the Pu-co[6 Pu/n (Pu3co)] and the Ce-CO[? Ce/? (ce24co11)] binary eutectics, thus allowing alloys of any plutonium concentration to be selected within the temperature limitation, the present investigation was begun for the purpose of accurately defining the "eutectic" valley. However, when the existence of three ternary intermediate phases and the consequent complexity of the liquid-solid equilibria were discovered, the investigation was expanded to cover the complete system, except for the corner between the isomorphous series PuCo2-CeCo2 and pure cobalt. 1) EXPERIMENTAL The purities of the alloy components are given in Table I. Alloy buttons, each weighing about 5 g, were prepared by arc melting the accurately weighed components (±0.1 mg) under a purified He-Ar atmosphere. Weight loss as a result of melting was usually less than 1 part in 1000; therefore, the buttons were not chemically analyzed. One half of each button was re-melted by induction heating in vacuo in a magnesia crucible and furnace-cooled, so as to produce a somewhat coarser microstructure than that of the rapidly cooled arc-melted button. About 20 g of each alloy to be examined by thermal analysis were remelted in vacuo in a tantalum crucible of the inverted thermocouple well type. The thermal, X-ray diffraction, heat treating, and micrographic methods used in this investigation were essentially the same as have been described in earlier reports.3,4 Two etchants for these alloys were found to be satisfactory: A) 8 (parts by volume) orthophosphoric acid, 5 ethanol, and 5 glycol; B) 7 tetraphosphoric acid, 36 water, and 57 2-ethoxyethanol. Etchant A was used at 5 to 6 v for 15 to 20 sec, and etchant B was used at 10 v for 10 sec. 2) THE LIMITING BINARY SYSTEMS 2.1) Plutonium-Cerium. The phase relationships of this system have been investigated by Ellinger, Land, and Cramer,5 and by Selle and Etter.6 The over-all agreement of their phase diagrams is good, but there are some discrepancies in the positions of some phase

Jan 1, 1967

By T. G. Digges, C. L. Vold, M. R. Achter

A study was made of the effect of the growth conditions on the perfection of single crystals of niobium (columbium). Dislocation densities, determined by means of double-crystal diffractometer measurements , were not greatly affected by the method of crystal preparation but could be reduced by annealing treatments. However, the size, sharpness, and tilt angles of the substructures, observed with X-ray reflection macrograph, were sensitive to variations in growth procedures as well as to subsequent thermal treatments. Although the dislocation density was the same in both types, there were more low-angle bound-aries in crystals grown by zone melting than in those prepared by strain anneal. Mechanisms to account for these observations are discussed in terms of dislocation movements. A planned study of the structure-sensitive properties of refractory metals required the use of single crystals of a high degree of structural perfection and, for ease of handling, of large cross section. It appeared that the strain-anneal technique could satisfy both of these requirements. First, crystals grown in the solid state have been reported to be more perfect than those obtained from the melt.&apos; Second, the diameters of rods which may be produced by zone melting should have a theoretical limit determined by specific gravity, thermal conductivity, and surface tension, while the diameter of strain-annealed rods is limited only by practical considerations. Previously it was shown that niobium (columbium) single crystals of 1 in. diam2 may be grown by strain anneal, compared to the 0.5 in. maximum diameter achieved by zone melting, as reported for molybdenum by Belk.3 The current research was undertaken to investigate and optimize the effect of various process variables on the perfection of 1/4 and 1/2-in.-diam niobium single crystals grown by strain annealing and to compare their perfection to those grown by zone melting. Characterization of these crystals was more conveniently accomplished by means of X-ray than by metallo-graphic techniques. EXPERIMENTAL PROCEDURE Specimen Preparation. Zone-melted crystals of 1/4 in. diam were produced by the standard electron-beam zone-melting technique. The swaged and cleaned rods were outgassed, in the solid state at a temperature near its melting point, at a rate of 12 in. per hr, and single crystals were grown by making two molten passes at 2 in. per hr. By maintaining a zone length of 4 to % in., very uniform single crystals several inches long were obtained. For the strain-annealed crystals, an induction heater was used, in preference to other types of heating, to take advantage of the good penetration of large sections. A five-turn coil, 1 in. long, operating at 10 kc and powered by a motor generator, was contained in a vacuum chamber. The rod, suspended from the upper end, was raised through the coil for both recrystallization and crystal growth. In preliminary work single crystals of the same material were also grown with single and multiturn coils powered by a 450-kc generator. A vacuum of 2 X 10-6 Torr was maintained at temperatures up to 2400°C. Starting with electron-beam-melted ingots of 21/2 in. diam, the analysis for which is given in Table I, the material was cold-swaged to the desired cross section of 1/4 and 1/2 in. diam and then recrystallized. The rods were then strained in a tensile machine and converted to single crystals by passing through the induction coil. As with zone melting, control of orientation is possible by the use of special procedures. Other investigators, see for example Williamson and smallman,4 have reported that orientation control may be achieved by a bending technique. In the present work the strained rod is partially lowered through the coil to start the growth of the crystal. Then it is removed and bent at a point in the poly crystalline portion. Finally, it is returned to the chamber and growth is continued "around the corner". This procedure has certain limitations. If the bending operation exceeds the critical strain, recrystallization may take place. Also, the amount of bending which can be imparted to the rod is limited by coil geometry, and up to now has been 10 deg. However, by repeating the bending and growing operations it should be possible to attain any desired orientation. In preparation for X-ray examination, single crystals were sectioned and planed by means of the spark-erosion technique. To obtain the maximum reflected intensity, the (110) plane was exposed for examination. They were then etched 3 to 5 min in a mixture of con-

Jan 1, 1967

By T. P. Floridis, J. H. Young

The need for close control of the phosphorus content of steels has led to numerous investigations on the equilibria of the dephosphorization reactions. Winkler and chipman1 have established the general conditions for effective dephosphorization of steel. They are high slag basicity, high oxygen potential, and low temperature. Other investigators have made additional contributions to the understanding of the dephosphorization process. The current status of the understanding of the dephosphorization of steel is concisely presented by Bodsworth2 and by Ward. The investigation reported in this communication was undertaken with the purpose of establishing the effect of additions of barium oxide and calcium fluoride on the dephosphorizing capacity of slags. Esin and Gel&apos;d4 proposed that barium oxide, being more basic than calcium oxide, should cause an increase in dephosphorizing capacity when added to steelmaking slags, or when used as a substitute for calcium oxide. Derge5 has also proposed that in the future conventional slags might be replaced by BaO-Al2O3 slags. There is, however, no experimental evidence confirming the effect of barium oxide on the dephosphorizing capacity of slags. The effect of calcium fluoride on the dephosphorization of steel is not clearly understood. It is generally recognized that additions of calcium fluoride are beneficial. It is not clear, however, whether calcium fluoride affects the equilibrium of the dephosphorization reaction, or whether it simply causes an increase in the fluidity of the slag and, consequently, faster approach to equilibrium. The experimental procedure consisted in equilibrating synthetic molten slags with liquid copper at 1550°C under a gas stream containing argon, hydrogen, and water vapor. In all experiments the argon to hydrogen ratio was approximately 4:1, and the hydrogen to water ratio was 5.42:l. Molybdenum crucibles were used as containers for the slag and metal. Under the above-described conditions of temperature and composition of atmosphere, there was no observable attack of the crucibles by the metal, slag, or atmosphere. Copper was used instead of iron, because iron attacks molybdenum. The equilibration was made in a tubular furnace consisting of a recrystallized alumina tube. The alumina tube was heated by electrical resistance. A Pt-40 pct Rh wire winding was used for most runs. A silicon carbide tubular resistor was also used for some runs. Temperatures were measured with a Pt-Pt-Rh (10 pct Rh) thermocouple and kept constant within ±5°C. Equilibrium was approached from both sides, i.e., by adding the phosphorus either as oxide in the slag or as a phosphorus-rich alloy of phosphorus and copper. The holding time at the equilibrium temperature was 6 hr. At the end of each run the crucibles were rapidly cooled and removed from the furnace. The slag and metal were separated and analyzed. The experimental results are shown in Table I. The phosphorus content of the slag is expressed both as percent phosphorus pentoxide and as percent phosphorus. The basicity ratio is computed by dividing the number of moles of basic oxides-oxides of barium, calcium, magnesium, and sodium—by the number of moles of acidic oxides- oxides of aluminum, phosphorus, and silicon. Calcium fluoride is not included in the computation of the basicity ratio; i.e., calcium fluoride is assumed to be neither basic nor acidic. The distribution ratio of phosphorus—percentage of phosphorus in the slag divided by the percentage of phosphorus in the metal- is plotted in Fig. 1 against the basicity ratio. The results indicate that slags containing barium oxide have greater dephosphorizing capacity than slags containing calcium oxide. The high dephosphorizing capacity of slags containing sodium oxide and the low dephosphorizing capacity of magnesia-containing slags which already have been reported in the literature2 are confirmed by the results of this investigation. It appears that calcium fluoride has a beneficial effect on the distribution of phosphorus between slag and metal in acid slags only. Although the obtained distribution ratios between the phosphorus contents of slag and copper are not directly applicable to the dephosphorization of steel, they are sufficient for evaluating the effect of slag additions on the dephosphorizing capacity of slags in general. An increase in the ratio of distribution of phosphorus between slag and metal indicates lowering

Jan 1, 1968

By A. K. Jena, M. O. Bever

The thermodynamic behavior of dilute solutions of copper, silver, and gold in liquid tin is not explained by current solution theories.1-3 The relative partial gram-atomic enthalpies, or heats of solution, of silver, gold, and platinum in liquid lead at 623°K have been measured in the investigation reported here. Their interpretation requires a modification of the regular-solution theory although lead-rich solutions may be expected to be simpler than tin-rich solutions. The liquid metal solution calorimeter used has been described.4 At the start of a run the bath consisted of 621 g of reagent-grade lead. Several calibrating additions of lead were made during a run; a value of 3.48 kcal per g-atom5 was taken for H623°K - H273°K Two runs were made with each of the solutes silver, gold, and platinum. Annealed wire samples of 99.999+ pct purity were used. Four to six samples each weighing 0.5 to 1,0 g were added from 273°K to the bath at 623°K and the associated heat effects were measured. The samples dissolved in approximately 3 min. No serious experimental difficulties were encountered. The heat effects of the additions in a run plotted against the average atom fraction of solute before and after each addition314 were linear functions of the concentration in lead of silver, gold, and platinum, respectively. By subtracting the difference between the heat contents of a solute element at 623° and. 273°K5 from the measuredheat effects, the partial gram-atomic enthalpy relative to the solid solute was obtained. The relativ: partial gram-atomic enthalpy at infinite dilution and the concentration dependence of the partial gram-atomic enthalpies obtained from the plots of vs atomic fraction x are listed in Table I. The relative partial gram-atomic enthalpy at infinite dilution of silver in lead of 5.46 kcal per g-atom at 623°K compares with values at 723°K of 5.74&apos; and 5.765 kcal per g-atom. These data yield a temperature coefficient of 3.0 cal per g-atom-OK. The partial gram-atomic enthalpy at infinite dilution of gold in lead at 623°K is 0.89 kcal per g-atom compared with 0.85 kcal per g-atoms (obtained by extrapolation) and 0.95 kcal per g-atom.5 By combining these values for 623°K with those for several other temperatures5" the temperature coefficient of was found to be 3.2 cal per g-atom-OK. No thermodynamic data are

Jan 1, 1968

By G. Y. Chin, M. T. Dolan, W. L. Mammel

We have obtained by computer methods the solutions of the Taylor analysis1 for axisymmetric flow in bcc metals. Four modes of slip have been treated in detail:2-4 (111), {112}(111), {123}( 111), and a mixture of all three. In addition to obtaining anisotropy of strength data for randomly oriented or textured polycrystalline material,4 we have also obtained the lattice rotation resulting from slip on the predicted systems. The patterns of lattice rotation are essential to an improved understanding of texture development during deformation. The present note shows the results of lattice rotation for 5 pct axisymmetric compression, for fifty-two axial orientations distributed at 5-deg intervals throughout the standard stereographic triangle. The data are presented in Figs. 1 to 4 for the four modes of slip. The lattice rotation is indicated by a vector, the ends of which correspond to the positions before and after the deformation. The different vectors emanated from a given orientation represent different combinations of five slip systems that satisfy Taylor&apos;s minimum shear criterion.&apos; Fig. 1 shows the results for {110}< 111> slip. They are similar to those obtained by Bishop5 for the case of (111 )(110) slip in axisymrnetric tension, since in axisymmetric flow the sense of lattice rotation for {111)( 110) in tension is the same as that for (110) (111) in compression. The pattern of lattice rotation may be divided into three regions: region A in which the rotation is toward [100]; region B, toward [lll]; and region C, toward either [ loo] or L lll]. Although some slip combinations in region C lead to [1101, the latter is not stable as continued activity of the same combinations will eventually move the axis toward [loo] or Llll]. It may also be noted that [loo] is itself not stable. However, close inspection of the vectors in region A shows that it is impossible for an axial orientation within A to rotate outside its boundary. Hence after heavy compression, all orientations from A and some from C are expected to form a diffuse band in the [100] vicinity. By far the largest region is B. Here all orientations rotate to [lll], which is a truly stable position. Thus the [111] texture component is expected to be rather sharp. For the case of (112 }( 111) slip, Fig. 2, region C is enlarged at the expense of A, with B about the same as for {110}( 111) slip. A significant difference, however, is that [loo], as well as [Ill], is now truly stable. Fig. 3 shows the results of slip on {123}( 111) systems. Region B is substantially enlarged at the expense of C, with A being intermediate in size between the {110}(111) and {112}(111) cases. Like the case of {110)( 111) slip, [100] is itself unstable. However, the lattice rotations here are confined to less than 5 deg; hence it may be considered "stable" for practical purposes. Finally, Fig. 4 shows the results of mixed slip. Here region B is the largest, and C the smallest, of all the cases considered. Both [loo] and [ill] are stable. The behavior of the three corner positions is, in fact, identical to the case of {112)( 111) slip, Fig. 2; it turns out that in these positions for the mixed slip case, all the active slip systems (based on Taylor&apos;s criterion) belong to pure {112}( 111) .4 In summing up the results for the four modes of

Jan 1, 1968

By W. A. Tiller, W. C. Johnston

A high-temperature electromagnetic stirrer is described in which heating and stirring are accomplished by independently controlled power sources. The appavatus is suitable lor use at temperatures up to 1700°C in a variety of ambient atmospheres. Some typical examples of the homogenizatimz capabilities of the system are given. THERE are few processes in solidification that are not markedly affected by motion in the melt during freezing. In many instances, the mechanisms are diffusion-controlled, and the transport in the melt may be greatly accelerated by deliberately stirring the melt. In zone-refining, stirring1 assists the removal of rejected impurities from the interface, so the process proceeds at a faster rate. The transition from a planar to a cellular interface is caused by constitutional undercooling in the melt ahead of the interface: and stirring delays its onset. Stirring is valuable for homogenization of melts: and chemical reaction with sluggish kinetics may be accelerated. Finally, it has been observed that grain refinement is related to motion in the melt. Fine grain castings are usually produced by the addition of catalysts to the -melt,&apos; catalysts which are thought to act simply as hetereogeneous nucleation centers. Even here motion is important. Richards and Rostoker 5 applied ultrasonic vibration to a solidifying A1-Cu alloy which had been innoculated with a catalyst and found that the grain diameter fell linearly with the amplitude, the peak acceleration and the power input to the melt from the transducer. Finally, mechanical and electrical stirring alone have been used to generate a fine-grained structure.6,7 Johnston ef a1.&apos; have carried out a series of systematic investigations of grain refinement by electromagnetic stirring in a number of low melting point alloys. They found, for example, that the number of grains per unit volume in Pb-Sn alloys could be increased several orders of magnitude by stirring an undercooled melt at the moment of recalescence. In general, a relation AT .H = constant prevailed for a given grain size, where AT was the undercooling of the melt and H the field strength. In more recent work, deliberate homogeneous nucleation of slightly undercooled melts established that the mechanism of refinement must be one involving crystal fragmentation and subsequent multiplication, rather than a "shower" of nuclei effect.9 It is the purpose of this note to describe a stirring device suitable for use up to 1700°C. At low temperatures mechanical stirring and direct-current methods are feasible, but at high temperatures the problem of a protective atmosphere and of electrode corrosion rules out such procedures. The most convenient method for high temperatures is to use externally generated ac fields for both stirring and heating. With rf induction heating alone, considerable stirring and agitation can be achieved, but in general the penetration of field into the melt is small, and the stirring cannot be controlled independently of the heating. In the present experiments, separate power sources of different frequencies for heating and for stirring were used. A susceptor design was chosen so that the 450 kc rf heating field was completely absorbed in the susceptor. The stirring frequency, 400 cps, hereafter called the af field, was chosen so that a high penetration of the melt proper was achieved. EXPERIMENTAL APPARATUS The apparatus, Fig. 1, consists of a quartz tube and end plates, surrounded by an rf induction coil and six equally spaced af stirring coils, four of which are shown in full and a fifth in section. Each af stirring coil is a transformer of which the secondary is a single-turn water-cooled copper loop and the primary is composed of two 10 amp-117 v Variac cores as shown. These cores are cooled by forced air, as each of the six pairs will carry maximum currents of 15 amp for short periods. Each set of Variac windings are connected in series, but opposite sets are connected in parallel with a three-phase 400 cps 400-v source. By properly phasing the coils in this way, a rotating field is produced. Capacitors C1, C2, and C3 in Fig. 2 are used to match this inductive load to the generator. Fig. 3 shows a cutaway view of the quartz tube. The sample (1 in. diam by 1 in. high) is placed in a tapered alumina crucible. An axial W-26 pct Re thermocouple, enclosed by a protection tube, is provided. The cruci-

Jan 1, 1968

By C. E. Brukl, E. Rudy, St. Windisch

The ternary-alloy system Nb-Mo-C was investigated by means of X-ray, melting point, DTA, and metallo-graphic techniques; a complete phase diagram for temperatures above 1500°C was established. Above 1960°C, niobium monocarbide and the cubic (Bl) high-tenzperature phase in the Mo-C system form an uninterrupted series of solid solutions. The ternary range of the pseudocubic q MoCl-, is very restricted. Dimolyb-denum carbide dissolves up to 44 mol pct Nb2C (2240°C), whereas the maximum solid solubility of MO2C in Nb2C does not exceed 5 mol pct. The order-disorder transformation temperatures in Mo2C and Nb2C are lowered by the mutual metal exchanges. Six invariant (p = const) reactzons occur in the ternary system; three correspond to class 11-type four-phase reactions involving a liquid phase, one to a class I (eutectoid)-type, and two further isotherms are associated with limiting tie lines. The results of the Phase diagram investigation are discussed, and the thermodynamic interpretation identifies the low relative stability of the binary sub-carbides in conjunction with the large stability diflerences between niobium and molybdenum carbides as the cause for the formation of a stable equilibrium between the monocarbide and the metal phase in the ternary reson. Due to their refractoriness, the carbides of the high-melting transition metals have received increased interest in recent years as base materials in composite structures for aerospace applications at high temperatures and for the development of self-bonded cutting tool materials; other novel fields of application include power reactors, where operation at high temperatures becomes essential for attaining high power efficiencies. In these applications, the increased reaction rates at high temperatures require a close consideration of the chemical interactions between the alloy constituents. As a consequence, a detailed knowledge of the phase relationships in the alloy systems is required in order to provide a sound basis for developmental -type work. Partly as a result of the considerable experimental difficulties associated with the investigation of this high-melting alloy class, no complete studies of ternary metal-carbon systems have been performed until recently. Even the high-temperature phase relationships in the binary transition metal-carbon systems have been delineated only during the past few years to a degree of accuracy required for a more detailed study of ternary or higher-order alloys. In recent investigations of binary and ternary systems of refractory transition metals with carbon, boron, and silicon,&apos; alloys from the ternary systems Nb-Mo-C became of interest because of the demonstrated possibility2,3 of obtaining compatible composites based on metal + monocarbide combinations. In the meantime, however, studies in other, but related, ternary metal-carbon systems, such as Ta-W-C, have shown that the solid-state equilibria may change significantly toward higher temperatures (>2000°C), and that extrapolations based on low-temperature equilibrium data are, in general, not very reliable. Although the lower-temperature (<2000°C) phase relationships in the Nb-Mo-C system are similar to those found in Ta-W-C, a cursory thermodynamic analysis of the equilibria indicated4 that complete solid-solution formation between Mo2C and Nb2C should not occur at higher temperatures. The present work was conducted in order to experimentally verify these expectations and, in addition, to provide phase equilibrium data in the melting range of the alloys. In the boundary systems, niobium and molybdenum are known to form a continuous series of solid solutions.576 The continuous solubility was also confirmed by Kornilov and Polyakova,7,8 who also observed a minimum melting point at 22 at. pct Mo and 2345°C. The phase diagram investigations of the Nb-C system by Storms and Krikorian9 and Kimura and Sasaki10 were recently supplemented by Rudy et al.11,12 The system contains a high-melting monocarbide with the B1 structure, Table I, and a subcarbide, Nb2C, which exists in at least two different states of sublattice order at low temperatures47"-&apos;3 and a disordered state above approximately 2500°C.11,12 The melting-point measurements by Rudy et al .11,14 are in close confirmation of the data by Kimura and Sasaki.10 The rather complex phase relationships in the Mo-C system were only recently Clarified.15,18 The system is characterized by three congruently melting, intermediate phases, MozC, ? MoC1-x and a Mol-,, Table I, of which only Mo2C is stable at temperatures below 1650°C. Substoichiometric MozC exists in several states of sublattice order which interconvert in homogeneous phase transitions. Hyperstoichiometric compositions cannot exist in the ordered state. Upon cooling through a critical temperature range, the

Jan 1, 1968