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By W. E. Quist, R. Taggart, D. H. Polonis, C. J. van der Wekken

An intermediate compound that has been identified as Niab is observed to form as a decomposition product from supersaturaled Ni-Nb solid solutions during aging at temperatures between approximately 300" and 500°C. On the basis of data from electron microscopy and selected-area diffraction, the structure of this compound has been determined as fct with a = b - 3a0 and c = a, wlzere a,, is the lattice parameter of the parent solid solution. The compound consists of close-packed layers with triangular ordering, where the niobiutrl atoms are separated by two nickel atoms ([long- close?-packed directions. A nine layer stacking sequence is required to describe the proposed structure. STUDIES of the Ni-Nb binary system have been limited primarily to phase diagram determinations,'-4 investigations of high-temperature equilibrium phases,5"1 and the determination of the influence of deformation on the structure of the equilibrium compound.8 The nickel-rich portion of the binary system is reported to be of the simple eutectic type in which the maximum solubility of 12.7 at. pct Nb occurs at 1282"c.' The two-phase field below the eutectic temperature is bounded by the a fcc solid solution and an orthorhombic Ni3Nb compound. No metastable phases have been reported in previous investigations. In transformation studies of certain nickel-base commercial alloys that contain niobium, two ordered metastable compounds containing niobium have been shown to precipitate from the solid solution, both of which have been identified as y' and have the composition NisNb or Ni,Nb. One compound has been reported to have the bct DOz2 type Al3Ti structure" and the other the cubic LI2 type Cu3Au structure.9,11 In the present work on Ni-Nb binary alloys a metastable y' compound has not been detected after conventional quenching and aging treatments. An anomalous behavior was noted in electrical resistivity measurements. in alloys containing between 7 to 12 at. pct Nb when aging treatments were performed below 500°C after fast quenching from 1250°C. Transmission electron microscopy has shown that this behavior is caused by the formation of a low-temperature precipitate of unreported structure type and composition. EXPERIMENTAL METHODS Several Ni-Nb alloys, containing up to 11.5 at. pct Nb. were prepared by either levitation melting and casting in copper molds or by induction melting in alumina crucibles; both techniques employed purified helium gas as a protective atmosphere. The purity of the nickel and niobium used to make the alloys was 99.98 wt pct Ni and 99.9 wt pct Nb. The composition and homogeneity of the alloys were checked by weight measurements and by electron microprobe analysis. The induction-melted alloys were homogenized for 100 hr at 1100°C. The resistivity specimens were prepared from rods swaged to 2.5 mm and the electron microscopy specimens were cut from sheet that was rolled to 0.4 mm and thinned using a modified Bollmann technique." The elevated-temperature solution treatments were carried out in a purified helium atmosphere followed by direct quenching into a 10 pct NaCl solution at 23°C. Additional protection against oxidation of the samples during solution treatment was accomplished by using tantalum foil as a "getter" in the furnace. The specimens were aged at various temperatures in salt baths controlled to +2oC. A Leeds and Northrup K5 potentiometer was used to make electrical resistivity measurements on specimens immersed in liquid nitrogen. Electron microscopy and diffraction studies were carried out with JEM-7 and Philips EM-200 microscopes operating at 100 kv. RESULTS AND DISCUSSION Ni-Nb alloys containing between 7 and 11.5 at. pct Nb that have been solution-treated in the range 1220" to 1280°C and quenched to 23°C undergo a precipitation reaction when aged in the temperature range 300" to 500°C. Precipitation was detected by selected-area electron diffraction after aging a specimen for as little as 30 sec at 350°C) whereas the reaction was well-advanced after aging for 150 hr at 475°C. Electrical resistivity measurements were used to monitor the progress of the precipitation reaction. In the present experiments the nucleation process for precipitation required a high solution temperature and a rapid quench into brine. The presence of aluminum, iron? and carbon in amounts totaling less than 1 wt pct was found by electron diffraction to completely suppress the formation of the low-temperature precipitate that has been detected in the binary alloy. Electron diffraction techniques were used to determine the structure of the precipitates that formed during the decomposition of the Ni-Nb supersaturated solid solutions. Figs. l(a) through l(d) show electron diffraction patterns oriented to the [loo], [110], [lll], and [I031 zone axes of the matrix. Areas of reciprocal space between these sections were investigated by slowly varying the orientations of the crystal under study; this procedure revealed no reflections other than those depicted in Fig. 1. The presence of super-lattice reflections at points coincident with the matrix reflections was confirmed by the examination of an almost completely transformed structure. On the basis of the accumulated diffraction data, the reciprocal lat-

Jan 1, 1970

By K. D. Sheffler, R. W. Hertzberg, R. W. Kraft

The Ni-Ni,Ti lamellar eutectic alloy responds to unidirectional solidification by alig-nment of the platelets of the two phases roilghly perpendicular to the moving solid-liquid interface. X-ray diffraction and zetallographic experimenls show that the following cryslallographic relationships exist in the directionally solidified ingots: lamellar interface w [0001] Ni3Ti 11 (111 }Ni local growth direction (1120) Ni3Ti n (110) Ni These relationships are rationalized on the basis of the crystallography of the constituent phases. In recent years researchers in several laboratories have shown that an aligned or "controlled" micro-structure can be produced in many binary eutectic alloys by forcing the alloy to solidify unidirectionally in an elongated crucible.'-3 It also has been found that a unique crystallographic relationship often develops between the two phases in such a structure.4, 5 Since a parallel arrangement of phase particles in an alloy should, and, in many cases, does, produce unique and highly anisotropic properties, the alloys so produced are of potential interest in a variety of applications. Studies of these alloys are also of interest at the academic level because they help to improve our basic understanding of the way in which metals solidify. This paper reports the metallographic and crystallographic structure of another of these alloys—that between nickel and Ni3Ti. In a separate paper, hot tensile and stress rupture properties as well as deformation and fracture behavior of the same alloy are discussed.19 EXPERIMENTAL DETAILS AND RESULTS General Characteristics of the Eutectic Reaction. Several small exploratory Ni-Ti melts were made to resolve a conflict in Hansen6 over the eutectic composition. The alloy containing 16.4 at. pct Ti, Fig. 1, appeared metallographically to be closest to the eutectic composition, although small dendritic patches of proeutectic nickel are present. The eutectic morphology is lamellar and appears quite sharp and angular, indicating a strong influence of crystallographic factors upon the microstructure. Some of the eutectic nickel platelets are continuous with the primary phase, whereas in hyper eutectic melts a "halo" of nickel generally surrounded the proeutectic inter-metallic platelets, indicating that nickel is the nucleating phase in this eutectic reaction. Extensive precipitation of Ni3Ti has occurred in the primary nickel phase in a Widmanstatten pattern which makes the dendrites difficult to differentiate from the eutectic mixture. This phenomenon may account for discrepancies among the eutectic compositions reported in the literature. The sloping solvus line bounding the nickel primary solid solution explains the presence of the intermetallic precipitate. Poole and Hume-Rothery7 report that the solubility of titanium in nickel drops from 13.8 at. pct at the eutectic temperature (1304°C) to approximately 9 at. pct at 700°C (Ni3Ti, on the other hand, melts con-gruently and exhibits essentially no solubility for nickel or titanium). Lever rule calculations based upon Poole and Hume-Rothery's phase diagram indicate an increase from 21 vol pct Ni3Ti in the eutectic mixture at the eutectic temperature to 61 pct at room temperature. Selective point count measurements of the volume fraction of eutectic Ni3Ti platelets in the eutectic microstructure yielded a value of 46 i 2 vol pct, which required extensive precipitation of Ni3Ti from the nickel solid solution onto the eutectic intermetallic platelets during cooling from the solidification temperature. Influence of Directional Solidification. Nickel bar and titanium sponge of 99.9 pct purity were induction-melted in zirconia crucibles under argon in the proportion of 13.8 wt pct (16.4 at. pct) Ti to produce 14-lb master heats which were cast into investment molds yielding 5/8-in.-diam pins approximately 10 in. long. These pins were remelted under argon in horizontal zirconia boats and unidirectionally solidified to produce 1/2-in.-sq ingots approximately 44 in. long. Details of the experimental techniques and apparatus are described elsewhere.'

Jan 1, 1970

By J. D. Livingston, H. E. Cline

Cu-Pb alloys in the vicinity of the monotectic composition have been directionally solidified under a high temperature gradient at rates up to 2 X l0-' cm per sec. Over a wide range of compositions, high growth rates yield a composite structure consisting of continuous rods of lead in a copper matrix. This range of compositions increases with increasing growth rate, in agreement with arguments based on the relative velocities of composite growth and the growth of copper dendrites or lead drops. The breakdown of the composite structure at slow growth rates is explained in terms of the relative interphase surface energies. The observed interrod spacings of the composite structure are large compared with the predictions of the Jackson-Hunt equations of eutectic growth. ThE directional solidification of many eutectic alloys produces fine composite structures of parallel lamellae or rods. There has been considerable interest not only in the fundamentals of this two-phase solidification process,'-3 but also in the interesting physical properties produced by such regular and aniso-tropic microstructures. Composite structures can be grown only over a limited range of composition, beyond which coarse primary dendrites of one phase appear. In organic eutec-tics, this composition range of composite structures has been shown to increase with increasing growth rate.7"10 These results were explained in terms of the relative velocities of composite (coupled) growth and dendritic growth. Although similar arguments should apply to metallic eutectics,11-13 suitable experimental results are lacking. In contrast to the work on eutectics, the directional solidification of monotectic alloys has received little attention. (The monotectic reaction is similar to the eutectic reaction, except that one of the resulting phases is a liquid, which subsequently solidifies in a separate reaction at a lower temperature.) Directional solidification of some monotectic alloys'4715 yields regular rodlike microstructures, whereas in other cases macroscopic separation of solid and liquid phases occurs.16 chadwick17 rationalized these results in terms of the probable relative magnitudes of the various interphase surface energies. A recent study of chill-cast Cu-Pb alloys18 revealed a fine rodlike microstructure in alloys near the monotectic composition. It was decided to investigate the directional solidification of such alloys, and to determine the range of composition and growth conditions yielding composite structures. The Cu-Pb system is convenient for such a study, both because it is simple metallurgically, with no compound formation and negligible solid solubilities, and because its basic properties are well-documented. Recent literature on the Cu-Pb system includes studies of bulk thermo-dynamic properties,'g surface energies,20"21 densi-ties,25 and diffusion constants.a6 A similar study of the directional solidification of Cu-Pb alloys was recently undertaken, independently, by Kamio and Oya." EXPERIMENTAL Alloys were prepared by melting 99.999 pct Cu and 99.999 pct Pb in a graphite crucible, stirring well, and pouring into a cold graphite mold. Rods 0.175 in. in diam were machined from the ingots. Alloy compositions studied ranged from 25 to 55 wt pct Pb. Samples were placed in graphite crucibles 5 in. long with 4 in. OD and 0.035-in. walls. They were melted under flowing argon in a vertical, two-zone. platinum -wound furnace. A voltage stabilizer was used to minimize fluctuations in input power. The narrow specimen diameter minimized convection. Directional solidification was achieved by driving the crucible downward into a +-in. hole in a water-cooled copper toroid. The toroid was located immediately below the narrow end zone of the furnace. The end zone was separately powered to maintain high local temperature. Therefore a high temperature gradient (approximately 300 deg per cm) was maintained in the specimen throughout the run. The crucible motion was screw-driven. and a wide range of drive speeds were available. The limited rate of heat removal caused a thermal lag in the specimens at high drive rates. However. temperature-time curves from thermocouples imbedded in a representative sample indicated that the average growth rate still approximately equaled the drive rate. Although the specimens were initially homogeneous, melting and re solidification redistributed the lead. producing composition variations of several percent along the specimen length. (During melting. lead melted first and ran down the sample surface. Rapid freezing tended to reproduce the resulting composition gr~dient, but slow freezing did not because a slow-moving interface tended to reject lead. as discussed later.) To determine local composition. ;-g samples were cut from regions exhibiting various microstructures and were chemically analyzed for lead content. Micrographs were taken on as-polished or lightly etched surfaces. Three-dimensional structure of the lead network was viewed with a scanning electron microscope after removal of some of the copper matrix with nitric acid. RESULTS Several different microstructures are observed, depending on composition and drive rate. Because melting and resolidification produced composition gradients, results are best presented in t&ms of final local composition, rather than initial or average composition. The ranges of local compositions and drive

Jan 1, 1970

By A. E. Dwight, Paul A. Beck

Known CsCl-type phases in transition metal alloys are shown to be formed preferentially at an electron concentration of approximately 6; a tetragonally distorted version of the CsCl structure is stable at -e/a = 6.5. The relationship of the distortion to electron concentration is verified by experimental data for the quasibinary systems Time ,Os) , Ti(Re,Ir), and Ti(Re,Pt) and the ternary system Ti-Ru-Rh. It was noted some years ago1j2 that the occurrence of CsC1-type ordered structures in equiatomic binary alloys, AB, of transition metals correlates with the location of the component elements in the periodic table. One of the components, A, must be to the left and the other component. B, must be to the right of the chromium group. It was pointed out2 that the CsCl structure does not occur in CrFe. is metastable in VFe,3 and is stable in TiFe up to high temperatures, and that this progressively increasing stability may correspond to an increase in electron transfer between the two components and to a correlated decrease of the magnetic moment associated with the iron atoms. Such a decrease was in fact later demonstrated by Nevitt. 4 The binary CsC1-type phases that are formed by a transition metal with either another transition metal or a lanthanide are listed in Table I. Thirty-three of these phases occur at an average electron concentration (average number of electrons outside of closed shells) of e/a = 6. The five phases that are formed at an electron concentration of e/a = 5.5 appear to be borderline cases. This is suggested for instance by the fact that TiRe is bcc. without any long-range order, while TiTc (at the same electron concentration) has a CsCl ordered structure. Six of the undistorted CsCl phases formed by transition metals with one another have an average electron concentration of e/a = 6. 5. Table I. With the exception of Zr4,Rhs4 and Hf48Rh52. the B component in these phases is a first long-period transition metal, i.e.. iron or cobalt. TiNi. with an electron concentration of e/n = 7, appears to undergo transformation at around room temperature to a more complex ordered struct~re.~ The five other CsCl structures known to occur at e/a = 6.5 have either scandium or a lanthanide metal as the A component; the difference in the elec-tronegativity and in the e/a ratio between the component elements is in these phases very large, A(e/a) - 7, and they may be expected to have a rather strong ionic bond component. Presumably the ionic bond stabilizes the CsCl-type structure at the equiatomic composition and it overrides the electron concentration requirement. When the B element is not from the first long period and the A component is not a rare earth metal. equiatomic (AB) alloys of transition metals with an average electron concentration of e/a = 6.5 in most cases have a crystal structure closely related to the CsC1-type, but slightly distorted tetragonally, Table I. This was first found6 for VRu, NbRu. and TaRu. It was noted6 that the undistorted CsCl structure is also stable in the three systems. but at a lower electron concentration, near e/a = 6. In these systems this requires a deviation from equiatomic stoichiometry. which is otherwise ideally suited for the CsC1-type structure. A similar situation was later found in the Ti-Rh7 and the Ti-Ir systems. 8 The Ti-Rh alloys undergo a second distortion and become orthorhombic when the rhodium content becomes somewhat larger than 50 pct. As mentioned previously. TiRe does not have the CsC1-type ordered structure. In order to investigate how close TiRe is to the borderline of stability of the

Jan 1, 1970

By K. Tangri, M. Chaturvedi

On water-quenching from within the (a + ß) phase region Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys can undergo w transfirmation. This transformation has been attributed to the enrichment of ß Zr phase, at the solution-treatment temperature, by the solute atoms. The ß—w transformation is accompanied by a hardening 01 the alloys and in the optimum condition w-bearing specimens of the ternary alloy are about 14 pct harder than those bearing only a' phase. On aging crl reverls thermally to more stable phases, which results in a considerable softening of the specirnens, as follows: w quenched -waged + ß Zr (enriched) waged - a + ß Nb The high-temperature bcc ß Zr in certain zirconium alloys can transform to a metastable w phase which has a primitive hexagonal structure. In the Zr-Nb alloy system w phase has been observed in the composition range of 7 to 15 wt pct Nb by a number of investigators.'-' In this composition range w can form by water-quenching ß Zr or by aging. known as athermal and thermal w, respectively. Hatt and Roberts' have found that thermal w has a constant c/a ratio of 0.622 +0.002 regardless of the composition and tha: the lattice parameters are of the order of a = 5.02A and c = 3.0 ?. The actual lattice parameters depend upon composition and the temperature of aging. The lattice constants of athermal w. however, depend upon alloy composition.3 The presence of w phase in zirconium alloys can cause substantial strengthening which has been attributed as probably due to the coherency between w and parent ß Zr phase. Robinson et al.' found that the presence of w can raise the strength of Zr-5.0 Nb-2.0 Sn and Zr-5.0 Mo-2.0 Sn alloys to about 170.000 psi: however. the ductility is considerably reduced. With the growing use of Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys in the nuclear reactors, and the significant effect of w transformation on the mechanical properties of zirconium alloys studies were undertaken to study the w transformation in Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys. This paper is concerned with the formation and reversion of w phase in these alloys. 1) MATERIALS AND EXPERIMENTAL TECHNIQUES The following two alloys were studied: 1) Zr-2.5 Nb-0.5 Cu (referred to as the ternary alloy): 2) Zr-2.5 Nb (referred to as the binary alloy). The chemical analysis of the alloys is given in Table I. The alloys, in the form of 3/8-in.-diam rods and 1/8-in.-thick sheets, were supplied by the Chalk River Nuclear Laboratories of the A.E.C.L. Initial fabrication for preparation of various specimens was carried out by cold rolling and swaging with intermediate anneals at 1000°C. All the heat treatments were carried out after the specimens were wrapped in zirconium foils and encapsulated in silica tubes under a vacuum of 5 x 10-% m of Hg. For optical metallography and hardness measurements disc specimens + in. diam by + in. thick were used. After the required heat treatments the specimens were mechanically and then chemically polished in a 45 pct HNO3, 45 pct H20. and 10 pct HF solution. The hardness measurements on the chemically polished specimens were carried out on a Vickers hardness tester using a 10-kg load. For each specimen at least fifteen indentations were made in order to obtain a representative hardness value. The phase identification and structural analysis were carried out using X-rays and electron diffraction techniques. Wires of 1.5 mm diam reduced to 0.12 mm diam by chemical etching were used for making Debye-Scherrer powder patterns using copper K-a radiations in 114.6-mm-diam camera. Thin films for transmission electron microscopy were prepared by electropolishing heat-treated a by 7 by 0.005-in.-thick strips using a modified Bollmann-Window technique. The 10 pct perchloric acid-90 pct methyl alcohol electropolishing bath was kepl at -5OCC and polishing was done at 5 to 10 v. The thinned specimens were washed in ethyl alcohol at -30" to -40°C and dried between filter papers. The thin films

Jan 1, 1970

By J. Szekely

A rrlethod is presented for the estitnation of oxygen pickup by teettzed molten steel streams. Of the mechanisrt~s considered for oxygen absorption, physical entrainment appeared to be the most significant. The estent of this physical entrainment may be calculated by the application of boundary layer theory. It was found that the values of oxygen pickup thus predicted were in wry good agreement with data reported in the literature, in the case of short, well-defined metal streams. This entrainment model does not provide an adequate description of oxygen pickup by long streams, exoosed to external winds and where breakup of the streatn is likely. WheN molten steel is teemed into a mold or tundish it will absorb oxygen contained in its environment. Recently interest has been expressed in the measurement and possible prediction of the extent of this oxygen absorption: this interest owes its existence to two principal factors: a) the desirability of providing a quantitative assessment of the role played by the absorbed oxygen in impairing the quality of (vacuum degassed. stainless, or continuously cast) steel; b) the need for providing guidelines for stream protection methods. The purpose of this communication is to suggest a simple method for the prediction of oxygen absorption by continuous metal streams and to compare the predicted values with the data available in the literature. It is hoped that this method may provide at least an initial framework for the interpretation of additional data that might become available. FORMULATION Consider a continuous metal stream falling into a molten metal pool as depicted in Fig. 1. This system may absorb oxygen due to the following two mechanisms: 1) physical entrainment of gases (including oxygen) into the molten pool where reactions. thus absorption, will occur; 2) mass transfer (convec:tive diffusion) of oxygen to the metal stream. 1) Physical Entrainment. The physical entrainment may be calculated by using the concept of the diaplace-ttzent tlzickness of the laminar boundary layer, which will be stated briefly in the following.&apos; When a fluid flows past a solid surface, a portion of the fluid close to the surface is retarded due to fric-tional forces. This region in which the velocity of the fluid differs from that in the main stream is called the boundary layer. An illustration of this is given in Fig. 2 where it is seen that the thickness of the boundary layer increases as we move along in the x direction. In the laminar region, when Re, < 10&apos; the thickness of the boundary layer is given by the following relationship: f =4.64 Re;1" [1] where 61 is the boundary layer thickness, x is the length along the leading edge, Us is the velocity of the free stream, (i is the viscosity of the fluid (gas in this instance).

Jan 1, 1970

By Mayumi Someno, Kazuhiro Goto, Masahiro Kawakami

The apparent rates of decarburization of liquid alloys of Fe-C, Fe-C-S, Ni-C, and Co-C systems and the rate of oxidation of solid graphite with pure carbon dioxide gas and with gas mixtures of carbon monoxide and carbon dioxide have been measured in the temperature range of 1000° to 1600°C. The cotnposition of carbon dioxide and carbon monoxide gas at the reaction surface has been measured by oxygen concentration cells with the ZrO,-CaO solid electrolyte. 1) The apparent rates of the carbon removal are essentially the same for all the cases of solid graphite, Fe-C, Fe-C-S, Ni-C, and Co-C systems under the same experimental conditions. 2) The apparent rates are independent of the carbon content in the high carbon concentration range but very much affected by the flow rate and the gas composition of the CO-CO2 reactant gas mixture. The ratio of the gas consumed by the reaction to the total quantity of the supplied gas is very large under the present experi~nental conditions. 3) There is a concentration gradient of&apos; carbon dioxide in the vicinity of the reaction surface and the content of CO, becomes extremely small at the reaction surface. 4) A large time fluc-tuation of the gas composition was observed. This jluctuation implies the presence of unstable flow in the gas phase in the vicinity of the reaction surface. THE decarburization of molten steel by an oxidizing gas or by slag may be one of the most important chemical reactions in steelmaking processes. Nevertheless, the kinetics of this heterogeneous chemical reaction do not seem to be well-solved even with the previous studies. Although the conditions for the reaction in steelmaking processes are quite different from those in the laboratory scale, some critical experiments may give information on the mechanism of the decarburization. From the previous work,&apos;-&apos; it is known that the rate of the decarburization is independent of the carbon content in liquid iron with more than about 0.2 wt pct C when the oxidizing gases are supplied to the surface of liquid Fe-C alloys on a laboratory scale. Two rate-controlling steps have been proposed for the decarburization of liquid iron with the high carbon content: one is the surface reaction control proposed by Swisher and Turkdogan;&apos; the other is that the rate is controlled by the gaseous diffusion through the gaseous stagnant layer. proposed by Baker. Warner, and Jenkins.7 and also by .Ito and Sano.2 In the present study, some experiments have been carried out for the evaluation of these rate-controlling steps in the decarburization of liquid iron with high carbon content. The apparent rate of decarburization of liquid iron has been compared with the rates of carbon removal of liquid Ni-C, Co-C, and solid graphite under the same experimental conditions. The composition of carbon dioxide and carbon monoxide gas at the reaction surface has been measured by oxygen concentration cells. I) EXPERIMENTAL PROCEDURE Fig. 1 shows the schematic diagram of the reaction chamber. Solid graphite and liquid metals were contained in an alumina or magnesia crucible of 32 mm ID and 35 mm in height. The samples were heated by high-frequency induction and the temperature was measured by the calibrated optical pyrometer. The temperature was held constant to within 10°C. The re-actant gases were supplied to the surface of the samples through the quartz tube of 8.0 mm ID. The distance from the end of the quartz tube to the surface was 20 mm. The block of high-purity graphite was cut and shaped to the inner profile of the crucible. The height of the shaped graphite was 18 mm, which corresponded to the depth of the liquid iron of 100 g. About 100 g of Fe-C alloy (4.20 to 4.40 pct C), Ni-C alloy (1.84 pct C), Co-C alloy (1.85 pct C), and Fe-C-s alloy (4.35 pct c, 0.5 or 1.0 pct S) were melted in the crucible. The reactant gases were pure CO, and gas mixtures of CO-CO,: the flow rates were controlled by capillary flowmeters with bleeders.

Jan 1, 1970

By H. J. Levinstein, H. J. Guggenheim, C. D. Capio

An optical incestigation of the phase transformations in RbFeF, has been conducted. Details of the ferromagwetic phase transition and the metamagnetic state are disczrssed. The three-dimensional magnetic domain structure existing in bulk crystals of RbFeF3 is described, as well as the effect of the magnitude and direction of the applied magnetic field on the domain struclure. The strong magnetoelastic interaction in RbFeF, is demonstrated. The results of the direct observation of the role of dislocation subboundaries and slip bands as nucleation sites and as pinning sites for magnetic domain walls are reported. THE perovskite compound RbFeF, has recently been shown to exhibit three magnetic states.&apos;&apos;&apos; Above 102°K it is paramagnetic, between 102" and 87°K it is antiferromagnetic. and below 87°K it is ferromagnetic (i.e., exhibiting remanence) with a modification of the magnetization occurring at 40°K. The changes in magnetic structure at 102o, 87°. and 40°K are accompanied by shear transformations to successively lower crystal symmetry classes.3 At 102°K a second-order phase transformation to a tetragonal crystal structure occurs. The c/o ratio increases with decreasing temperature to a value of 1.0034 at 87°K where the crystal undergoes another shear transformation to an ortho-rhombic crystal structure. The magnetic modification at 40 K is also accompanied by a shear transformation to a lower crystal symmetry. RbFeF3 is unique in that in the ferromagnetic state it is transparent in the bulk to visible light, has a low saturation magnetization, a large magnetic rotation. and has good optical properties.&apos; All these features make it an ideal material for the investigation of magnetic domain structures in bulk crystals. In addition a dislocation etch has been developed which reveals the point of emergence of dislocations with the (100) surfaces of RbFeF3,5 making it possible to determine the dislocation arrays in the material. As a result domain wall dislocation interactions can be observed in the bulk crystal. In this paper we report on 1) the crystallography of the phase transformations in RbFeF3. 2) the domain configurations as a function of magnetic field and crystal orientation. 3) the interaction of dislocations and magnetic domain walls in RbFeF,. EXPERIMENTAL PROCEDURE Since the temperature range of interest is below room temperature the dewar shown in Fig. 1 was employed. It was designed to fit on the stage of a Leitz panphot metallograph, and permitted examination of the sample at a magnification up to 150 times. The dewar consists of a double-walled glass cylinder bent into an L shape. The space between the walls is evacuated. The viewing wirldows were made of four optically flat quartz discs aligned parallel to each other and sealed to the sides of the concentric cylinders. The specimens were mounted in a holder attached to a flexible plastic shaft. Various holder designs were employed depending upon the type of observations to be made. For studies of the phase transformation between the antiferromagnetic and ferromagnetic state the holder shown in Fig. 2(d) was employed. This holder permitted accurate temperature control at 82" to 88°K by balancing the heat input from the carbon resistance heater against the heat loss to a heat sink immersed in liquid N2. Magnetic field studies were conducted by employing the holders shown in Figs. 2(a) and (b). The holder shown in Fig. 2(n) has a solenoid imbedded in it. such that the magnetic field direction is in the same direction as the incident light. The magnetic field provided by the holder in Fig. 2(b) is perpendicular to the incident light direction. The holder in Fig. 2(r) was employed when observations of the specimen were desired while an elastic bending stress was applied to the sample. The stress was applied to the sample by pulling a wire from outside of the dewar. The wire was attached to a lever pivoted on the holder. which caused the knife edge of the lever to push against the sample. The crystal growth and the sample preparation were described previously.&apos; PHASE TRANSFORMATIONS The first phase transformation in this system is from the cubic perovskite structure to n tetragonal

Jan 1, 1970

By Lyman Johnson

Single crystals qf copper in "single slip" orientatiorzs have been deformed in compression. During defortnation all of the independent deformation parameters have been measured. These parameters consist of thefive strain components and three components descrihing the lattice rotation. By a finite strain analysis these pararmeters , forrming a deformation gradient martrix, are related to the amounts of slip on each of the twelve slip systems. The results show that the amount of secondary slip is about equal to the amount of primary slip. This is an order of magnitude larger than has been believed previoutsly. ACCORDING to early theory and experiments, when a single crystal of a fcc metal is deformed in tension or compression it should deform by slip on only one slip system until the stress axis reaches a symmetrical orientation.&apos; However. the observation of a large increase in the secondary dislocation density during ..single slip" makes it clear that some slip does occur on secondary systems. Knowledge of the amount and distribution of this secondary slip is essential to a complete understanding of the mechanisms of single-crystal deformation. Ahlers and Haasen 2 and Mitchell and Thornton1 have tried to detect the amount of secondary slip in single crystals of silver and copper, respectively. Each simultaneously measured the angle A, between the tensile axis and the primary slip direction and the length 1 of a gage section of the specimen after incremental amounts of deformation in tension. The measured A, was then compared with the theoretical single slip angle hp. given by sin Ap = j sin . hO where ?o was the initial angle between the tensile axis and the primary slip direction and lo was the initial gage length. In both sets of experiments a small but systematic difference between ?e and ?p was found. This difference must be due to the occurrence of secondary slip. However, as Mitchell and Thornton1 pointed out. nothing quantitative can be said about the amount and distribution of this secondary slip from the measurements that they made. The reason that no quantitative conclusions could be made is because no unique solution for the distribution of slip on the twelve fcc slip systems can be determined from only two measured deformation parameters such as A and 1. There are, in fact, eight independent macroscopic deformation parameters that can be measured when a single crystal undergoes a homogeneous deformation. Physically these can be thought of as the five finite strain components and the three angles describing the crystal lattice rotation. All eight of these parameters were measured by Taylor4,5 for aluminum deformed in tension and compression. At that time the concern was to show that slip occurs on {111 (110) systems in fcc metals, and the mathematics were not available to determine what slip distributions were compatible with the measurements. In this paper the mathematics6,7 are developed that allow the slip distribution to be determined from these measurable macroscopic deformation parameters. The analysis is applied to the measurements of the strain and lattice rotation of copper single crystals deformed in compression. The results show that the amount of secondary slip is an order of magnitude larger than had previously been thought. CRYSTALLOGRAPHIC DESCRIPTION OF A HOMOGENEOUS DEFORMATION The deformation of a solid body can be represented by a transformation matrix F that transforms the un-deformed state into the deformed state. Consider a vector X connecting two material points in the unde-formed material and the vector x connecting the same two material points after deformation, where both vectors are referred to the same set of Cartesian axes. The final vector x is related to the initial vector X by the equation: X = FS. [2] Eq. [2] can be considered as the equation defining F, which is called the deformation gradient matrix. Its components are: If the deformation is homogeneous, the transformation is linear and the components of F are constants. Using subscript notation, if P is the unit vector in the initial direction of a material line, the components of the unit vector p in the direction of the same material line after deformation are given by:

Jan 1, 1970

By Morris E. Fine, Richard A. Kloske

The stress-strain beha1,ior in tension of Au-Ag alloy single crystals containing nominally 1,3, 5, 95, 97, and 9.9 at. pct Ag was studied uS strain role and lektlperalure down lo 4.2K. A slrain aging yield hob lc,rrs observed on aging under stress in the temperature range 30° lo 75°C. The species diffusing to the dislocation is thought to be a divacancy-solute complex wilh /he. solute then pinning the dislocation by short-Range ovdering and possibly Suzuki locking. At room tempevullcve the critical reso1ved shear stress .follo~c.s an empirical equation of the form tc = A +Bc&apos; where A and B are 90 and 420 ,g per sy mm tor gold base alloys and 60 and 510 g p~r sq )11ttz .tor si1ver base alloys. This strengthening was attributed to a long-range size ejtect. The alloy slrenglliening at i.2&apos;K is greafer than at room temperature. The additional atrloutzf was altribuled lo a uwak skovl-range i?~levacIion between the solute and dislocallon cove. The activation energy .tor deformation a/ 4.Z0h7 decreases on alloying. 11 increases irr other fee systetns. hi the 5 at. pcl Air-Ag- alloy in pmvlicular there is very extensive easy glide and lavge overshooting- of the synmetry line at 4.Z°K. There is also u sharp decrease in the vale of. strain hardenitlg near the widdle of stage II with the new rate bezng abold the sarne as tlzal In easy glide. This uws attributed to a sudden reduction of activity on the pvitrrarg sgstertr. In Au-Ag alloys the critical resolved shear stress Tc. at room temperature is essentially a parabolic function of atomic concentration:&apos; the 50 at. pct Ag-50 at. pct Au alloy is about 10 times stronger than the pure metals. This strengthening occurs in spite of the similarity in valence and atomic size. Solid-solution strengthening in fcc metals is characteristically greater at low temperatures than at room temperature. In Ag-10 at. pct Au and Au-10 at. pet Ag; r, (20.4K)/Tc (300°K) is 2.3 and 2.5. respectively.- compared to 1.2 in pure silver.3 Suzuki&apos; and Flinn5 explained the alloy strengthening at room temperature as a combination of Suzuki locking and short-range order hardening. The agreement between the computed and measured strengthening was very good: however, in Ag base-A1 alloys Hendrickson and Fine6 concluded that Suzuki locking resulted mainly in a yield drop effect. Recent reviews of the solid-solution strengthening in fcc metals by Fleischer2 and Haasen8, 9attributed the strengthening at room tenl-perature to the combined effect of atom nlisfit and change in shear modulus 011 the long-range stress field surrounding a dislocation. The parameter used by Fleischer was The shear tnoduli of Ag-0 to 6 at. pct Au alloys were measured by Pur-wins. Labusch. and Haasen.In While addition of 4 at. pct Au caused an appreciable increase in the C,, elastic constant. there was a decrease on increasing the solute to 6 at. pct Au: C,, and C :: were not measured in the 6 at. pct alloy. The greater solution hardening at low temperatures implies the presence of short-range interactions between the solute atoms and dislocations These include size and electronegativity effects. Whether the extra alloy strengthening at low temperatures is due to locking or friction hardening is still a controversy."".&apos; In the present research the stress-strain behavior of Au-Ag single crystals containing 1 to 5 at. pct Ag and 1 to 5 at. pct Au was measured vs strain rate and temperature down to 4.2&apos; K. Particular attention was paid to yield drop effects. These have not been previously reported. PROCEDURES The alloys were supplied by the Engelhard Industries. Inc.. as strips 0.050 by 0.250 by 12 in. The compositions are given in Table I. Single crystals were grown under static vacuum using a traveling molten zone technique. Crucibles shaped to the size of the strips were made of high-purity graphite prebaked for 24 hr at 1100°C in vacuum. The crucible and its charge were placed in

Jan 1, 1970

By R. M. Asimow, J. J. Svitak

Tile crilical resoll,r,d shear stress, CRSS, for slip of slow1y cooled Ag-All single crystals was measured at 201K. Tlze importance of short-range order and Suzuki sogregalion was inrestigated by deter~rlining llze ciic.cl of quenching on the CRSS. Thr obserz.ed rfrtenching effects were generally small and indicate that short-range ordcr and Slczuki srgregation do no1 account for n significant portion of the solid-solution strengthening in this system. However, all alloys slzozcwi yield drops ulhich rilere heat- lreat nrenl-sensiline THERE have been a number of mechanisms proposed to explain alloy strengthening in the Ag-Au system. The effectiveness of some of the proposed mechanisms, especially those that depend on Suzuki segregation1" or on short-range order,2-6 SRO? should be strongly dependent on heat treatment. Since internal friction investigations in this system indicate that there is stress-induced order,&apos; we must also consider the possibility of local atomic rearrangement or short-range ordering in the stress field of a dislocation. There is not complete agreement whether these mechanisms should be most effective in raising the flow stress, in causing yield drops. or some combination of these two effects. This work was undertaken to determine the effects in the Ag-Au system of mechanisms that depend on segregation or order by using known but drastically different cooling rates following a high-temperature treatment. By comparing the critical resolved shear stress, CRSS, for slip measured after a very slow cool and again after a fast cool (quench) of the same crystal, it is possible to determine the effect of cooling rate on CRSS* and on yield drops. The single-crystal CRSS was determined at several compositions because there has been no data published across the phase field since the publication of room-temperature flow stress by Sachs and Weerts.&apos; EXPERIMENTAL PROCEDURE Starting with materials of 99.999 pct purity, single-crystal specimens of random orientation 1.25 mm square by 4 cm long were prepared with 0, 10, 14, 25, 30. 50, 70. 75. 80, 90. and 100 at. pct Au. Crystals were grown in an inert atmosphere using a modified Bridgman technique with a growth rate of 1.5 cm per hr in a gradient of 3°C per cm. To obtain a maximum practical degree of SRO, the crystals were cooled at a rate of 275&apos;C per hr to about 310°C and then cooled at 6°C per hr to approximately 200°C. Specimens were tested in tension immersed in an alcohol bath at -72°C (201°K) at a strain rate > = 0.003 mm"1 to determine the CRSS (7). The activation volume of several crystals was determined from the flow stress increment (72 - T1) for a tenfold increase in strain rate. The total shear strain varied from 0.05 to 2.3 pct but was generally less than 0.4 pct. After testing, each crystal was then held at 500°C for 10 min, quenched into brine at -35°C and. after storage at -71°C for 1 day, retested in tension using a procedure identical to that of the first test. The quenching apparatus, see Fig. 1, was designed to enable accurate temperature control and rapid quenching while preventing specimen contamination and accidental deformation. Test data from a specimen with a fine thermocouple spot-welded to it indicated a cooling rate of 104C per sec. The relatively low heat-treatment temperature was chosen to minimize quenching defects caused by a large vacancy

Jan 1, 1970

By B. Chalmers, L. Bäckerud

The growth temperature of dendrite tips has been measured in a binary alloy system, Al-Cu, as a function of rate of formation of solid phase. Metallographic examination has rerealed the dendritic structure dereloped during different stages of the freezing process and factors influencing the undercooling, such as solute buildup, surface tension, and kinetic effects, have been estimated. Practical implications of the results are discussed. THE development of a dendritic growth pattern during the solidification of metals and alloys has an influence at least equal to that of grain size on the properties of the casting or ingot being produced. and consequently much interest has been devoted to the study of this phenomenon. A review of the knowledge gathered up to 1964 has been given by Chalmers.&apos; Among recent publications in the area the works of Jackson et al 2 and Flemings and his coworkers3,4 are of special interest. This work is an experimental study of the factors which influence the development of a dendritic pattern during the solidification of a typical binary alloy system, A1-Cu, when there is no heat source in the liquid. This means that the temperature and the temperature gradient are controlled by the heat extraction from and latent heat evolved by the growing crystals. EXPERIMENTAL PROCEDURE Samples were prepared from a master alloy made up of aluminum, 99.995 pct, and copper. 99.99 pct, of the concentrations given below and of sufficient size to fill cylindrical graphite molds of 35 mm ID and 35 mm height. Bottom. wall. and lid of the graphite molds were all of 6 mm thickness, Fig. 1. Thin ceramic thermocouple shields could be introduced at different positions in the sample through holes in the lid. The assembly was heated in a small vertical resistance furnace to a suitable temperature. 700° to 720°C. which is well above the melting point of the sample. It was thereafter taken out of the furnace and chilled at selected rates, and the temperature changes in the sample were recorded. For high sensitivity in the temperature recordings, a reference thermocouple was held in a solidifying pure aluminum melt. which was very well insulated and therefore was at a constant temperature. 660°C, for nearly half an hour. The thermocouples were connected so that the difference voltage was fed into the recorder of which the most sensitive range was 1 mV. Temperature variations could be read to within 0.l°C. The thermocouples used were of the Chromel-Alumel type (28- or 38-gage wires) and the recorder a two-pen Moseley strip chart recorder. The different rates of chilling were obtained in the following way. a) The graphite crucible was wrapped in one or two layers of an insulating. feltlike material (Fiberfrax paper from the Carborundum Co.). The drop in temperature per unit time after complete solidification could thereby be varied from 8" to 18°C per min. This corresponds to a rate of heat extraction in the given samples from 0.08 to 0.17 cal cm-&apos; sec-I. 6) The bare graphite mold was placed free on a piece of firebrick. The measured temperature drop was around 50°C per min (-0.5 cal cm-" &apos; sec- I). c) By blowing compressed air through a tube with many holes surrounding the mold in a ring shape. the rate of heat transfer could be raised to 150&apos; to 200°C per min (-1.7 cal cm-3 sec- I). d) Cooling rates around 1000 C per mill (-9.4 cal cm-3 sec-l) were obtained by substituting the air in the device used in c by water. The range of cooling rates used in these experiments correspond to technical casting processes from slow solidification as obtained in sand molds to very rapid freezing which occurs in some continuous casting and welding processes. EXPERIMENTAL RESULTS The A1-Cu system was chosen for the experiments: the relevant part of the phase diagram is reproduced in Fig. 2. Observation of the Center Temperature. When a thermocouple is inserted in the center of the melt. the type of cooling curves obtained at slow solidification rates differs little from numerous published curves from similar systems. and the solidification pattern

Jan 1, 1970

By Morris Cohen, Robert C. Ruhl

The phases in Fe-C alloys over a wide composition range have been studied after splal quenching from the liquid state. Binary alloys containing 0 to 5.1 wt pel C as /cell as a large number of ternary Fe-C-Si alloys with 2.5 to 5.0 wt pct C and 0.3 to 5.1 wt pct Si were attlong those sludied. Olher Fe-C-X alloys, zcilh X being Co, Cr, ,Wn, Ni, and Ru , were also inrestigated after splat quenching. At high carbon contents, a new hcp phase (designaled 6 phase, but different from e carbide) is retained upon splat quenching. The .fraction of this phase varies up lo 97 pcl for a Fe-4.K-1.9% alloy. The composition of the E phase ranges from about 3.8 10 4.8 wt pet C, and the corresponding laltice parameters increase linearly ulith carbon content, while the c/a ratio remains essentially conslc~nl. The E phrtse appears to he a solulion of carbon in E iron, the latler being nornially found only at high pressures. It is deduced that the unit cell of the E phase corresponds to the formula Fe12C3, and llzal il is relaled to tlie ordered slruclures 0.f 6 iron carbide and c iron nilride. The E phase is compared and contrasted to the olher known carbides and nitrides of iron and nickel. An exlrapolaTion of the atomic volume 1,s carhon conlent of /he E pllase lzcts giz.en a neu7 estitnale jor /he alomic volume of E iron, 11.30 cu A, a1 atmospheric pressure and temperature. Other alomic volume relalionships lead to /he co~zpositioti Fe2.iC tor E iron carbide, /he unit cell fortr~ula being -Fe2rClo The E phase undergoes a lulo-slage decomposition upon healing, .forrning firsl rnarlensile plus E carhide, a/ler 1 hr at 140" lo 200°C, slid then ferr ite Plus cementile, after 1 lir a1 330" to 460°C. A1 carbo,l contents between 1.5 and 3.0 LC/ pcl, (he predo.wirzar/t plzase alley quenching is fcc austenite. The retained carbon content of this phase increases with itlcreasing silicon in certain concentration ranges, reaching a maximum of 2.37 wt pct C itz a Fe-2.6C-4.OSi alloy. This is the highest carbon conten1 reLaitled in austenite to date. These high-carbon aus-tenites can be partially tm?zsforttled lo tnartensile hy severe deformation in the temperature range of — 190 to -50°C. TECHNIQUES for splat quenching from the liquid state have been utilized in numerous recent investigations to produce metastable phases in a variety of alloy systems. Among the several ways of splat quenching, the shock-tube method appears to yield the highest cooling rates1-3, 7, 8 and was adopted here. Estimated cooling rates attained in the present experiments ranged from 10&apos; to 10 80Cper sec.&apos; As a part of a research program on interstitial al- loy phases, the Fe-C system was selected for splat-quenching studies. It was hoped that splat quenching would allow high metastable supersaturations of carbon to be retained in solution. Also of considerable interest were the conditions governing the occurrence of the various intermediate phases upon solidification. The alloys investigated included both binary Fe-C compositions as well as six ternary Fe-C-X alloy systems. The known phases in the Fe-C system are summarized in Table I.* Only the ferrite and graphite are ent investigation.14 and is described in detail herein Table II summarizes corresponding data on Fe-N phases, which are also of interest here because of their similarity to the Fe-C phases. EXPERIMENTAL PROCEDURES Alloys were prepared by melting the elements. 99.9 pct purity, in an inert gas nonconsumable electrode arc furnace. The buttons. weighing about 5 g. were remelted twice. were then fragmented. and their interior surfaces were examined for uniformity: if any doubt existed. they were remelted again. Chemical analyses were performed on all the alloys. the accuracy being about k0.05 wt pct. At high temperatures, carbon-containing alloys react with alumina crucibles as follows: If the atmosphere in the splat-quenching furnace does not contain sufficient carbon monoxide the alloy can be depleted of carbon and contaminated with aluminum. Calculations and experimental observations showed that 50 to 100 torr CO partial pressure effectively blocked the above reaction in all the alloys investigated. The splat-quenching equipment in Figs. 1 and 2 provides for evacuation and back-filling with CO-Ar mixtures. The furnace is capable of operation up to 1650°C. and the gettering action of the graphite heating element reduces the oxygen partial pressure in the furnace atmosphere to below 10-5 torr, thus preventing oxidation of the specimens.

Jan 1, 1970

By Morris Cohen, Robert C. Ruhl

Fe-Ni-B alloys were inresligated by X-ray diffraclion after splat quenching. Although this rapid cooling did not produce a measurable supersaturation of dissol1ed boron in either binary Fe-B or Ni-B alloys, a marked increase in dissolved boron was achieved by the splal quenching of lernary Fe-Ni-B alloys. Boron additions up to 2.0 wl pcl to Fe-13Ni and Fe-24Ni alloys cause an increase in retained austenite from 0 to 59 pet and to 92 pet, respectively. The martensitic phase in the Fe-13Ni-9 at. pct B alloy is .found to conlain about 0.6 at. pet interstitial boron, as evidenced by the tetragonality of this phase, and about 3.6 at. pct substilutional boron, us deduced from the lattice paratrieler decrease, the balance of the boron being precipitated as (Fe,Ni)aR. The austenitic Fe-Ni-B plznses also contain both interstitial and substitutional boron. Iron-Boron System. The phase diagram and related information about the solid solutions and intermediate phases in the Fe-B system are given by ansen,&apos; Elliott.2 and Pearson.1 The equilibrium solid solubility of boron in both ferrite and austenite is quite small: the maximum solid solubility of 0.02 wt pct B (0.11 at. pct B) occurs in iron at the eutectic temperature of 1150°C. while the maximum solubility in a iron is less than 0.01 wt pct.l There is disagreement in the literature as to whether these solid solutions are substitutional or interstitial. Based on the measured activation energy of boron diffusion in iron, it appears likely that boron moves interstitially in austenite. 4 However, this observation does not rule out the possibility that some of the dissolved boron may also occupy substitutional sites in this phase. Investigations of internal friction of boron in iron have suggested that boron is interstitial there.&apos; Shevelev, 5 on the other hand, reports a decrease in the ferrite lattice parameter with dissolved boron content, thus indicating a substitutional solution. Unfortunately. Shevelev&apos;s methods and analysis are unclear. Goldhoff and spretnak6 have measured the lattice parameters of pure iron and an Fe-0.005 wt pct B (0.027 at. pct) alloy at both room temperature (ferrite) and from 920° to 1200°C (austenite). They reporte: a lattice contraction in the ferrite of 0.0003 ± 0.0002 due to boron, thus indicating a predominately substitutional solution. Appreciable (0.0010 to 0.0060) lattice contractions were found in boron-containing austenites, but these effects were much larger than could be accounted for by the substitutional solution of all the boron present. Nickel-Boron System. Elliott 2 gives a summary of thephase diagram and recent work on this system. The solubility of boron in nickel reaches a maximum of 0.03 wt pct at the eutectic temperature of 1080°C. Iron-Nickel System. ansen&apos; presents the equilibrium diagram for this system and also discusses the metastable austenitic and martensitic phases along with the a=? transformations. The lattice parameters of Fe-Ni ferrites have been recently redetermined by Abrahamson and Lopata.7 Their data, Fig. 1, show an unusual change in slope at about 1.7 at. pct Ni. Very few measurements of the lattice parameters of Fe-Ni martensites are available in the literature. The values of Roberts and Owen8 are considerably higher than those measured in the course of the present work: they find that a, is constant at 2.874? from 8 to 30 at. pct Ni. Both the slope change in the ferrite parameters and the occurrence of a maximum plateau in the marten-sitic lattice parameters vs composition denote the

Jan 1, 1970

Jan 1, 1970

By G. J. W. Kor, F. D. Richardson

The sulfide capacities a1 1500°C of slags in the systern CaO + CaF, +Al2O3 have been measured by bringing samples into equilibrium with gas mixtures made from CO + CO2 + SO, + N2. The capacities are negligzble both JOY CaF, and CaF, + Al2O3 melts but when lime is present even at 20 wt pct capacilies as high as 30 x lom3 can be obtained. The results have yielded information concerning the liquidus at 1500°C and fluoride-oxide melts in general. The slags used in electroslag refining contain substantial amounts of calcium fluoride. Robinson and Grainger&apos; report that good results can be obtained by using slags of the approximate composition 70 wt pct CaF2 + 30 pct Al2O3. It has also been reported2 that a mixture of 50 pct CaF2 + 30 pct A12O3 + 20 pct CaO makes a good slag for general use. In view of the fact that very low sulfur contents are claimed for ingots produced by the electroslag process. studies have been made of the sulfide capacities [c, = wt pct S (PO2/Ps2)1/2 at 1500°C. of slags consisting of pure calcium fluoride, and of its mixtures with lime and/or alumina. These studies have also provided information on the phase diagrams. The slags were brought into equilibrium with gas mixtures prepared from NZ + CO + CO2 + SO2. Because CaFz is volatile and also easily hydrolyzed. measurements were made of its loss in weight in various flowing gas mixtures. EXPERIMENTAL Procedure. The apparatus and procedure were the same as described previously. Both platinum and iridium crucibles were used (height and inner diameter about 1 cm), but those made from iridium were stronger and able to withstand a greater number of experiments. In each experiment two crucibles were set in an alumina support and 1 g of slag mixture was divided between them. The materials from each crucible were analyzed separately after being brought to equilibrium with the gas, and then quenched in the water-cooled end of the reaction tube. Materials. Alumina and lime were prepared as de-scribed previously.4 Calcium fluoride was obtained by igniting Analar calcium fluoride in a platinum boat at about 800°C in a stream of pure dry nitrogen: the product was cooled in a desiccator. The slags were produced from mixtures of the powdered constituents, these being melted together in the course of each equilibrium experiment. The sulfur analyses and the X-ray diffraction meas- urements on the solidified slags were made as described previously.4 From the latter it appears that the CaF2 contained less than 0.5 mol pct CaO. RESULTS Values of (ps2/po2) for the Gas Mixtures. The re-sults have been represented in terms of (ps, /Po,)"? for the various gas mixtures used. The va1;es of this function at 1500°C are listed in Table I: they were calculated from the data cited previously for calculations at 1350°C 4 Weight Losses with CaF, and CaF, + CaO. The re-sults of the weight loss measurements* at 1500°C for

Jan 1, 1970

By Shiro Ban-ya, John Chipman

Using the same experimental method previously described, the activity of sulfur in a number of quaternary and more complex liquid iron alloys at 1550oC is determined. A "cross-product term zj.zk is determined for each of the quaternary alloys to take account of the mutual effects of the combinalion on the indi-vidual contributions to In 5 The interactiotl coefjicienls delermined for ternary solutions and cross-pvoducl coefficierlts from the quaternary systems are used to estimate the activity coefficient of sulfur in more complex alloys. OUR recent papers 1,2 have described new determinations of the activity of sulfur in Fe-S binary and Fe-S-j ternary solutions, but the practical applications of these data require the estimation of activity of sulfur in more complex systems. Previously, Sherman and chipman3 devised an approximate method for calculating the activity of sulfur in complex systems by summing the contributions to the logarithm of the activity coefficient of each element in the alloy. This procedure is more successful, as one of us has pointed out,4 when the concentration variable is taken as zs = ns/(nFe — nS t CLJ~ nj) rather than the conventional atom fraction x~ = nS/(nF, + nS + Enj). In the former definition, which is based on analogy to interstitial solutions, the value assigned to uj is normally +1 for metallic and -1 for nonmetallic elements." In addition, however, the effects of mutual interaction between alloying elements must be considered for the accurate estimation of the activity of sulfur in more complex systems. The purpose of this investigation has been to devise a more reliable method for estimating the activity coefficient of sulfur in nmIticomponent iron alloys. The measurements were carried out on some quaternary systems Fe-S-j-k for combinations of silicon, chromium, nickel, cobalt, molybdenum, and tungsten. Manganese was not included since the experimental method has been found unsatisfactory for volatile, sulfide-forming elements. The experimental method was the same in every detail as described previously. 1,2 The concentrations of sulfur, silicon, chromium, and copper in the quenched metals were determined by chemical analysis but those of nickel, cobalt, molybdenum, and tungsten were calculated from charged amounts. Iron sulfide was included in the charges in amounts estimated to furnish approximately the final sulfur content. CALCULATION W2 have shown2 that when zs and zj are the concentration variables in a ternary Fe-S-j solution the logarithm of the activity coefficient S = a~/zg is a linear function of the variables as and aj defined as follows:

Jan 1, 1970

By M. Garfinkle, W. D. Klopp, W. R. Witzke

The tensile properties of binary W-Re alloys containing up to 33 at. pct Re were determined at temperatures from 78" to 3630°F. Elongations as high as 260 pct were observed in electron-beam-melted tungsten containing 23 at. pct Re when tensile-tested at 3630°F after a 1-hr anneal at 3090°F. All alloys tested under these conditions with rhenium contents between 20 and 28 at. pct exhibited elongations of at least 200 pct. These alloys also showed enhanced grain growth rates. The values of strain-rate sensitivity ranged between 0.2 and 0.3 for the alloys exhibiting high elongations. Houlezler, for swaged alloys tested just above the re-crystallization temperature, strain-rate sensitivities as high as 0.8 were observed. The high-temperature strengths of the high-rhenium-content alloys were also less than those of alloys with intermediate rhenium contents. ALLOYING with rhenium has two distinct effects on the mechanical properties of tungsten. One effect occurs in dilute solid solutions and is most clearly seen in the lowering of the hardness and of the ductile-brittle transition temperature at about 5 at. pct Re.1-3 The second effect occurs in the concentrated solid solution at about 25 at. pct Re, and is characterized by twinning and a larger decrease in the ductile-brittle transition temperature. An aspect of the high-rhenium-content alloy which has received little attention is the higher-than-usual tensile ductility at 3500°F.3 Anomalous high ductility has been observed in several other alloy systems. usually associated with a phase transformation or a solid-solubility limit.4&apos;5 The phenomenon is now commonly termed superplasticity. High-strain-rate sensitivities6 and grain boundary sliding7 have been reported for materials that undergo superplastic deformation. Also, it has been found that superplastic materials usually have a very fine grain structure (less than 10 pm).8 and it has been proposedg that such a structure alone is sufficient for superplasticity. The purpose of the present study was to further characterize the high-temperature tensile properties of W-Re alloys to determine the conditions under which they might be superplastic. Tensile and selected strain-rate sensitivity tests were conducted on alloys containing 0 to 33.3 at. pct Re at temperatures from 78" to 3630°F. Grain-growth rates were also determined on most alloys at 3630°F. Fig. 1 indicates the compositions evaluated and their locations relative to the solvus line in the tungsten-rich end of the W-Re phase diagram. Twelve binary W-Re alloys containing up to 33.3 at pct Re were double electron-beam-melted and subsequently fabricated to 0.3-in. rod by extrusion and swaging. The starting materials were commercial powders of tungsten and rhenium with impurity levels as listed in Table I. The various blended W-Re compositions were hydrostatically pressed at 70 ksi. The pressed bars were drip-melted into 1.5-in.-diam ingots and then remelted into 2-in.-diam ingots at a chamber pressure of 10~5 torr. The ingots were canned in molybdenum and extruded to round bar at a reduction ratio of 6:l. The alloy extrusions were swaged to about an 80 pct reduction at 2100" to 2900°F. Buttonhead tensile specimens with a reduced section diameter of 0.16 in. and a gage length of 1.0 in. were centerless-ground from the swaged bars

Jan 1, 1970

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

The least work analysis of constrained deformation, first proposed by Taylor and recently modified by Chin, Hosford, and Mendorf to include twinning, has been applied to the problem of axisynmetric flow. Calculations were made, for the case of (111}(110) slip and {111}(112) twinning applicable to fcc deformation. The solutions were obtained by linear programming techniques and with the aid of computer methods. The results show that in axisymnzetric tension the onset of twinning tends to decrease the plastic strength anisotropy as compared with the case of defortnation by slip alone. In compression, however, the strength anisotropy is increased as a result of twinning. The [100] and 11111 orientations and their vicinity are unfavorable for twinning in tension and compression, respecticely. With slight modifications, the present results may be applied to bcc metals deform ing by {110}(111j slip and (112} (111) twinning. Implications of present results to wire texture development of low stacking fault energy fcc materials arc discussed. PREVIOUSLY we&apos; have applied linear programming techniques to obtain solutions to Taylor&apos;s least work analysis2 of constrained deformation of single crystals undergoing axisymmetric flow, such as occurs in an axially loaded polycrystalline aggregate. These solutions yielded values of the parameter M = o/.r = Y/w which relates the external stress-strain (u-c) curve to the shear stress-shear strain (7-y) curve for the crystal. The value of M is dependent on crystal orientation and hence enters significantly in texture strengthening considerations.3&apos;4 Recently the Taylor hypothesis was modified by Chin, Hosford, and Mendorf&apos; to include mechanical twinning. The latter has now been applied to the problem of axisymmetric flow and the results are reported in this paper. Calculations have been made for {111 )( 110) slip and {111 )(112) twinning applicable to fcc crystals. With minor changes in the indices of the slip and twin systems, the present results are also applicable to bcc crystals which slip on (110) (111) and twin on {112)(111) systems.

Jan 1, 1970

By A. S. Yue, A. E. Vidoz, F. W. Crossman

Tensile specimens were prepared from a single grain of an epitaxially grown Al-CuAl2 eutectic ingot. The eutectic lanzellae were oriented parallel and perpendicular to the tensile axis of the specimens. Since the composite was of the eutectic composition, the aluminum-rich matrix could dissolve up lo 5. 7 wt pct Cu in solid solution and, therefore, was amenable to strengthening by precipitation hardening. The tensile properties of the eutectic single crystals were determined at room temperature as functions of interlamel-lar spacing, platelet orientation, and thermornechanical trealment. The obserced variations in composite stress and modulus with respect to the level of&apos; composite strain are discussed in terms of premature fracture of CuAlz platelets, a distribution function for the strength of the lamellae, and unequal strains due to localized fracture of&apos; platelets. The discontinuous fiber composite model of Kelly and Tyson is modgied to account for a changing distribution of fiber lengths during composite loading. The tensile properties at elevated temperatures were determined for the direc-tionally solidified eutectic oriented with platelets parallel to the tensile axis. The observed properties are attributed to the onset of plasticity of the CuAL2 phase above 150°C. DURING the investigation of whisker- and fiber-reinforced metallic matrix composites in recent years, two major problem areas have developed: 1) The fabrication of the composite involves tedious handling techniques in order to obtain a unidirection-ally aligned and uniformly spaced set of whiskers in the metal matrix. 2) Due to weak interfacial bond strengths or because of the formation of additional embrittling phases at the metal-fiber interface during long-time exposure or fabrication at elevated temperatures, many composite systems have exhibited considerably lower strengths than those predicted by a law of mixtures analysis.&apos; These problems have been bypassed by the technique of growing whiskers and plates of high-strength materials in a ductile metal matrix by controlled unidirectional eutectic solidification.2 The tensile properties of directionally solidified A1-CuA12 eutectic are presented here. This alloy consists of a ductile aluminum matrix, containing up to 5.7 wt pct Cu in solid solution, which is amenable to precipitation hardening by heat treatment and a reinforcing high modulus CuAlz intermetallic phase. The two phases are present in the form of alternating platelets or lamellae. The microstructural stability of this unidirectionally solidified alloy at elevated temperatures has been studied extensively.3.4 and preliminary tensile and bend tests have been reported. 5-7In the present investigation the tensile properties of the A1-CuA1, eutectic have been studied as a function of several ther-momechanical variables: solidification rate. heat treatment. rolling at elevated temperatures. and lamellar orientation. It was felt that the uniformity of structure and excellent interfacial bonding would give tensile properties concomitant with the metal matrix composite theory of strengthening proposed by Kelly and coson. 8-9 The tensile properties that were obtained point to a wide distribution of strengths for the CuAlZ platelets, which leads to large deviations from the predicted mechanical behavior for this composite. EXPERIMENTAL PROCEDURE Epitaxial Growth of Eutectic Alloy. The A1-CuA1, eutectic alloy was prepared by an epitaxial growth process. Sections of a master alloy ingot (total impurity content <0.008 pct) were placed in an alundum boat, melted. and directionally solidified to obtain a multigrained plate 12 by 2 by 4 in. This plate was tapered at one end to mate with a seed crystal 1; in. long and 4 by $ in. square. Then the seed-plate combination was placed in an alundum boat which sat in a quartz tube passing through the center of a horizontal resistance wound tube furnace. A dried argon atmosphere was maintained. The temperature gradient in the furnace was such that the liquid-solid interface of the eutectic alloy was located near the end of the furnace and could be observed through the quartz tube. Single-crystal plates were formed by melting the material back to the midpoint of the seed crystal of the desired platelet orientation and then epitaxially growing the plate from the seed by withdrawing the alundum boat from the furnace at a constant rate. This technique was used to produce aluminum and CuA12 lamellae parallel and perpendicular to the transverse direction on the plate. Metallographic examination showed that both phases were continuous across the original liquid-solid interface. It was also possible to grow a plate from two seeds placed side by side: and, although the lamellae of one seed were oriented at 90 deg to those of the second seed, the interface between the two grains remained parallel to the growth direction along the entire ingot length. Maintenance of a straight intergranular boundary during the solidification process was possible as long as both seeds were oriented with their original growth direction parallel to the solidification direction of the plate. Eutectic plates were directionally solidified at rates of 0.2, 1.0. and 4.7 cm per hr and sectioned transversely to the solidification direction to determine the apparent inter lamellar spacing of the lamellae. Metallographic examination was also employed

Jan 1, 1970