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By J. Campbell

Various models are discussed for the evaluation of the negative pressures which may occur in solidifying materials which exhibit various deformation modes: elastic-plastic, Bingham, viscous, or creep flow. The inadequacy of the previously proposed elastic-plastic solution for solidifying metals is revealed by comparison with the more reliable creep results which are given graphically for aluminum, copper, nickel, and iron. The maximum tensions experienced in the liquid phase of solidifying spheres ranging in size from large castings to submicron powders are in the range from —10' to —105 atm for these metals. THERE has been much recent interest in the negative pressures associated with the volume change on solidification and in the possibility of the occurrence of cavitation. Considering the freezing of a highly supercooled liquid, an attempt to evaluate the stresses in the liquid ahead of the rapidly moving solidification front has been made by Horvay1 on a microscale and by Glicksman2 on a macroscale. In a casting of a wide freezing range alloy, the pressure differential due to viscous flow of residual liquid through the pasty zone has been discussed by Piwonka and Flemings,3 In a previous publication4 the author has attempted to estimate the negative pressure occurring in the residual liquid of a spherical casting, employing an elastic-plastic model to describe the collapse of the solidified shell under the internal tension. An earlier model assuming a rigid shell was shown to be inaccurate by many orders of magnitude. The elastic-plastic model is critically reviewed here, and other models are developed which are thought to be more closely related to metals and other materials near their melting points. The spherical geometry (Fig. 1) is chosen because the highest shrinkage pressures would be developed, although the analyses are readily adaptable to cylindrical geometry. A parallel sided casting experiences little internal tension because of the relatively easy dishing inward of the sides. (This commonly observed phenomenon has previously been attributed solely to atmospheric pressure.) Furthermore, small regions of confined liquid in a large solidified volume of a casting approximate reasonably well to spherical geometry. ELASTIC-PLASTIC MODEL The author has shown4 that as solidification proceeds the internal hydrostatic tension builds up until the elastic limit of the shell is exceeded. At this point the internal pressure is closely -2Y/3. Subsequently a plastic zone spreads from the inner surface toward the outer surface of the shell. When the whole casting is deforming plastically a rather more generalized analysis taking account of the externally applied pressure PA + 2y/b gives the internal pressure as: P = Pa + 2y/a + 2ys/b - 2 Y In(b/a) [1] The 2y/a and 2ys/b terms result from the tendency of the liquid-solid and solid-vapor interfaces to shrink, reducing their energy, and thereby helping to collapse the solid phase and compress the liquid phase. The 2y/b term would be important only for powders. The last term arises because of the plastic restraint of the solid, resisting collapse and so effectively expanding the residual liquid. From Eq. [I] it is easily shown that there is a minimum in the pressure at the radius amia= y/Y [2] which is of the order of 103K for the metals aluminum, copper, and iron, and corresponds to the minimum pressure Pmin = 2 Y[l-ln(bY/y [3] The results of a fully worked out elastic-plastic solution are given in a previous reporL4 The main criticism which may be leveled at this analysis when applied to metals at their melting points is the strong dependence of the yield stress on the strain rate. The strain rate varies with both solidification conditions (e.g., whether chill-cast or slowly cooled) and during solidification, as is indicated in the following section. Thus an appropriate choice of Y is very arbitrary. Before proceeding to a discussion of models which are strain-rate-dependent, it is necessary to evaluate the strain rate as a function of the rate of solidification. SOLIDIFICATION RATE Various empirical relations have been deduced5 for the rate of thickening of the solid shell by pour- out tests on partially solidified spheres. These, however, are unsatisfactory for our purposes since they become very inaccurate when the liquid core is very small. A theoretical approach is therefore necessary, and some solutions are set out below. Making the assumptions of constant surface temperature of the casting during freezing, no superheat, and a material freezing at a single temperature, Adams8 deduces the approximate solution: which becomes when b » a: Employing a semiempirical approach vallet6 finds:

Jan 1, 1969

By E. A. Steigerwald, G. L. Hanna

The influence of work-hardening exponent on the variation of fracture toughness with material thickness was studied for high-strength steel, aluminum, and titanium alloys. The results indicate that, when materials are compared at similar fracture toughness to yield strength ratios, the material with the lower work-hardening exponent undergoes the transition from flat to slant fracture at a larger thickness than material with a high work-hardening exponent. In the thickness range where complete slant fracture is obtained the reverse is true and a lower work-hardening exponent results in a lower fracture toughness. The influence of work-hardening exponent on fracture toughness is, therefore, dependent on the particular fracture mode. In the transition region a low work-hardening exponent is beneficial for fracture toughness while in the 100 pct slant region it is detrimental. THROUGH the use of fracture mechanics analyses, the influence of geometric variables on the crack propagation resistance of structures can be interpreted with reasonable consistency. However, in order to gain a more complete understanding of the fracture process, the influence of material parameters on crack propagation must be defined and coupled to the macroscopic fracture mechanics approach. The work-hardening exponent, which characterizes specific material behavior, may serve as an effective parameter to allow some degree of coupling to be accomplished. In the extension of a crack in a specimen of suitable dimensions the propagation process occurs in a stable manner when the crack extension force is balanced by the resistance to crack extension, which exists in the material at the crack tip. As the applied stress, and therefore the crack extension force, on the specimen increases, the resistance also increases primarily because the effective plastic zone at the crack tip, which is the main energy absorption process, becomes larger. Since the work-hardening rate of a material influences the stress-strain relationship, it will also influence the energy absorption process in the plastic enclave at the crack tip and hence should have an effect on crack propagation. A number of studies have been made correlating the strain-hardening exponent or the strain to tensile instability with the ability of a material to resist fracture. Gensamer1 concluded that a low-strain-hardening exponent would result in a steep strain gradient at the base of a notch. He reasoned that a large work-hardening coefficient would result in high-energy ab- sorption due to the increased area under the stress-strain curve. Larson and Nunes2 experimentally observed in Charpy tests on steels heat-treated to below 200,000 psi yield strength that the energy to failure in the fibrous mode, i.e., above the brittle-to-ductile transition temperature, was logarithmically related to the strain-hardening exponent. In order to avoid the complicating effects present in studying materials which undergo a brittle-to-ductile transition, Ripling evaluated the notch sensitivity of a variety of fcc metals with varying work-hardening exponents.3 The results indicated that the relative notch sensitivity, as determined from tests on specimens with a sharp notch, decreased with increasing values of strain hardening. Although the energy required for ductile or fibrous fracture increases with increasing work hardening, high-strength steels often exhibit improved crack propagation resistance when heat-treated to obtain low values of strain hardening.4,5 An analysis of whether low strain hardening is beneficial or detrimental to crack propagation resistance must depend on the particular fracture criterion involved. At temperatures where the material is relatively ductile and the development of a critical strain is required for fracture, high strain hardening increases the energy required to produce failure. In the transition region and below, however, a critical stress law appears to be valid6 and a low rate of work hardening may produce superior resistance to semibrittle crack propagation. The experimental program is aimed at studying these possibilities and determining the specific influence of strain hardening on the crack propagation resistance of several high-strength materials. MATERIALS AND PROCEDURE The alloys, chosen as representative of various classes of high-strength materials, are summarized in Table I. The heat treatments evaluated along with the smooth tensile properties are presented in Table 11. Pin-loaded sheet tensile specimens were employed to determine the smooth tensile properties. A strain gage extensometer (measuring range 0.200 in.) was used at a strain rate of 0.02 in. per in. per min. The work-hardening exponents were determined from the stress-strain curves generated in the smooth tensile tests and the assumption that the portion of the stress-strain curve beyond the yield point can be described by the power relationship: where a is the true stress, P is the true plastic strain,

Jan 1, 1969

By V. Pasupathi, R. E. Hummel, J. W. Koger

Specimens of P1 brass were plastically deformed at room temperature to various degrees of deformation and subsequently cooled in order to transform them to low-temperature martensite. Deformation shifts Ms. A, , and the temperature of minimum resistivity to lower temperatures, and also decreases the temperature coefficient of electrical resistivity. These properties change rapidly up to about 15 pct reduction but vary very little with higher deformation. The possible relationships between martensite formed by deformation and the M, temperature of low-temperature martensite are discussed. Evidence is given that deformation martensite delays the formation of low-temperature martensite. BETA' brass undergoes at least two different types of martensitic transformations. One of these transformations (B1- B2) was first observed by Kaminski and ~urdjumov' and occurs when 81 brass with a zinc content between 38 and 42 wt pct (quenched from the single-phase region) is cooled below room temperature. Jollev and Hull' determined the structure of 0" from X-ray and electron-diffraction data as ortho-rhombic. Kunze came to the conclusion that the super-lattice cell of 0" is one-sided face-centered triclinic (pseudomonoclinic). The second martensitic transformation (B1-A1) occurs when the specimens are deformed at or somewhat above room temperature. This type of martensite will be called deformation martensite. Horn-bogen, Segmuller, and Wassermann4 determined the structure of deformation martensite to be bct. (An intermediate phase, az, occurs before the final phase appears.) At deformations higher than 70 pct, a, transforms into a.4 A critical temperature Md exists above which no transformation occurs during deformation and is estimated to be around 400°C in P1 brass.5 This martensite has elastic properties.6 When the sample is stressed, martensitic plates appear; when the stress is released, the plates disappear. The present paper studies the effect of deformation martensite on the formation of low-temperature martensite. The experiments involved samples of 8, brass which were plastically deformed by various amounts and were subsequently cooled below the transformation temperature. EXPERIMENTAL PROCEDURE The 13 brass investigated was made from 99.999 pct pure copper and 99.9999 pct pure zinc and contained 38.8 wt pct Zn. The specimens, consisting of foils 0.1 mm in thickness, were heat-treated at 8'70°C for 15 min in an argon atmosphere and then quenched into ice water. They were then deformed by cold rolling and subsequently cooled at a rate of 1°C per min. The martensitic transformation that occurred during cooling was followed by electrical resistivity measurements. The resistance measurement technique and its accuracy have been described in a previous paper. Because the transformation 81 —-8" occurs below room temperature, the samples were placed in a cryo-stat which contained isopentane as a cooling medium. The isopentane was cooled by liquid nitrogen pumped under pressure through a 15-ft coil of copper tubing which was immersed in the isopentane. The nitrogen flow was regulated by a temperature controller using two thermistors in the cooling medium. The cryogenic liquid could be heated with an immersion heater. The useful temperature range with this device was from +25° to approximately -155~C. EXPERIMENTAL RESULTS Resistivity Measurements. The following abbreviations are used in this paper to label the characteristic temperatures during the martensitic transformation. M, is the starting point of the martensitic transformation and is defined as that temperature where the resistivity vs temperature curve on cooling first deviates from a straight line. Mf is the temperature at which the martensitic transformation is completed. On reheating, the transformation from martensite to the parent phase starts at a temperature A, and ceases at a temperature Af. Fig. 1 presents five different resistivity vs temperature curves corresponding to the transformation of brass from Dl to 8" after different degrees of reduction in thickness. The following observations can be made from these curves. 1) With increasing degree of deformation the Ms temperature is shifted to lower temperatures. This shift ranges up to 35°C compared to the undeformed state. This is also indicated in Fig. 2, where AM, (the shift of Ms, compared to the undeformed state) is plotted vs the degree of deformation. AM, increases rapidly until a reduction of about 15 pct is reached. With higher deformations, no additional increase in AM, was found. 2) With increasing degree of deformation the temperature of minimum resistivity (M) is also shifted to lower temperatures. The shift, attains a maximum of about 61°C compared to the undeformed state. In Fig. 3, AM is plotted as a function of deformation. It can be seen that, as in 1 above, AM increases rapidly and no further shift of M occurs for deformations greater than 15 pct. 3) The temperature coefficient of resistivity, is given by the slopes (dp/dT) of the linear portions of

Jan 1, 1969

By Amitava Gangulee, Michael B. Bever

The effects of changes in short-range order on some thermodynamic, electrical, and mechanical properties of the silver-rich Ag-Mg solid solutions have been investigated. The heats of formation at 273°K of several alloys after two different thermal treatments were measured by liquid metal solution calorimetry. Their electrical resistivities were measured at temperatures ranging from 4° to 773°k. Tensile tests and microhard-ness measurements were carried out at room temperature. An effective short-range order parameter and an effective interaction energy are defined. The measured changes in the properties are interpreted in terms of these quantities. The short-range order parameter is evaluated separately from calorimetric and X-ray data; the results are in fair agreement. Increasing short-range order lowers the resistivity of- the silver-rich Ag-Mg alloys but does not measurably affect the temperature dependence of the resistivity. Short-range order increases the yield stress of these alloys but does not affect the ultimate tensile stress. Changes in the effective short-range order parameter independently obtained from measured changes in the heat of formation, resistivity, and yield stress are in fair agreement. THE thermodynamic, electrical, and mechanical properties of solid solutions have been investigated as functions of composition in many alloy systems, but the effects of configurational variables, such as short-range or long-range order, on these properties have received much less attention. The silver-rich terminal solid solutions in the system Ag-Mg are well-suited for the investigation of the effects of order. The composition of these solid solutions may be varied over a range extending to about 27 at. pct Mg.1 Appreciable short-range ordering is likely to take place in these solid solutions; the degree of short-range order can be varied by suitable thermal treatments. In the alloy Ag3Mg long-range ordering is possible.2 This paper is primarily concerned with the effects of short-range order on some thermodynamic, electrical, and mechanical properties of the silver-rich Ag-Mg solid solutions.-The long-range ordering transition is the subject of a concurrent paper.3 1) EXPERIMENTAL PROCEDURES 1.1) Preparation of Specimens. Ag-Mg alloys containing up to 26.4 at. pct Mg were prepared by melting 99.99 pct pure Ag (obtained from Baker Chemical Co.) and 99.99 pct pure Mg (obtained from Eastman Chemi- cal Co. or Johnson Mathey & Co.) in graphite crucibles under molten potassium chloride. The solidified ingots were homogenized at 823°K for 10 days in evacuated Vycor capsules. A surface layer 6 in. thick was then removed by machining. The ingots were swaged into rods (2.5 mm diam) which were drawn into wires (1.0 mm diam). The wire specimens were annealed at 773°K for 24 hr in vacuum and either quenched into iced brine or cooled to room temperature over a period of 15 days. The average grain diameter (obtained from the linear intercepts) was about 0.1 mm. 1.2) Calorimetry. The heats of formation of the alloys were measured in a tin solution calorimeter as the difference in the heat effects of alternate additions of samples of an alloy and the mechanical mixture of its component elements. The additions were made from 273°K to the bath at about 623°K. The heat effect on addition of the alloys was approximately 5.5 kcal per g-atom. The procedure and the method of calculation have been described.4 1.3) X-Ray Diffraction. Short-range order parameters of two alloys were calculated from diffuse scattering intensities,' obtained with briquettes made from powdered specimens. A GE-XRD 5 diffractometer with filtered copper radiation was used. The absolute intensities were based on calibrations with paraffin. 1.4) Resistivity Measurements. The electrical resistivities were measured by a potentiometric method in which the potential drop across the specimen was compared with that across a standard resistance. For measurements below room temperature, the specimens were immersed in liquid helium (4°K), liquid nitrogen (78°K), dry ice and trichloroethylene (195°K), or ice and water (273°K). Measurements above room temperature were made on specimens held in a furnace under helium or in vacuum. The resistivity measurements were reproducible to ±0.5 pct or less. 1.5) Mechanical Tests. Tensile tests were carried out at room temperature with a Tinius-Olsen XY elec-tromatic Universal testing machine. Wire specimens (nominal gage length 1 in.) were used. The strain was measured with an extensometer, which had a sensitivity of 4 x 10-5; the strain rate was 10-3 min-1. Microhard-ness measurements were made with a square-pyramid indenter and a 100-g load.11 2) RESULTS AND DISCUSSION All measurements were made on two parallel sets of specimens, one quenched and the other slowly cooled from 773°K. The reported values are the averages of at least three measurements. In the range of magnesium concentration substantially below that of the composition Ag3Mg, quenched alloys have less short-range order than slowly cooled alloys. As the magnesium concentration approaches that of Ag3Mg, slow cooling develops long-range order. Quenching sup-

Jan 1, 1969

By Amitava Gangulee, Michael B. Bever

The order-disorder transition in Ag-Mg alloys in the range 17 to 26 at. pct Mg was investigated and some thermodynamic, electrical and mechanical properties of ordered Ag-Mg alloys were measured. A modification of the phase diagram is proposed on the basis of measured transition temperatures. The stability of the ordered structure is analyzed in terms of the quasichemical theory. Kinetic aspects of the order-disorder transition were also investigated. The energies of ordering at 273°K of Ag-Mg alloys were measured by liquid metal solution calorimetry. The electrical resistivities and the tensile Properties of the ordered alloys were measured. The exothermic energy of ordering increased with magnesium concentration up to the composition Ag3Mg; it is discussed in terms of the quasichemical theory. Pronounced order hardening was observed and can be explained by several strengthening mechanisms, which are particularly effective because of a change in crystal structure associated with the ordering transition. IN silver-rich solid solutions containing approximately 17 to 26 at. pct Mg, long- range ordering may occur during appropriate therma1 treatments. Clare-brough and Nicholas1 first detected such long-range order in the solid solution of composition Ag3Mg and suggested that the unit cell of the superlattice contained 256 atoms. x-ray2 and electron diffraction3 investigations confirmed the occurrence of long-range order, but indicated a shifted Ll2-type structure. On the basis of recent X-ray measurements the crystal structure of ordered Ag3Mg has been confirmed as type D023.4 The present paper is concerned with the order-disorder transition in silver-rich Ag-Mg solid solutions and some thermodynamic, electrical, and mechanical properties of the ordered alloys. The effects of short-range order on the properties of silver-rich solid solutions containing up to 26 at. pct Mg are discussed in a concurrent paper.5 1) EXPERIMENTAL PROCEDURES 1.1) Preparation of Specimens. Specimens of Ag-Mg alloys containing from 18 to 26 at. pct Mg were prepared in the form of 1.0-mm-diam wires as described in Section 1.1 of Ref. 5. Disordered specimens were prepared by annealing at 773°K for 1 hr and quenching into iced brine. Ordered specimens were prepared by annealing at 773°K for 1 hr and slowly cooling to room temperature over a period of 15 days. The specimens were stored at 78°K. 1.2) Resistivity Measurements. Electrical resistivi- ties were measured by a potentiometric method.' Equilibrium values were obtained at several temperatures. The kinetics of ordering were investigated by following the time-dependent changes of the resistivity of initially disordered specimens during annealing. The specimens were enclosed in copper capsules and immersed in salt pots; the heating up was accelerated by injecting preheated helium into the capsules. 1.3) Calorimetry. The energies of ordering of the alloys were measured in a tin solution calorimeter as the difference between the heat effects of additions of completely ordered and disordered specimens of the same composition. The procedure and the method of calculation have been described.6 Magnesium was used for thermal compensation in most calorimetric runs in order to improve the accuracy.677 1.4) Mechanical Tests. Tensile tests were carried out with wire specimens at room temperature as described in Section 1.5 of Ref. 5. Microhardness measurements were also made.? 2) RESULTS AND DISCUSSION The characteristics of the order-disorder transition of silver-rich Ag-Mg solid solutions will be discussed first. Some thermodynamic, electrical, and mechanical properties of the ordered alloys will then be considered and compared with the corresponding properties of dis-disordered alloys. 2.1) The Order-Disorder Transition. 2.1.1) Thermodynamic and Structural Aspects. The transition temperatures Tt were determined from discontinuities in the slope of the equilibrium resistivity vs temperature curves. Normalized curves for three compositions are shown in Fig. 1. The transition temperature increases with the magnesium concentration and reaches a maximum at the stoichiometric composition Ag3Mg. The transition temperature of the alloy Ag3Mg was measured as 665° ± 2°K and compares with published values of 660°,1 663°,8 and 66°k.3 In the slopes of the resistivity vs temperature curves of alloys containing 22.2 and 22.5 at. pct Mg, discontinuities were observed at two temperatures. Such upper and lower discontinuities indicate a two-phase field. A modified form of the published phase diagram9 is shown in Fig. 2. A two-phase field was found only on the low-magnesium side of the composition Ag3Mg. On the high-magnesium side, the existence of a two-phase field could not be established because of insufficient resolution, but such a field must be present. This part of the phase diagram can be made complete by a eutec-toid-type reaction a = (a' + ß). The existence of a boundary between the two-phase fields (a + ß) and (a' + ß) is also in accord with published lattice parameters.3 The crystal structure of ordered Ag3Mg (type Do23)

Jan 1, 1969

By M. Yokoi, K. Yoshida, T. Ishihara

The solubility of aluminum in cryolite melts has been determined by measuring the weight loss of aluminum submerged in melts contained in closed silicon nitride crucibles. The solubility in pure cryolite has also been checked with radioactive isotopes. The solubility decreases with increasing A1203 concentration, decreasing temperature, and decreasing cryolite ratio. The change of aluminum solubility in pure cryolite with the activity controlled by using Al-Ag and A2-Cu alloys supported the assumption that aluminum would dissolve in the form of monovalent ions according to the reaction 2AI + A1' —3 A1'. DURING the electrolytic production of aluminum, the electrodeposited metal redissolves in the bath in some form and is subsequently oxidized back to the former ionic state. This dissolution is believed to cause a decrease in current efficiency. Reported data on the solubility of aluminum in cryolite melts are scattered. The mechanism of the dissolution has not yet been determined, although certain tentative interpretations have been proposed. Many previous studies were performed under different test conditions. Studies were performed in open Systems and closed systems,'o and some were apparently not given sufficient time to Saturate." The solubility can be determined either analytically or gravimetrically. The analytical determination involves measurement of hydrogen gas evolved when frozen cryolite bath saturated with aluminum is treated with sodium hydroxide or hydrochloric acid. The gravimetric determination, adopted by many investigators,'-37'3>14 involves measurement of weight loss of the metal after equilibration with a known amount of salt. It appears that the former technique yields much smaller values of solubility than the latter technique. This large difference in results strongly suggests that the dissolved metal in the cryolite melt is no longer in the metallic state. In the present study, the solubility was determined by the gravimetric technique at various activities of aluminum using A1-Ag and Al-Cu alloys. An attempt was also made to determine the ionic (or atomic) state of the dissolved metal from the solubility data. 1) EXPERIMENTAL Apparatus. Preliminary tests demonstrated that silicon nitride was suitable as a container material since it is compatible to both the fluoride melts and molten aluminum. Crucibles were made from silicon powder nitrided with Nz gas at 1600°C. The good compatibility between crucible and melts is due to the near perfect nitrogenation of silicon. Only small amounts of silicon were found in the aluminum metal after the experiments. The crucible dimensions were 90 mm inside height, 38 mm ID, 105 mm outside height, and 50 mm OD. Close-fitting lids made of silicon nitride were also prepared. The crucible was placed in a close-fitting graphite crucible heated by a vertical Kanthal furnace, Fig. l. The lid and crucible are fastened tightly together by the condensation of vaporized bath in the narrow gap between them. The temperature was measured in the wall of the graphite crucible and controlled within * 5°C. The temperature difference between the measuring point and the interior of the melts was initially determined. Materials. The materials used were: Greenland cryolite (impurities CaF,, 0.5 pct; AlzO3, 0.2 pct; high-purity aluminum metal (purity 99.99 pct); A1203 and AlF, (provided by Nippon Light Metal Co., Ltd); reagent-grade NaF and CaF2; and 99.9+ pct Ag and Cu. Powdered reagents were carefully dried before use. A1-Ag and A1-Cu alloys were prepared before the experiments; their composition was checked by analysis

Jan 1, 1969

By M. J. Pool, P. J. Spencer

Using liquid tin as the solvent metal, the heats of solution of silver and gold in dilute Ag-Au-Sn alloys have been determined at 723°K by solution calorimetry. From the values obtained, the various enthalpy interaction coefficients in the liquid solutions have been calculated. The self-interaction coefficients Au and Ag values of 2413 and -6322 cal per g-atom in binary Au-Sn and Ag-Sn alloys, respectively. For an alloy containing 0.04 mole fraction Ag, Au is equal to 2443 cal per g-atom, and in an alloy containing 0.04 mole fraction Au, Ag has a value of —8883 cal per g-atom. Ag Au has been determined as —2650 cal per g-atom. ADDITION of a second solute to a dilute binary alloy may cause striking changes in the thermodynamic characteristics of the alloy. As yet, such influences cannot be reliably predicted and it is only by accumulation and critical assessment of data that definite patterns can be established. Solute interactions in dilute solutions were first treated mathematically by Wagner,1 who derived expressions for the Gibbs energy interaction coefficients. For an alloy consisting of. two dilute solutes, i and j, in a solvent s, these coefficients are given by: where yi is the activity coefficient of component i and x represents the mole fraction of solute or solvent. Other interaction coefficients have since been established in terms of the various partial molar thermodynamic properties of an alloy; the definitions of these are summarized in a recent paper by Lupis and Elliott.2 Mathematical and experimental studies of such coefficients have received much attention recently, providing a very useful theoretical background to an understanding of the nature of solute interactions taking place in dilute solutions. One of the definitions to arise from such work has been the extension of Wagner's treatment by Lupis and Elliott4 to include an enthalpy interaction coefficient, defined as. where ?Hi is the relative partial molar enthalpy of component i. ?Hi is related to the relative partial molar enthalpy at infinite dilution, ?Hio, and to the enthalpy interaction coefficients by the expression: Values of ?Hi can often be obtained with a high degree of accuracy by calorimetry, and, if the composition of the alloy is carefully controlled by making calculated solute additions throughout a series of measurements, application of Eq. [2] will enable reliable enthalpy interaction coefficients to be determined. Partial heats of solution must be investigated as a function of solute concentrations down to xi = 0, and, since ? values are obtained from the slope of mi plotted against x, great care must be taken to obtain precise data. The data for binary systems which are presently available often require extrapolation of ?Hi values from compositions as great as xi = 0.05 or even 0.1 to obtain limiting heats of solution. Reliable enthalpy interaction coefficients are thus impossible to obtain from published data for most of the alloy systems already investigated. The aim of the present work was to determine values of Au, Ag, Ag, and Au in the limit xsn -1 for the relatively simple Ag-Au-Sn system and from these results to gain some insight into the possible influence

Jan 1, 1969

By Richard W. Greene, Joel S. Hirschhorn

MATERIALS containing very large amounts of porosity, in the order of 50 to 90 vol pct, are extremely difficult to evaluate according to conventional mechanical behavior tests. This is usually a consequence of their low strength and brittleness, and the difficulty of preparing standard test specimens. A simple fatigue test has been developed and used successfully on porous nickel sheet materials requiring no specimen preparation; it could also be used for configurations other than sheet or plate. Samples of the sheet. 2 in. sq and about 0.025 in. thick, are placed in a vessel that is evacuated to about 50 f Hg by a mechanical pump. Distilled and deairated water is allowed to enter the vessel thereby filling the porosity (virtually all the porosity in these nickel ma-

Jan 1, 1970

By E. A. Metzbower

STEREOGRAPHIC projections of hcp metals depend upon the c/a ratio. Hu1 has recently published standard projections for titanium (c/a = 1.5873) of the (000l), the basal plane, the (1010), the most prominent slip plane, and the (1012), the most prominent twinning plane. Caplan et a1.2 and Hul have also published tables of angles between poles of planes, between directions, and between directions and poles of planes. From these tables standard projections and double stereograms can be constructed. However, by using the on-line CALCOMP plotter with a CDC 3800 computer, direct plots of standard stereographic projections of titanium (c/a = 1.587) have been made.

Jan 1, 1970

By A. U. Seybolt, R. E. Honnemon

THE purpose of this communication is to present activity data of nickel in the Ni-A1 binary system that was obtained for oxidation studies in this system. This data has been calculated from previous aluminum activity data at 1000°C1,2 by two types of Gibbs-Duhem integrations. These results should be of particular current interest to oxidation studies of aluminide coated nickel-base superalloys3 as well as other problems. The aluminum activity of Steiner and Komarek1 is plotted in Fig. 1 as a function of composition in Ni-A1 alloys at 1000°C. The dashed parts of the aluminum curves represent our approximate extrapolations from the directly measured regions. In the terminal solution regions (a and L phases) standard parabolic solution functions of the following form are assumed: where i is the component subscript, Ki is a constant at a fixed temperature, and xi is the atom fraction. The resultant nickel activity curve at 1000° C from the Gibbs-Duhem integration4 of the aluminum data is shown in Fig. 1. The solubility ranges of the phases present in the Ni-A1 binary at 1000°C are also denoted at the bottom of Fig. 1. Corresponding values of the nickel activity, aNi and yNi, are listed in Table I at useful compositional intervals. including phase boundary values. The aNi and yNi data listed in Table I are subject to the uncertainties present in the experimental aluminum activity data1 and to deviations from parabolic solution behavior in the terminal solutions. The errors introduced into aNi and yNi due to the parabolic solution assumption are expected to be relatively small (<20 pct of the values given) in all the phases except the liquid phase. Some deviation from parabolic solution behavior in the liquid phase is indicated

Jan 1, 1970

By W. White, C. S. Barrett, R. M. Mueller

THIS communication reports film-recorded low-energy (-100 kev) proton blocking patterns&apos;-3 for hcp and wurtzite crystals. In particular the sequence of relative line intensities* (corresponding to planar blocking) of common hcp metals is compared with those of the wurtzite semiconductors CdS and CdSe. The results show that the sequence is characteristic of structure. For closely similar structures such as those of CdS and CdSe, differences of line intensity attributable to the difference in atomic number, 2, of the scattering centers may also be seen. The apparatus and techniques used in this work were identical with those reported previously in a study of cubic systems.3 All metals were given standard metal-lographic etches, and the semiconductors cleaved, immediately prior to introduction to the vacuum system. Fig. 1 shows the unit cell of the wurtzite structure, and the geometry used to obtain the [1120] and [0001] patterns of Fig. 2. The axial ratio, c/a, for all ma- terials of Fig. 2 is close to 1.63. The patterns are then directly comparable in linear dimensions inasmuch as they have been reproduced at the same overall magnification of the patterns. (Overall magnification is proportional to the product of film-to-target distance and printing magnification. Thus the 1/6-in.-diam beam passage hole in the center of the film need not appear of uniform size in the prints.) The patterns for a given orientation, though similar

Jan 1, 1970

By R. L. Skaggs, R. A. Molsberger

DARKEN1 has pointed out that in most binary liquid metal systems the thermodynamic behavior is relatively simple in the terminal regions. Between the terminal regions. the thermodynamic behavior is more complex. To describe this behavior. he has defined the stability and the excess stability as shown: Qualitatively, an inflection in the free energy of mixing curve results in a change in sign, from positive to negative, of the stability function and is associated with a phase separation. Conversely, a large, positive stability over a narrow composition range is indicative of compound formation. Darken has shown a number of binary systems where the excess stability is constant in the terminal regions and varies sharply in between. The work described in this communication was undertaken to study the variation of stability and excess stability in a binary system as a function of temperature. In addition, this work provides means for the direct comparison of previously existing equilibrium (activity) and calorimetric (molar heat of mixing) data on the Hg-Sn system. Hildebrand2 measured the vapor pressure of mercury over Hg-Sn alloys at 595°K. Calorimetric studies were made by Kawakami3 at 523°K and by Kleppa4 at 423°K. In this work, the equilibrium vapor pressure of mercury was measured over a series of Hg-Sn alloys at 453" and 503°K by use of fused salt manometer with one arm exposed to mercury vapor. Deflections of the fused salt manometer were balanced by argon pressure on the other arm. Argon pressure was subsequently measured with a dibutyl phthalate manometer. Tin activities were determined by graphical integration of the Gibbs-Duhem equation making use of the familiar Darken5 a function. The activities of mercury and tin showed positive deviation from Raoult&apos;s law at 453 and 503°K as Hildebrand observed at 595 K. In Fig. 1 are plotted log activity of mercury and tin vs reciprocal temperature. The plot contains activities obtained from calorimetric studies. Hult-gren6 has calculated activities from Kleppa&apos;s4 data, and Kubaschewski7 has similarly treated Kawakami&apos;s7 data. Presumably. in each case a regular solution was assumed. For regions where a two-phase field

Jan 1, 1970

By H. R. Thresh, A. F. Crawley

THIS paper reports the results of viscosity measurements on cadmium and indium undertaken as part of a continuing program of liquid metals research at the Mines Branch, Department of Energy, Mines and Resources, Ottawa, Canada. For several years this work has continued with support from the Canadian Zinc and Lead Research Committee and the International Lead and Zinc Research Organization. In the literature, existing viscosity data for cadmium are conflicting and for indium are scarce. To achieve reliable viscosity values, further work was deemed desirable. Viscosities were measured using an absolute technique applied to an oscillating closed right cylinder of liquid metal. The apparatus, constructed in this laboratory, and the techniques involved in its use have been described in an earlier paper.&apos; In calculating viscosity coefficients from the experimental data. Roscoe&apos;s2 mathematical analysis was used. Because accurate viscosity determinations are dependent on accurate density data, and since the available data in the literature are conflicting, the densities of liquid cadmium and indium were also determined in this laboratory by a pycnometric technique. This increased the control exercised over all the variables contributing to the accuracy of viscosity measurements. For both density and viscosity determinations, cadmium and indium of 99.999 pct purity were used. These materials were obtained from Cominco Ltd. In order to suppress excessive evaporation losses and help prevent oxidation, a technique developed in measuring the viscosity of zinc&apos; was used for cadmium. This involved maintaining an atmosphere of purified hydrogen in the viscosimeter during the rela- tively long periods while the experimental temperature was changed and stabilized between individual determinations of the logarithmic decrement. Then, immediately prior to a decrement measurement. the viscosimeter was rapidly evacuated to a pressure of about 2 X 10-6 torr. On the other hand. with indium. no difficulties arise from vaporization or oxidation. In this case. therefore. the apparatus was held under a constant vacuum of about 2 x 10-6 torr. When all sources of random and systematic error have been evaluated, the technique is considered to yield data accurate within 10.5 pct. The viscosity of cadmium was measured to a maximum temperature of 451.8°C and that of indium to 346.5°C. These temperatures are. respectively. 130° and 190°C above the melting points. Two separate runs were made on indium. The results are given in Table I and, when plotted. yield the smooth curves evident in Figs. 1 and 2. Values of A and C to fit the well-known Andrade1 equation were determined by a regression analysis. The equation is: where n = coefficient of viscosity. i = specific volume. T = temperature in °K Despite its empirical nature. Andrade&apos;s equation has been shown by many workers to describe accurately the temperature dependence of viscosity. During the current research program, this equation has been found to fit the data precisely for all the metals and alloys that have been studied.5,6 The A and C&apos; values for cadmium and indium are given in Table 11. Comparing the magnitude of the present data with those of previous workers,7-10 as has been done in Fig. 1, it is evident that the results of Fisher and Phillips10 deviate markedly from the rest. With the exception of one point from Gering and Sauerwald,7 all other data agree within ±2.5 pct. However, the agreement between the present results and the work of Menz and sauerwald9 is especially good over the entire experimental temperature range. Such agreement may be expected since Menz and Sauerwald&apos;s

Jan 1, 1970

By C. R. Watts

The feasibility of producing composites containing nickel or molybdenum fibers in a beryllium matrix was inrestigated. The composties studied were jabricaled by powder mallurgical techniques. The 1-mil-diarr nickel fibers reacled completely below 900°C, converling the fibers .from nickel to Ni5Be2,. As the /LO/-pressing temperalure as raised above 1110oC, tlie nickel diffused outward from the beryllide fibers. The solid solubility of nickel in beryllium was clboul 20 wt pet at the 1100°C pressing temperature a1 the zone-fiber interface. The 1.5-mil-diam molybdenum fibers slzolred no evidence of reaction and little evidence of diffitsion after pressing at 900°C. Between 1000° and 1050°C pressing conditions, the fibers began lo react , producing 1ayers of MoBe2 and MoBe12, respectively surrounding the molybdenurn core. The struture remained the same at 1100°C with no evidenre of solid solubility of the molybdenum in the berylium or vice versa. In recent years a considerable amount of attention has been devoted to the determination of methods for improving the mechanical properties of materials through the use of fiber or whisker reinforcement. Previous work with metal matrix composites indicates that the study of the chemical compatibility of the fiber and matrix is an area requiring greater understanding. The metal-metal or ceramic-metal interface is frequently subject to chemical reactions that may result in the formation of hard brittle intermetal-lic compounds and/or low melting point eutectic compositions. The reaction products may reduce both the low-temperature and elevated-temperature strength of the composite by weakening the fiber-matrix bond, producing premature failure at the interface. It is well-known that most metal-metal systems and many metal-ceramic systems of interest for structural composites are thermodynamically unstable,&apos;-" particularly at elevated temperatures. If, however, the rate of reaction under the conditions of fabrication is sufficiently low. composites can be fabricated that can be used efficiently for indefinite periods at low temperatures and for short periods at elevated temperatures. This paper presents the results of a series of tests to determine the compatibility of nickel and molybdenum fibers with beryllium at various hot-pressing temperatures. Nickel was selected as a candidate fiber material primarily because the relatively ductile fibers might be useful as crack arresters in applications such as ballistic impact where crack growth can result in catastrophic failure. The high density and the reactivity of nickel were primary factors detracting from its selection as a possible reinforcement. Molybdenum with a modulus of elasticity of 52 Xlo6 psi is one of the few metallic materials having a modulus higher than beryllium (42 X lo6). Its high modulus, coupled with its refractory characteristics, made molybdenum an attractive candidate for a relatively stable fiber reinforcement for beryllium. Its density, being higher than that of nickel and over five times that of beryllium: detracted from its other characteristics. EXPERIMENTAL PROCEDURE The specimens were prepared from beryllium powder with a dispersed phase of fibers by powder metallurgical techniques. P-20 grade powder, Table I, from Berylco was used as the matrix material. Short lengths of 0.001-in.-nominal-diam nickel fibers supplied by the Sigmund Cohn Corp. and 0.0015-in.-nominal-diam molvbdenum fibers obtained from the General Electric Co. were used as the dispersed phase. The composite constituents were combined under an argon atmosphere by mechanically mixing the powders and fibers. The compositions used were nominally 1 vol pct fibers. After mixing. the composites were hot-pressed into a-in.-diam pellets under an argon atmosphere at 900°, 1000". 1050". and 1100°C at a pressure of 6000 psi with no hold time at these temperatures so that a comparison could be made between the resultant microstructure and hot-pressing temperature. The billet was heated at a rate on the order of 30°C per min to the desired temperature and then cooled at a somewhat slower rate. The microstructure obtained should be considered as characteristic of the integrated time-temperature history of the sample, as well as the maximum temperature attained. Upon removal from the hot-pressing dies. the specimens were cut. mounted. and polished by standard procedures. No etchant was used in specimen preparation. Photomicrographs, electron microprobe scans, and electron back-scatter pictures were made. X-ray dif-fractometer patterns were made of several of the specimens. but only the lines for beryllium could be resolved. Specimens for optical and electron microprobe examination were selected partially for the roundness of the cross section. A round cross section was taken to indicate that the body of the fiber was approximately normal to the surface and that therefore effects due to fiber material immediately below the surface could be neglected. RESULTS AND DISCUSSION The microprobe scans indicated that nickel reacted as low as 900°C, converting the entire fiber cross section to NisBe21. Fig. l(a). There was no evidence of further reaction from the optical or the back-scatter pictures, Figs. 2(n) and 3(a).

Jan 1, 1970

By Ashok K. Sinha

During the course of an in~*estigation into the occurrence of ordered AB3 structures, the following new phases have been found —CrRh3 (AuCu3 type), CrCo3 (MgCd3 type), HfCo4 (Ths Mn23 type), and WPt, MoPh type). The composition of the TiPt3-x phase (TiNi, type) is close to Ti23Pl77. The alloy chenzistry of transition rnetal AB3 structures is rezliewed in the light of electron concentration correlations of hex-agonality recently obtained for quasi-binary alloys. The relatizte colurne contraction in the AB3 structures increases with increasing difference in volume of the conzponents. A family of ordered close-packed layered structures is formed by stacking identical layers of composition AB, in various sequences, such that the coordination is twelvefold throughout and there are no A-A contacts. Previous work&apos; on quasi-binary AB3 alloys has led to the conclusion that the stacking sequence of the AB, structures changes with increasing radius ratio RA/RB from a purely cubic, through different mixtures of hexagonal and cubic stacking to a purely hexagonal stacking. However. for binary AB3 alloys, a correlation between the type of the crystal structure and the position of the components in the various volumns of the periodic table has been noted.2-5 It has been noted6 that this correlation appears to hold even though the radius ratio RA/RB may vary over a considerable range with the location of the components in the three long periods. Another study7" of several quasi-binary systems led to the conclusion that an increase in hexagonality of the stacking is associated with increase in the electron concentration e/a. as defined by the average per atom of the total number of electrons outside the inert gas shells. In apparent conflict with this conclusion, it is known that seven binary alloy structures isotypic with TiNi3 which is 50 pct hexagonal occur at a higher electron concentration (e/n = 8.5) than that (e/a = 8.25) for the 100 pct hexagonal MgCd3 type structure present in seven binary AB3 alloys. Table 111. In the present work, an investigation into the occurrence of binary AB3 structures in transition metal alloys was made, and a survey of binary AB3 structures is presented. EXPERIMENTAL The starting materials were pure metals of 99.9 wt pct purity. The alloys were arc-melted under partial pressure of argon and annealed in sealed silica capsules lined with molybdenum foil under argon at- mosphere. The total weight loss upon melting and subsequent annealing was always less than 1 pct and hence the alloys will be referred to by their intended (unanalyzed) compositions. Wherever the constitution permitted. the alloys were given a homogenizing treatment at 1200°C (3 days) prior to annealing. Unless otherwise stated all alloys were annealed at 900°C for 1 week and water-quenched. Sometimes the final annealing treatment was carried out on powders to accelerate the attainment of equilibrium. X-ray powder patterns were taken using a Guinier-de Wolff focusing camera (CuK, radiation) or an asymmetrical focusing camera (Co or CrK, radiation). For lattice parameter determination. internal silicon standards were employed. The intensity calculations were made using a Fortran IV program written by Jeitschko and parthe.9 RESULTS Twenty AB3 and three AB4 alloys were investigated. Table I lists the crystallographic data on some of the intermediate phases encountered in the present work. Table II contains the X-ray data for HfCo, (Th,,Mn,, type). The positional parameter, x. was assumed to be 0.378. the value for Th6Mnn2310 The X-ray pattern of ZrCo, was very similar to that of HfCo, and the previous structure determination of ZrCo, by Kuzma el al." was confirmed. Ordering in the alloy CrCo could be ascertained by the presence of only one weak super lattice line (101). the others being too weak presumably owing to the small difference in the scattering powers of chromium and cobalt. This line was observed in the X-ray pattern of powder from the massive sample annealed at 830°C (7 days) after the powder had been reannealed at 600°C (24 hr). The diffraction pattern of the powder similarly reannealed at 830°C (24 hr) contained only the lines due to a mixture of hcp and fcc Co(Crj solid solutions. Therefore, it appears reasonable to assume that O2 and/or N2 contamination which would be less likely to occur during the 600°C anneal was not responsible for the observed weak reflection. Also. this reflection cannot be identified with any of the strong lines of the neighboring s phase which is present in the Co-Cr system at higher chromium contents. The composition corresponding to the TiNi3 structure observed by Raman et al.12 in the two-phase alloy Ti,zt,, has been established in the present work as being between There was satisfactory agreement for the low-angle lines (up to d = 1.997A) between the observed diffraction pattern of TiCua and that calculated assuming the ZrAu, structure. as recently proposed by Pfeifer-et a1.I3 However. some of the superlattice lines. e.g., at d = 1.937 and 1.919A. predicted by the ZrAu, structure were not actually observed eve? though neighboring lines. at d = 1.947 and 1.986A. of comparable calculated intensity were present. The ZrAu

Jan 1, 1970

By D. T. Peterson, T. Carnahan

The diffusion coejTicients of carbon, nitrogen, and oxyget were determined in $ thorium over the tempernilcre range 1440" io 1715°C. The diffusion coyfiicir?zls are given by: D = 0.022 exp (-27,000/RT) jor carbo)~, D = 0,0032 exp(-l7,00Q/RTj for nitrogen, and D = u.0013 expt(-11,UOU/RT) for oxygen. Cavl~orz was found to increase the hardness of thoriunz nearly linearly with concentration over the range 100 to 1000Ppm carbon. ThORIUM has a fcc structure up to 1365°C and a bcc structure from this temperature to its melting point at 1740°C. Diffusion of carbon, oxygen, and nitrogen in bcc thorium was of interest in connection with the purification of thorium by electrotransport.&apos; In addition, it was possible to measure the diffusion of all three of these interstitial solutes in the same bcc metal. Only in niobium, tantalum, vanadium, and a iron have all three interstitial diffusion coefficients been measured in a given bcc metal. Diffusion coefficients have been measured for carbon and oxygen in a thorium by Peterson2, 3 and for nitrogen by Gerds and Mallett.4 Activation energies for diffusion are reported by the above authors to be 38 kcal per mole for carbon, 22.5 kcal per mole for nitrogen, and 49 kcal per mole for oxygen. Values of the diffusion coefficients of carbon and nitrogen in 3 thorium have been reported by Peterson et al.&apos; However, these were secondary results of their investigation of electrotransport phenomena in thorium and it was hoped that the present study could provide more precise data. EXPERIMENTAL PROCEDURE The specimens used in this study were the well-known pair of semi-infinite bar type. The couple was formed by resistance butt welding two 0.54-cm-diam by 3.0-cm-long bars of thorium together under pure helium, the concentration of the solute being greater in one cylinder than that in the other. The finished couple then contained a concentration step at the weld interface and diffusion proceeded only along the axis of the rod. The thorium used in this study was prepared by the magnesium intermediate alloy method.5 The total impurity content was less than 400 ppm. The major impurities were: carbon, 100 ppm: nitrogen, 50 ppm; and oxygen. 85 ppm. The total metallic impurity content was less than 150 ppm. The high solute concentration portions of the diffusion couples were prepared by adding the solute to the high-purity thorium in a non-consumable electrode arc melting procedure. Carbon and nitrogen were added in the form of spectroscopic graphite and nitrogen gas while a Tho2 layer was dissolved by arc melting to add oxygen. High-purity thorium formed the low concentration portions in the carbon and nitrogen couples. The low oxygen portions were obtained by deoxidizing high-purity thorium with calcium for 3 weeks at 1000°C according to a method reported by Peterson.3 The high C-Th contained 400 ppm C, the high N-Th contained 400 ppm N, the high 0-Th contained 220 ppm 0, and the low 0-Th contained 25 ppm O. The high O-Th was brine-quenched from 1500°C to retain most of the oxygen in solution at room temperature. These concentration levels were all below the solubility limits in 0 thorium at 1400°C. A resistance-heated high-vacuum furnace was used to heat the couples. The samples were mounted horizontally on a tantalum support which had small grooves near each end. Spacer rods of thorium, 0.4 cm in diam, were placed in these grooves to prevent contact between the sample and the tantalum support. This arrangement should have prevented contamination of the sample by contact with the support. In further effort to reduce contamination, the oxygen diffusion couples were sealed inside evacuated outgassed tantalum cylinders lined with thorium foil. Thorium rings around each end of the samples acted as spacers in this case. Pressure during diffusion runs was about 10-6 torr after an initial outgassing stage. Temperature measurements were made by sighting on black body holes in the sample support adjacent to the samples with a Leeds and Northrup disappearing-filament optical pyrometer. Temperatures were constant during a diffusion anneal to ±5C. The observed temperatures were corrected for sight glass absorption after each diffusion run. The pyrometer was checked against a calibrated electronic optical pyrometer and a calibrated tungsten strip lamp with the electronic pyrometer being taken as the standard. All temperature readings agreed to within ±3C over the temperature range 1450" to 1690°C. Time corrections due to diffusion during heating and cooling were necessary because of the short diffusion times. The diffusion times ranged from 6 min for the oxygen sample run at 1690°C to 90 min for the carbon sample run at 1500°C. A series of temperature vs time plots were made for heating and cooling of the samples to the various diffusion temperatures. This data was then used in a method according to shewmon6 to determine the time corrections. The corrections amounted to

Jan 1, 1970

By H. J. Levinstein, H. J. Guggenheim

RbFeF3 is a transparent ferromagnet with a large faraday rotation which permits the direct observation of magnetic domain structures in bulk crystals. If the position of dislocations within the crystal were known a study of domain wall dislocation interactions would be possible. In this paper an etch pitting solution is described which gives a one-one correspondent between the etch pits and the dislocations intersecting the (100) surfaces of RbFeF3 crystals. In addition, RbFeF3 is shown to have sufficient plasticity so that known arrays of dislocaticms can be put into the crystal. It is concluded that RbFeF, is an ideal material for the investigation of domain wall dislocation interactions. THE influence of lattice defects on the motion of domain walls in ferromagnetic materials has been studied extensively both theoretically and experimentally.4 The experimental data obtained has been qualitative in nature due to the lack either of a quantitative knowledge of the magnetic domain array in the material or of the nature of the dislocation configuration in the material. At room temperature RbFeFJ has the cubic perovskite structure. It is ferromagnetic below 86°K.5,6 It is also transparent and has a large faraday rotation in the magnetic state.7 The domain structure of this material is observable in three dimensions in crystals which are millimeters in thickness.&apos; Since RbFeF3 satisfies the requirements of having an easily visible three-dimensional magnetic domain structure, it was of interest to determine if the dislocation structure of the material could be determined and if the material was sufficiently plastic so that dislocations could be introduced into the as grown crystals. This is a report of the results of that investigation. EXPERIMENTAL PROCEDURE The crystals of RbFeFs used in this experiment were grown by a technique described elsewhere.5 In order to obtain flat strain-free surfaces, the samples were first cut to the crystal orientation desired and the surfaces were polished by employing standard metallographic techniques. The samples were then immersed in a solution of 1 to 3 g of KOH in 100 ml of water. This solution removes the damaged surface layers uniformly leaving a film on the surface of the sample. Etch pits are obtained by immersing the sample in an approximately 50-50 solution of 30 pct peroxide and concentrated acetic acid. The film formed during the KOH treatment turns deep red in color in the etching solution and flakes off the sample. The etching time required depends upon the pit size desired and

Jan 1, 1970

By D. G. Westlake

Resistance was measured as a function of temperature, 77° to 350°K,for niobium with hydrogen concentrations up to 3.76 at. pct. Cooling curves exhibited a discontinuity in slope that was attributed to the initiation of hydride formation. A thermal hysteresis was explained in terms of diffusion and the effects of mechanical constraint on nucleation and growth. The heat of solution of hydrogen in saturated Nb-H solid solution was 2. 737 * 0.045 kcal per g-atom, including an unknown stress contribution. The critical ternperatures for hydride formation during cooling and for embrittlement are correlated but not coincident. This, too, has been explained in terms of mechanical constraint and stress-itlduced tratzsfomtation. Nb-H alloys undergo a ductile-brittle transition with decreasing temperature.&apos;-3 Similar behavior has been noted for v-H.1,5 For the latter system, there is evi-dence6 suggesting coincidence of the temperature of minimum ductility and the critical temperature T, for initiation of metal hydride precipitation during cooling. A decrease in ductility due to the presence of brittle hydride particles is readily explainable. The fact that ductility begins to decrease at some temperature higher than T,4,5 has led us6 to suggest that precipitation may be strain-induced locally during tensile testing. Wood and Daniels&apos; have reported that a plot of temperatures for 10 pct reduction in ductility vs hydrogen concentration in niobium is not coincident with, but parallels, the solvus.7,8 They believe this indicates a correlation of precipitation and embrittlement. Unfortunately, the two reported solvuses 7,8 to which Wood and Daniels refer are not in agreement with each other. Furthermore, neither solvus was determined experimentally: each was a plot of calculated solubilities. Thus. neither is highly dependable for determining whether embrittlement coincides with initiation of hydride precipitation. Longson 9 studied phase relations in the Nb-H system between 77° and 473°K by dilatometry. Comparison with metallographic observations convinced him that discontinuities in the slopes of his dilatometric plots were coincident with the initiation of hydride precipitation. Plotting the logarithm of the hydrogen concentration vs the reciprocal of the temperature at which precipitation was initiated during cooling, he obtained a straight line shown in Fig. 1. Walter and chandler,&apos;&apos; using X-ray techniques to detect the for mation of a second phase? have obtained three experimental points in the temperature range 295° to 338°K which fall within Longson&apos;s experimental error for this plot. A similar plotg of temperatures at which dissolution was completed during heating was a parallel line shifted toward higher temperatures. A straight-line plot was obtained, also, for the temperatures of the ductile-brittle transitions,9 but it was shifted toward still higher temperatures. Longson calculated an activation energy of 2.980 * 0.300 kcal per mole for the hydrogen embrittlement mechanism, compared with 3.180 * 0.400 kcal per mole for hydride formation. Despite this agreement, he concluded, "It seems unlikely that embrittlement results from hydride formation..." because hydrogen embrittlement was observed at temperatures above the solvus. It is our contention that embrittlement is correlated with precipitation of the hydride.6 Because only Longson has investigated Nb-H phase relations at sub-ambient temperatures, we deemed it advisable to re-determine the temperature dependence of the solubility using the resistometric technique already proven successful for the V-H system.6 Coincidence of our solvus with temperatures for embrittlement would

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.&apos;-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&apos;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.&apos;

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-&apos; 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,&apos;-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&apos;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,&apos;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