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By R. J. Wasilewski

Elastic constants have been determined on single crystals of maximum-melting-temperature NiAl compound (50.6 at. pct Al) at 25°C. Temperature variations of Young&apos;s modulus in the three principal crystal directions between -150°C and +800°C exhibit no anomalies similar to those repovted for the isomor-phous compounds brass and AuCd. It is concluded that the ordered CsCl structure is stable throughout the temperuture range investigated NICKEL aluminide is a congruently melting inter-metallic compound of CsCl (B2) structure, stable over a rather wide composition range (40 to 54 pct Al). It exhibits a lattice-parameter maximum of a. = 2.883A at 50.6 pct 1,&apos; at which composition the excess aluminum atoms are randomly distributed on the vacant nickel sites;&apos; at higher nickel compositions the excess nickel atoms replace substitutionally in aluminum sites, while at higher aluminum compositions vacancies are present in the nickel sublattice.&apos; For compositions close to NiAl no significant variation in order has been observed at temperatures up to 900°C.3 At high nickel (65 pct) compositions the existence of a martensitic structural transformation has been reported.4 NiAl is isomorphous with the ordered B-brass structure and with the high-temperature structures of AuCd and TiNi alloys close to equiatomic composition. All of these exhibit anomalous elastic behavior as the temperature of their diffusionless transformation to lower-symmetry structures is approached.&apos;-&apos; This behavior has been explained as indicative of low stability of the CsCl structure in these compounds.8 The present paper reports the experimental determination of the elastic compliances of single-crystal NiA1. EXPERIMENTAL METHOD 1) Specimens. Polycrystalline NiAl was prepared by induction melting of carbonyl nickel and high-purity aluminum and casting into solid copper molds in gettered helium atmosphere. The castings thus produced were canned in mild steel and impact-extruded at 1100 to 1200°C to 38-in.diam rod. After the steel sheath was pickled off, the rods were converted to single crystals by repeated high-frequency satisfactorily produced the desired orientations in the form of crystals 38 in. diam by 8-12 in. long. These were subsequently center less-ground to 7 to 8 mm diam by 12 cm long, electropolished (in 80 pct methanol, 20 pct HzSO, at 25°C and 15 v), and their axial orientation determined to 1 deg by the Laue back-reflection technique. The rod dimensions were determined to 0.01 mm and the densities determined from geometrical dimensions and the weight of every specimen. The average density was 5.905 0.005 g per cu cm. 2) Elastic Data Measurements. Room-temperature elastic compliances Sii were obtained by determining to 0.1 cps the fundameGtal resonance frequencies in the longitudinal and torsional vibration modes. At these frequencies the following relations hold: E =p(2lfif =pvl= Sn- 2r[Sll- S12- -£„«] g -P(2iftf = pvf = (s<« - fr(sii - sw - ±s»y where vl and vt are the velocities, and f andft the resonance frequencies of the longitudinal and torsional waves, respectively; 1 is the length of the specimen, and r is the direction cosine factor—{cos2 a • cos2 0 + cos2 a . cos2 y + cos2 0 ¦ cos2 Y}. Given three crystals of orientations (loo), (110), and (111)) i.e., F = 0, 14, and 13, we can obtain six equations for the three unknown compliances, thus providing adequate internal check. Mode mixing can readily occur in the relatively long wavelength resonance utilized here. This could cause a significant error in the torsion wave velocities calculated from the apparent resonance frequenies. Therefore, throughout the present work, a method, described previosl,&apos; was used to ensure that no significant mode mixing took place at resonance. The apparatus used was originally designed for the Elastomat, Magnaflux Corp.. Chicago, 111. determination of flexural vibration resonance frequencies. At room temperature torsional vibration mode could be established by suitable positioning of the thin wires from the input and pick-up transducers, and longitudinal mode by cementing thin magnetic foils at the ends of the specimen, and inducing the vibration by means of electromagnetic transducers. Only flex-ural resonance frequencies could, however, be obtained reproducibly over the temperature range of -150 to 800°C. These data are considerably less accurate, the error in the modulus being estimated at -0.6 pct from the specimen dimensions and the room-temperature Poisson&apos;s ratio. The reproducibility of the resonance frequencies in repeat runs was, however, to better than l in 15,000, hence the relative elastic modulus error less than 0.02 pct. Therefore, the temperature variation of the principal Young&apos;s moduli obtained by flexural resonance, relative to room-temperature values, is believed to be satisfactorily accurate.

Jan 1, 1967

By Thomas L. Garrard, James A. Sprague, Rex B. McLellan, Samuel J. Horowitz

A simple rnodel for interstitial solid solutions has been devised in which each solute atom interacts with the solzlent lattice in such a way as to exclude an integral number of nearest-neighbor sites from being occupied by solute atoms. The vemaining interstitial sites ave not affected. In calculating the entropy of such a solution the increase in the number of interstitial sites available to solute atoms due to the ozlerlap of blocked sites is taken into account. The ther-modynamic functions and the solute activity fov the solution rnodel are calculated. It is shown that considevation of the overlap sues vise to a correction term of the order 8&apos; ( where 0 is the ratio of solute to solvent atoms) in the thevmodynamic functions of the solution and the solute activity. The theo,uetical equation for the solute activity was used to analyze the available data for the activity of carbon in austenite at three different temperatures. The use of a computer enabled many comparisons to be made between the experimental activity data and those predicted by both the complex blocking model and the simple, widely used blocking model in which the ozierlapping of blocked sites is not taken into consideration. Even though at the highest concentvations oj carbon the value of 0 is small it is shown that the best fit to the experimental data is found with the complex blocking model. RECENTLY it has been shown by analysis of gas-solid equilibriums that many interstitial solid solutions can be described by the quasi-regular solution model in which the partial energy of solution E, is independent of composition and, apart from a composition-independent term due to nonconfigurational factors, the partial entropy is that of an ideal solution. Dilute solutions of nitrogen in both bcc and fcc iron,&apos; dilute solutions of oxygen and silver,&apos; carbon in a Fe,&apos; and dilute solutions of carbon in ? Fe2 are all quasi-regular. In addition many hydrogen-metal solutions, both dilute and concentrated (with the notable exceptions of the H-Ni and H-Pd systems), can be described by this model.3 However, it has been shown that, at various temperatures, solutions of carbon in ? Fe begin to depart strongly from quasi-regular behavior at low concentrations. A plot of the activity of carbon in ? Fe relative to pure graphite vs atom fraction of carbon (Cc) calculated from Smith&apos;s data4 begins to depart from the (nonlinear) plot for a quasi-regular solution at about Cc = 0.02. Several statistical models have been proposed to account for the thermodynamic behavior of carbon in y Fe. Darken and smith4 set up a model in which a carbon atom in its octahedral site in the fcc lattice has either one or no neighboring carbon atoms in the twelve nearest-neighbor octahedral sites. It is assumed that the number of carbon atoms having more than one nearest neighbor is negligibly small and that a solute atom does not interact with another solute atom located at a greater distance than the first shell of interstitial sites. Thus this model contains two characteristic interaction energies. Darken and Smith showed that this model gives reasonable agreement with the measured variation of carbon activity ac with composition. It is noteworthy that Darken and Smith concluded from their analysis of the activity data using this model that there was a repulsive force between carbon atoms in austenite which reduces the concentration of C-C pairs below that corresponding to random mixing. Recently Aaronson, Domian, and pound5 have considered in detail the statistical model developed by Lachere and by Fowler and Guggenheim7 and have shown that it is compatible with the activity data for carbon in austenite. In Darken&apos;s model the energy of solution of carbon in austenite is dependent on temperature and concentration; the departure from Henry&apos;s law arises from the concentration dependence of the energy of solution. In Aaronson, Domian, and Pound&apos;s discussion of the Lacher, Fowler, Guggenheim model the interaction energy wy between adjacent carbon atoms in austenite is strongly temperature-dependent. Aaronson et al . suggested that this temperature dependence of wy and hence the energy of solution of carbon could be explained either by the formation of clusters of carbon atoms (despite the repulsive nature of wy) or by the dual occupancy of the octahedral and tetrahedral sites by the carbon atoms. However, a temperature-dependent heat of solution

Jan 1, 1968

By M. F. Burr, F. V. Lenel, G. S. Ansell

The mechanism of materials transport was studied in a model experiment in which a silver sphere was sintered to a silver plate under an applied external stress near 800° C. For this purpose the rate of growth of the neck between the sphere and the plate which is proportional to the strain rate of the creep process involved in a neck growth was determined by measuring the change in electrical resistance of the away. This rate of neck growth will depend upon the magnitude of the applied external stress. By measuring the rate just before and just after a change in the external load and therefore the applied stress, the relationship between strain rate and stress can be derived. For material transport by slip in a pure metal it is expected that the strain rate should be proportional to the 4.5th power of the stress, while for transport by diffusional flow the strain rate should be proportional to the first power of the stress. It was found that for stresses above 5 X 10&apos; dynes per sq cm (70 psi) the exponent in the stvain rate-stress relationship was near 4.5 while for a stress of 0.7 x 108 dynes per sq cm (10 psi) an exponent of 1.8 was determined indicating that under the conditions of the experiment the transport mechanism for creep in silver shifts from diffusional flow to slip in the region of 0.7 x 106 dynes per sq cm (10 psi). The importance of these findings for an understanding of the mechanism of material transport in conventional sintering (without an external load) is pointed out. In a paper by Early et al.,&apos; the mechanism of material transport during sintering was treated as a creep process. It was pointed out that two possible types of creep processes could operate during sintering. In one of these material transport takes place by vacancy diffusion due to a gradient in chemical potential. The vacancy sinks in this diffusion process may be either grain boundaries, as in Herring-Nabarro microcreep, or regions beneath a surface of small radius of curvature. The diffusive transport may be by volume, grain boundary, or surface diffusion. In the other mechanism, creep takes place by slip, i.e., dislocation motion with its rate generally controlled by dislocation climb. One method of distinguishing between these two modes of materials transport is to investigate the creep rate, i.e., the rate of material transport as a function of stress. For material transport by vacancy diffusion, the creep rate is proportional to the stress.&apos; For creep by slip, controlled by dislocation climb, the creep rate is proportional to a power of the stress. For pure metals this power is 4.5;374 for dispersion-strengthened material a fourth power relationship can be derived.5 This difference in the stress dependence of the strain rate is illustrated schematically in Fig. 1, in which the logarithm of strain rate is plotted vs the logarithm of the stress for material transport in a pure metal by vacancy flux and by dislocation climb. The stress range in which the two types of material transport overlap would be expected to be relatively small. In the experiments described by Early and coworkers compacts pressed from spherical copper powder were sintered isothermally under an applied axial external load in a dilatometer. During sintering the magnitude of the applied external load was changed and the rate of axial shrinkage of the compact before and after the load change determined. In order to find the stress dependence of the strain rate the magnitude of the stress acting on the neck area between the spherical particles before and after the load change had to be calculated. This stress is the resultant of the stresses due both to the externally applied load and to surface tension. The stress due to surface tension is a function of the geometry of the powder compacts. It can be determined, at least for simple geometrical arrays,&apos;76 but it is difficult to obtain exact values of the stresses for actual powder compacts when only average values of particle diameters and of neck diameters between particles are available. The calculation of stress and the strain rate-stress relationship is therefore only very approximate.

Jan 1, 1968

By Larry Kaufman, Joan B. Berkowitz-Mattuck, Edward V. Claugherty

The oxidation characteristics of hafnium and zirconiunr diboride were measured between 1200 and 2200&apos;K by a thermal- conductivity method which continuously ttzeasures the rate of reaction of oxygen with the diboride and by a metallographic air oridation method zuhich provides a measure of the total arr7ount of bovide conuerted to oxide for a given time interval. The oxidized specimens obtained from the tl~eritaal-coi~ductitrity method were also examined by quantitatire metal-lographic procedares. The significant results obtained in this investigation reveal that metal-rich compositions of lzafi~iutil diboride proride the most oxidation-resistatzt material up to 2000°K; hafnium diboride is tmre oxidation- resistant titan zirconium diboride at all tempevatures examined; the morphology of the oxide formed on H/B2 and Llie temperature coefficient qi. the oxidation rate constants change at the temperature of ttze monoclinic to tetragonal phase transition] in HfO2; the oxidation of neither HfB2 nor ZrB2, results iN catastrophic Jazlure at lorc. oxygen pressures; and pvefevetztial gvaLti boundary oxidation was not obsevued for either HfBi or ZYB, A comprehensive study of the high-temperature characteristics of refractory transition-metal di-borides is currently in progress. This program has included investigation of the physical, thermal, and thermodynamic properties of TiB2, ZrB2, HfB2, NbB2, and TaB2. In addition, aspects of the synthesis and fabrication of such materials have been studied. In view of the diverse nature of this research, a number of other laboratories have actively participated and contributed specific capabilities for analysis and characterization of these materials. As a consequence, an extensive description of the relevant properties of these compounds has emerged which is central in evaluating their high-temperature (1200" to 2500°K) performance. To date, information on thermodynamic stability, specific heat, and vaporization characteristics,1 hot hardness and electrical resistivity,1, 3 therma1 expansion:&apos;4 and thermal conductivity 1, 5 has been presented. This information has been generated on materials of the highest purity (98.5 to 99.9 wt pct Me + B) and density currently available. Samples fabricated by zone melting6 and high-pressure hot pressing"3&apos;7 techniques have been used to generate suitable specimens for all of the aforementioned studies. dation characteristics of the most oxidation-resistant of these materials, hafnium and zirconium diboride, is presented and a description of the synthesis and the experimental procedures used to prepare and characterize specimens is given. The high-temperature range under consideration (1200" to 2200°K) and the known dependence of oxidation characteristics on sample chemistry, density, and oxidation conditions required a close coupling of the synthesis, fabrication, and evaluation procedures.8 This was accomplished by continual surveillance of chemical composition of starting materials before and after specimen fabrication and by evaluation of density, phase constitution, and microstructural features prior to and after oxidation exposure. I) PROCUREMENT AND CHARACTERIZATION OF STARTING MATERIALS In view of the current state of the art in fabricating refractory boride materials, the methods used in preparing samples for the present study are given in detail as follows: starting materials were purchased in high-purity powder form and fabricated by high-pressure hot pressing into 0.40 by 1.00 in. bars from which oxidation specimens were obtained. The hafnium diboride used in this study was purchased from Wah Chang Corp.; the zirconium diboride from U.S. Borax and Chemical Co. These powders were routinely characterized by quantitative chemical analyses for metal, boron, carbon, oxygen, nitrogen, and iron, by qualitative emission-spectrographic analysis for trace impurities, by X-ray procedures for extraneous phase identification, and by powder densitometry for comparison with X-ray (theoretical) density. Hafnium and zirconium metal and elemental boron were also purchased as high-purity powders and characterized for impurities by emission-spectrographic analyses. The hafnium diboride was procured in three shipments which were designated as HfBl.g7(1), HfB1.88(2A), and HfB2.12(2). The indicated stoichiometry is based on the atomic ratio of total boron to total hafnium; the number in parentheses identifies the shipment number. Shipment 1 was 5 1b, shipment 2A, 1 1b, and shipment 2, 8 1b. The zirconium diboride was procured as a 20-1b shipment and designated as ZrB1.89(1). A small quantity of purified zirconium diboride was also supplied and designated ZrB1, 9(P). The averaged results for chemical analyses which were generally performed according to the procedures set forth in the compilation by KrieGe9 are presented in Table I. Qualitative spectrographic analyses indicated that Ca, Cr, Ti, Si, Zr (in H~B~), and A1 were present at levels between 0.01 and 0.10 wt pct. Other metallic elements were found to be less than 0.01 wt pct. Since it is virtually impossible to purchase these materials in the desired quantities (5 to 20 lb) as single-phase compounds it is necessary to obtain

Jan 1, 1968

By I. R. Kramer

The stress associated with the surface layer was deter-minedfor iron and molybdcnum. These measurements show that the surface layer plays a very important role in the plastic deformation of bcc metals. From the relationship between strain arui the surface layer stress and the back stress imposed by the internal obstacles, it is Proposed that the term of of the Hall-Petch equation is related only THE large number of studies on plastic flow of bcc metals have failed to delineate clearly the nature of the obstacles influencing dislocation movement. Since there have been numerous articles which have reviewed the mechanisms proposed to explain various aspects of plastic flow of bcc metals (for example, conrad1 and Keh and weissmann2), no attempt will be made here to discuss their conclusions. However, it should be to the surface stress while the term Kfd is the stress required to move dislocations through internal obstacles. Further, it was possible to derive the well-known empirical equation C = , and it is shown that the terprz a is comprised of two terms: a term Cs associated with the surface layer stress and Ci related to internal dislocation obstacles. pointed out that in all of these publications it has been tacitly assumed that plastic flow occurred uniformly throughout the cross section of the specimen (except, of course, for the Luders band region), It has generally been further assumed that the nature and distribution of the obstacles opposing the motion of dislocations are of the same type throughout. The role of the surface layer in bcc metals has not, as yet, been considered even though its role in fcc metals has been shown clearly to be of primary importance.3-5 Investigations on fcc metals have shown that at the onset of plastic deformation a surface layer acts to

Jan 1, 1968

By J. H. Bucher

The effect oj strengthening by the precipitation of CbC on the toughness offerrilic steel (0.11 pct C, 0.74 pct Mn, 0.02 pct Cb) was stlrdierl. A g-veater degree of pvecij~itatiotz strengthning is ohtaznable by aging the so11ution-treated steel rtt velaticely Low temperature was (1050" to 1l0O3F) than at higher temperatures (I150" to 1200°F). Tlzis ),eslrlts because a finer dispersion of CbC is obtained. However, the ductile-brittle transition te~ripertltlo.e is raised approximately tlzjeee tirues as IFLIIC~ perv a given yield stress increase) by aging in the lower tempareture range, as opposer1 to the Irigher temperaturee range. This is proposed to result pl-iir~avily because different precipitation distributions ar-e associated with the different aging -treaments. For tile steel aged at 1050" and 1100 F, heavy: precipitation on disloccctions occurs resulting in a high degree of rlislocution pinningh tl~us lowevirzg- the toilguiess. in tll c(.rise of the steel aged at 1150 &apos;a?zd 1300"F, Pvecipitatiol~ is irzove g-enerally distributed throughout the matrix rather limn preferentially on dislocations.; and Ike detlsitj, oj mobile dislocations is not sttyong-ly affected. THE most effective means of strengthening hot-rolled ferritic steels are grain refinement and precipitation hardening. It is well-known that grain refinement increases both the strength and the toughness. However, the upper limit of yield strength obtainable by grain refinement alone is about 60,000 psi,&apos;12 i.e., without the formation of acicular structures of lower toughness. The most economical means of obtaining higher strengths is the addition of small amounts of elements such as columbium, vanadium, and titanium which form carbonitrides. These precipitates are very effective in increasing the strength, but the ductile-brittle transition temperature is is Although it is agreed upon that precipitation decreases the toughness, there has been little effort to isolate the causative aspects of embrittlement. The purposes of this research are to assess quantitatively the embrittlement associated with different precipitate distributions and morphologies, and to correlate the degree of the embrittlement with the microstruc-ture. In this manner, a more fundamental insight into the mechanism of embrittlement may be gained. The steel selected for this purpose was a low-carbon, columbium-bearing steel which, after solution treatment, was aged in various time-temperature cycles while attempting to hold all other microstructural variables constant. MATERIALS AND PROCEDURE The composition of the steel used in this investigation is given in Table I. Starting with commercially hot-rolled material 0.074 and 0.280 in. thick, tensile blanks were cut from the thinner material and sub-size Charpy blanks from the thicker stock. The steel had been rolled using a low finishing temperature and rapid cooling after hot rolling to produce grain refinement. Also, chemical-extraction analysis and electron-microscopic examination revealed that the cooling was rapid enough to suppress any significant precipitation of CbC. Therefore, the as-rolled steel is essentially in the solution-treated condition in terms of the precipitation reaction, and it was used as the base material in this investigation. The room-temperature tensile properties of the as-rolled steel aged at 1100", 1150", and 1200°F for various time intervals were determined, see Table 11. The ductile-brittle transition behavior was studied for smooth and notched tensile specimens, and longitudinal half-size (0.200 in. thick) Charpy V notch specimens aged at 1100" and 1150°F for various times, also Table 11. In addition, tensile specimens aged for single times at 1050" and 1200°F were evaluated with respect to their ductile-brittle transition. Only one 1050°F treatment was used because of the extremely slow kinetics of precipitation at this temperature and because no significant difference between 1100°F aging was observed. Only one treatment at 1200°F was used because aging at this temperature alters the shape of the grain boundary cementite particles present in this steel and complicates analysis. The design of the tensile specimens is given in Fig.1.

Jan 1, 1968

By P. Chiotti, R. J. Hecht

Vapov pressure data of solid and liquid alloys along with phase diagram data WERE employed in calculating the thermodynamic properties for the Ca-Zn system. The relative excess partial molal free energy for zinc in the liquid alloys was fitted to the relation wnvre i is we tvmpvrutiirc in ucgrees neivtu ana in is the tr~ole fvaction. Relations for the standard free energy of fortnation for the eight cornpounds were developed. ThE phase diagram for the Ca-Zn system as determined by Messing et a1.&apos; is reproduced in Fig. 1. Earlier work on the phase diagram has been summarized by Hansen and Anderko2 and the very limited thermodynamic information for this system has been reviewed by Hultgren et a1.3 The heat of fusion for the compound CaZn2 has recently been determined to be 9260 cal per mole.4 Solubility data for calcium in zinc for the temperature range of 650" to 725°C were fitted by the relation and reported by Johnson and Dilloon5 In the present investigation zinc vapor pressure data were obtained by both the dewpoint and Knudsen methods in order to calculate the thermodynamic properties for liquid and solid alloys. The dewpoint method as employed in this work requires that only one of the components be significantly volatile, in this case zinc, and that the vapor pressure of pure zinc be known. For calculation purposes the vapor pressure data for pure zinc3 were fitted to the relations; for temperatures below 692.7K, the melting point of zinc. The corresponding free-energy relations for vaporization and sublimation of pure zinc are respectively. Some of the calculations also required knowledge of the free energy of fusion of zinc and calcium and the free energy for the 721°K allotropic transformation of calcium. The free energy of fusion of zinc was taken to be the difference between Eqs. [4] and [5], or The free energy of fusion and transition for calcium were based on the enthalpy values of 2040 and 222 cal per g-atom, respectively, as reported by Chiotti et al. 4 and the heat capacity data given by Hultgren et a1.3 and are represented by the relations, In the latter expression the difference in heat capacity between the a and 0 forms of calcium was neglected. MATERIALS AND EXPERIMENTAL PROCEDURES The metals used in preparing the alloys were Bunker Hill slab zinc, 99.99 pct pure, and Ames Laboratory calcium. The average impurities in the calcium in parts per million by weight, ppm, were N-17, H—77, Ba—216, Mg-15, Si—7, C—157, and Fe—9. The zinc was cleaned in dilute nitric acid and the calcium was used as received. The calcium was stored and handled in an argon-filled glove box and when removed from the dry box it was kept in an argon-filled desiccator. Tantalum was found to be a suitable alloy container. The alloy components were contained in a 1-in.-diam cylindri-

Jan 1, 1968

Jan 1, 1967

Jan 1, 1968

By Y. C. Liu, G. A. Alers

The anisotropy of Young&apos;s modules has been measured in a series of Cu-Zn alloys containing from 1 to 28 wt pct Zn and cold-rolled to 97.5 pct reduction in thickness. These modulus data show a transition in texture as a function of zinc content. In alloys containing 3 pct Zn or less, the texture is essentially that of copper. A transition is clearly seen to develop in 5 and 8 pct Zn alloys, and to have been completed in a 12 pct Zn alloy. Increasing the proportion of zinc to 28 pct does not produce any further changes in texture. The ease and clarity with which this texture transition can be seen in the modulus data is in great contrast with the inccertainty in the visual detertiination of the range of texture transition from a series of pole figures. It is also shown that the rolling texture of the high zinc content alloys cannot be a one-COIIIponenl texture, characterized by the (110)(112) ideal orientations, as usually assumed. A second, quite important component must also be present. Furthermore, the present data are inconsistent with matrix orientations on which the "oriented growth" hypothesis of the Formation of recrystallization textures is based. THE rolling textures of binary fcc alloy systems have generally been considered to lie between two terminal textures. One of these terminal textures is the copper type and the other is the 70:30 brass type. Alloys of intermediate solute content exhibit a texture which is said to be a transition between the two terminal textures. On the basis of efforts to correlate rolling and recrystallization textures, two concepts have been advanced in the literature on the nature of these transition textures. One school1 suggests that the addition of alloying elements, e.g., zinc, causes a gradual rotation about one of the (111) poles from the (358)( 523) ideal orientation of the copper-type rolling texture to the terminal {1 10)( 112) ideal orientation used to describe the rolling texture of 70:30 brass. This concept is depicted in Fig. 1 with the angle of rotation corresponding approximately to the zinc content and is designated here as the "rotational concept", since the ideal orientation of the transition textures varies continuously with alloy content. The other concept,&apos; the "mixture concept", suggests that the copper-type rolling texture is a two-component texture. {358) (523) and {110)(112). The addition of alloying elements affects only quantitatively the relative amounts of these two components—a decrease in the amount of material in the {385}(523) orientation and an increase of that in the {110)(112) orientation with increasing solute content. An experimental distinction between these two concepts is essential to a better understanding of both the mechanism of formation of rolling textures and the subsequent formation of recrystailization textures in fcc metals. Because of the lack of resolution inherent in pole figures, it has not been possible to differentiate between the two concepts on the basis of this data. Both concepts, in fact, make use of the same set of pole figures but interpret them in a manner acco-rding to the view held.&apos;" Experimental results obtained from inverse pole figures or selected-area electron diffraction3j4 simply -yield more refined pole figures but do not eliminate the ambiguity of interpretation. Quantitative measurement of an entirely different, orientation-dependent property must thus be considered

Jan 1, 1967

By C. M. Sellars, F. Maak

Electvomotie -force measurements hazle been made on ten Au-Ni alloys at temperatures 7754 825O, 900O, and 935°C using galvanic cells with solid electrolyte. Partial and ivtegral thermodynamic functions are derived from the determined activities. which show la9,ge positive deviations jvom Raoult&apos;s law. The enthalpies of mixing- are in close accord with calorinletric data and are of comparable accuracy. The results are used to determine the miscibility gap and the spinoidal curL1e, and are qlalitaticely interpreted in terms of the large size difference and electvonic interactions between gold and nickel. Tabulated nalues of the actielities and partial and inteqal free energies, entropies , and enthalpies of mixing are given in the appendix. The Au-Ni binary alloys are of particular interest thermodynamically as despite the large size difference between gold and nickel atoms they have a simple phase diagram.&apos; This shows complete solid solubility above a miscibility gap with a maximum at about 71 at. pct Ni and 810°C. Activity measurements (over a range of temperature) have been made previously on this system using an electrolytic-cell method with a liquid chloride eletrolyte. The accuracy of this data has, however, been queried3&apos;4 as it leads to values for the enthalpy of mixing which are too high when compared with calorimetric measurements.46 Also, the entropy of mixing is too high when compared with the values derived from specific-heat measrements. It was therefore considered worthwhile redetermining the activity data using a galvanic cell with a solid electrolyte (0.85 ZrO, 0.15 CaO). This method is based on a technique introduced by Kiukkola and aner&apos;&apos; and has been used to determine the thermodynamic data for Cu-Ni alloys.&apos; THE ELECTROLYTIC CELL The cell used to determine the activity of nickel in Au-Ni alloys was: The cell reaction, which is discussed for a similar cell by Rapp and Maak, gives a direct measurement of the activity of nickel in the alloys as where aNi(alloy) is the activity of nickel in the alloy, 3 is Faraday&apos;s constant, E is the measured electromotive force, R is the gas constant, and T is absolute temperature. For the valid application of this type of cell the arrangement must be ther modynamic ally stable, which means that: 1) The free energies of formation of the oxides fulfill the relation 2) The vapor pressures or dissociation pressures of each component must be so small that during the time of the experiment the amount of any material transported by the gas is negligible. 3) Side reactions which change the alloy composition must be negligible. These reactions are possible both with the surrounding gas atmosphere and with neighboring solid material. The first two conditions are fulfilled by the above cell.&apos; The third condition is approached by care in the experimental technique. The factors which control the attainment of equilibrium in this type of cell have been considered by Rapp and aak, who conclude that metallic diffusion in the alloy is the rate-control ling step. A lower temperature limit for the practical application of this cell is therefore set by the condition that equilibrium must be attained in reasonable times. EXPERIMENTAL MATERIALS AND TECHNIQUE The starting materials were mint gold (99.99 pct Au) and Mond nickel (99.9 pct Ni). Alloys, nominally 5 at. pct Ni and then each 10 at. pct across the system, were produced by vacuum-levitation melting compacts of mixed drillings of the pure metals. The nickel used for the reference tablets in the cells was melted in a similar manner. The alloy ingots, cast in boron nitride molds, were cold-forged and

Jan 1, 1967

By H. Morgolin, M. A. Greenfield, P. A. Farrar

In attempts to correlate mechanical properties with microstructure, it is desirable to determine the effect of ß grain size. Determination of 6 grain size is not a problem where a has precipitated at ß grain boundaries. However, in alloys which have been processed to produce an equiaxed a +ß structure the ß grain size is difficult to determine because ß grain boundaries do not etch. The ß grain size can sometimes be laboriously determined by searching for grains which have some a! outlining the 6 grains. This difficulty of ß grain size determination in equiaxed a + ß alloys can be overcome by heating below the ß transus temperature for a short time and water quenching. Fig. 1 shows the starting equiaxed a + ß structure of the Ti-5.25A1-5.5V-0.9Fe-0.5Cu alloy used in this investigation, and Fig. 2 shows the microstructure of the same sample water-quenched after heating 1 min at 1675"F, which is 25°F below the 0 transus. The heavily etched microstructure consists of some equiaxed a! in a structure showing considerable amounts of transformed ß. The ß grain

Jan 1, 1969

By H. I. Aaronson, W. L. Winterbottom, G. M. Pound

On the basis of a statistical thermodynamic treatment of the data of smithz2 on the activity of carbon in austenite, a Darken and smith23 deduced that the interaction energy, wy, between carbon atoms occupying nearest-neighboring interstitial sites is repulsive. Aaronson, Domian, and pound 24 (ADP) have applied two different statistical thermodynamic analyses to the same data. One analysis, a composite of the treatments of Darken and Smith,23 Speiser and Spret-nak," and Kaufman, Radcliffe, and Cohen,26 primarily takes account of the positional entropic effects of wy in dilute solutions. The other, due to Lacher 27 and Fowler and Guggenheim," is also applicable to more concentrated solid solutions. The equations which both analyses produce for a,, can be rearranged so that, when In a,, is plotted against the appropriate function of the carbon content of the austenite, the plot will have a slope of unity when the correct value of wy is employed. Since no other constant need be simultaneously evaluated in order to make this determination, ascertaining the most probable value of wy from Smith&apos;s data on the bases of these analyses was an elementary problem in least-squares statistics. ADP demonstrated that the wy&apos;s obtained from both analyses vary with temperature in a statistically significant manner. McLellan, Gerrard, Horowitz, and sprague 29 (MGHS), on the other hand, have recently made a new statistical thermodynamic study of a,, and have concluded from this study that wy is independent of temperature. The equation which they derived for a,, is: where 0 = ratio of the mole fraction of carbon atoms to that of iron atoms, z = number of sites from which occupancy is excluded by the presence of a carbon atom and the repulsive character of its interaction with nearest-neighboring carbon atoms (z = zb in Ref. 24), and x = a geometric parameter taking account of overlapping of the exclusion shells of nearby carbon atoms. z is 1 when wy = 0 and 13 when wy = Q. For a given value of z there is a definite value of X . In the absence of an equation connecting z and X, MGHS computed geometrically the value of x at several integral values of z; x increases from 0.47 to 8.67 as z is increased from 2 to 5. MGHS fitted Eq. [I] to the a,, data of Ellis, Davidson, and Bodsworth30 at 925° and 1050°C and to the data of smithz2 at 1000°C. Confining z to integral values in the range 4 to 6, and also somewhat restricting AG, the values of z, X, and AG were sought which would minimize the average root mean square percent deviation, D, of the calculated a,, values from those obtained experimentally. Finding that D was almost always smallest when z = 5, they concluded that z is independent of temperature. Since:23, 26 where R = gas constant and T = absolute temperature, and w,, was considered to be similarly invariant, this statement cannot be strictly true. The temperature range employed, however, was small and the interval between the values of z tested was sufficiently large so that the results could still be consistent with a constant The statistical treatment of the a,, data applied by MGHS is clearly quite approximate. Since more than three values of a,, are available at each temperature, z, and X, and AG have actually been over deter mined, rather than underdetermined as their treatment implies. In order to make more complete and exact use of the a,, data, we shall employ the Legendre Method of Least Squares, as described by Whittaker and Robinson. This technique permits conversion of the overdetermined system of equations produced by the substitution into Eq. [I] of the available pairs of a,, and 0 data at a given temperature into an even determined system. Solution of the latter system yields the most probable value of each of the three constants. This procedure will allow a statistically rigorous test to be made of the conclusion drawn by MGHS in respect of the temperature dependence of In order to obtain linear iiequations of conditions" from Eq. [I], the exact values of the three unknown constants are replaced by G/RT = =/RT + y, where z, i, and aG are estimated values of the constants. This allows the conditional equations to be written in the form:&apos;l where fi -f6, i, =/RT) = the right-hand side of Eq. [I] for a given pair of ayi and values. When applied to Eq. [I], these relationships become: The i equations of condition are reduced to three "normal equations" in the manner discussed by Whittaker and Robinson. The normal equations are then solved for a, ß, and y. Both of these operations are most conveniently performed by means of the computer technique of Winterbottom and Gjostein.32 This technique was used iteratively until the value of 0 (the most sensitive of the three correction factors) became 10.01. (Until they became small, a, 8, and y

Jan 1, 1969

By D. Webster

It has been found that cold work and annealing treatments greatly retard subsequent grain growth and re-crystallization in thoria-dispersed nickel and NiCr. It is suggested that this is due to cavities which form at the particle-matrix interface durilzg cold work. Some of&apos; the cavities formed are stable at elevated temperatures and prevent the mouement of large-angle grain boundaries. The movement of these boundaries at elevated temperatures is known to start recrystal-lization and when they are pinned by voids recrystal-lization is made more difficult. TD Nichro,me had a maximun resistance to recrystallization when the initial structure was coarse-grained and had a low dislocation density. A precipitation reaction occurred between 700° and 1200° F in cold-worked TD Nichrome. It has been shown by Grant and his coworkers1-3 that a large part of the strength of dispersion-hardened materials is a result of the cold-worked structure produced during fabrication. In a comparison of the strength of cold-worked copper with and without an oxide dispersion, Grant4 has demonstrated that there is little difference in strength at room temperature. The function of the dispersion is essentially to maintain the cold-worked structure at higher temperatures than is possible in the base material. The dispersion accomplishes this by making recrystallization more difficult. The mechanism by which this is achieved is discussed in this paper. The fabrication history of the material has been found to greatly influence recrystallization behavior and this aspect is discussed in some detail. I) MATERIAL Three different thoria-dispersed materials were examined. 1) DuPont 2 pet thoria-dispersed nickel (TDNi) in 0.050-in. seeet. Thoria particle size is in the range 100 to 2500A. 2) DuPont 2 pct thoria-dispersed nichrome (20 pet Cr, 80 pct Ni) (TDNiC) in 0.125-, 0.060-, and 0.040-in. sheet. In the as-received condition these sheets have

Jan 1, 1969

By R. E. Fryxell, E. C. Duderstadt, P. K. Conn

Permeation rates for nitrogen through are-cast tungsten were measured in the temperature range 1800° to 2600°C at nitrogen pressure differentials of 1.0 and 0.1 atm. Gas chromatography was used to measure the quantities of nitrogen which permeated through 0.26-and 0.53-mm-thick membranes. Permeation rates were proportional to the difference between the square roots of the nitrogen pressure on each side of the membrane and inversely proportional to membrane thickness, indicating that atomic nitrogen was the difjusing species and that diffusion was the rate-limiting step of the process. Permeation coefficients calculated from the rates may be represented by the equation: where P is expressed in the units cu cm (stP)-mm per 1/2 sq cm-min-atm"&apos; and 70.0 * 0.8 is the activation energy in kcal per mole of permeating species. In a previous paper&apos; permeation rates were reported for hydrogen and for oxygen through tungsten. The permeation cells used for the hydrogen measurements were suitable for measuring nitrogen permeation under similar conditions, so the study was extended to include nitrogen. Permeation rates of nitrogen through the polycrys-talline arc-cast tungsten membranes were measured in the temperature range 1800" to 2600°C using two membrane thicknesses and two nitrogen pressure differentials. Permeation coefficients were calculated and the activation energy of the process determined. THEORY For diatomic gases which dissociate prior to diffusion, the quantity of gas which permeates through a membrane or wall at an elevated temperature where bulk diffusion is the rate-limiting step of the process is given by the equation: where: q = total amount of gas permeating a membrane, P = temperature-dependent permeation coefficient, A = area of membrane, t = time, p1 = gas pressure of high-pressure side, p2 = gas pressure of low-pressure side, and d = thickness of membrane. When bulk diffusion is the rate-controlling step, and the solubility of the gas obeys Henry&apos;s law, a further relationship is that, for any given temperature and pressure, the permeation is the product of diffusivity and the amount in solution. The permeation coefficient, P, is temperature-dependent according to the equation: where: Po = permeation constant, Q = activation energy, R = gas constant, and T = absolute temperature. EXPERIMENTAL PROCEDURES Cell Design and Experimental Apparatus. Two of the permeation cells used for the previous measurements of hydrogen permeation are shown in Fig. 1. Each consisted of a cylindrical chamber divided into two parts by a permeation membrane fabricated from high-purity (>99.9 pct W) wrought tungsten sheet made from vacuum arc-cast tungsten. The membrane in one cell was 0.26 mm thick and the other was 0.53 mm. The fabrication of these cells was described previously.&apos; The permeation cells were mounted vertically in a high-temperature tungsten resistance furnace and the membranes and the center portions heated to test temperatures. Temperatures were measured by sighting an optical pyrometer on the edge of the membrane through a small hole in a tungsten plate which was fastened to the side of the cell. This arrangement provided an approximate black-body condition and eliminated the need for emittance corrections. Appropriate corrections were made for the quartz window through which the temperatures were measured. Argon gas flowed through the 10-liter volume of the furnace at the rate of 4.7 liters per min. An impurity

Jan 1, 1969

By S. A. Levy, J. D. Wood, J. F. Libsch

A study of the preparation and characteristics of a sevies of Fe-C-N alloys has been conducted. X-ray, microhardness, and metallographic data from a series of single-phase alloys produced by controlled nitriding of carburized foils are reported. The existence of a, y&apos;, and E carbonitride has been confirmed and it has been found that y&apos; is the hardest of these three phases. These results have been used to interpret the structures observed in nitrided steels. ThIS paper represents the second phase of a three-part study concerning combined nitriding and induction hardening. The first part&apos; reports the engineering properties (depth hardness, fatigue, corrosion, and tempering resistance) obtained when seven different steels were subjected to the combined surface hardening treatment. The present paper concerns the nature of a series of single-phase alloys prepared by controlled nitriding of carburized foils. The objective was to simulate the various layers present in the as-nitrided steels. A third paper will report data concerning induction hardening of the foils and will attempt to set forth a mechanism for the high hardness obtained with the combined treatment. Studies of nitriding generally fall into three categories. The first were performed with bars using conventional nitriding procedures, thus yielding a nitrogen gradient and a series of phase layers near the surface.2-6 Such studies were complicated by the presence of the nitrogen gradient, for it was difficult to correlate microhardness and X-ray data with chemical composition. There were always complications due to penetration of the X-ray beam which yielded data influenced by underlying layers. Similar difficulty was experienced in taking hardness readings on a tapered specimen or in isolating various layers by machining. A second group of investigations was primarily concerned with establishment of regions in the Fe-N7,9,10 or the Fe-C-N phase diagram.8 These were performed with powder specimens which were nitrided in am-monia-hydrogen mixtures to produce a series of single-phase specimens. Unfortunately, the powder specimens were not amenable to either microstruc-ture or microhardness study. The third was essentially a compromise of the first two. Bose and Hawkes," using 0.005-in.-thick iron foils, were able to obtain microhardness, microstruc-ture, composition, and X-ray data on Fe-N alloys of eutectoid composition by equilibration in ammonia-hydrogen mixtures at 1290°F. To obtain a better understanding of the nitriding process and the characteristics of the phases in ni-trided plain carbon steels, the present investigation was undertaken. The investigation was patterned after that of Bose and Hawkes with the exception that nitriding was performed at 930°F (below the eutectoid temperature), as in the case of conventional nitriding. In addition to gaining some understanding of the source of high hardness of nitrided and induction-hardened steels, found in a previous study,&apos; this investigation was concerned with the possibility of controlling the phases present in a nitrided case through the use of ammonia-hydrogen mixtures. Knowledge of the properties of each phase and the conditions necessary for its production could make it possible to avoid undesirable surface layers in nitriding. EXPERIMENTAL PROCEDURE Specimen Preparation. A specimen thickness of 0.005 in. was selected to insure that a uniform nitrogen concentration could be obtained in a reasonable length of time by equilibration in an ammonia-hydro-gen atmosphere. The starting material was a vacuum-cast ingot of pure iron, the major impurities of which were: C, 0.003 pct; 0, 0.014 pct; S, 0.004 pct; N, 0.003 pct.* &apos;Compositions are quoted as weight percent except as noted. The ingot was heated to 1500°F in a nitrogen atmosphere and forged to a slab 0.25 in. thick. The material was then reduced to 0.005 in. by alternately cold rolling and annealing in a nitrogen atmosphere at 1600°F. This material was divided into five groups of specimens which, after carburizing, yielded the following carbon contents: 0.003, 0.28, 0.36, 0.46, and 0.74 pct C. The carburizing was performed by the Leeds and Northrup Research Laboratories, using a controlled "carburizing potential" which yielded a uniform carbon concentration throughout the cross section.

Jan 1, 1970

By Alan Arias

On grinding in pure water, zirconium, tantalum, iron, and stainless-steel powders were extensively comminuted and simultaneously oxidized with hydrogen release, whereas nickel, copper, and silver powders did not react with water and their particle sizes increased. On grinding nickel, copper, and silver in water pressurized with oxygen, nickel and copper became extensively comminuted and were oxidized, whereas silver did not react with oxygen and its particle size increased. From these results and other considerations , it is hypothesized that for extensive comminution of ductile metals and alloys to occur on grinding they must react with the grinding media. UlTRAFINE metal and alloy powders are finding an ever-growing number of applications in metallurgy and in other fields.&apos; Of particular interest are ultrafine metal and alloy powders suitable for dispersion strengthening.2&apos;9 Various research programs on dispersion strengthening are being carried out and in some of these programs the ball-milling method is being used to produce dispersion-strengthened materials. This method usually involves the simultaneous grinding of metal or alloy and a dispersoid followed by consolidation of the resulting powder mixture. To obtain the ultrafine powders required for dispersion strengthening,&apos; grinding is carried out in many liquids, including aqueous and nonaqueous media, with or without grinding aids.4&apos;5 Nonaqueous liquids usually contain water as an impurity and some grinding aids may contain water of hydration.5 The water present may affect the grinding process. The writer has shown5 that. on ball milling chromium in water, the chromium is oxidized and hydrogen is released. It was surmised that the same reaction may occur on ball milling other metals and alloys in waterbearing liquids. Therefore, the investigation of ball milling in water was extended to metals and alloys other than chromium. In the course of the investigation, however, it became apparent that the data-to-gether with the results from a few additional experi-ments—could be used to postulate a comminution mechanism for ductile metals and alloys. A well-known comminution theory is that of smekal.7 According to this theory, comminution is possible because of the weakening effects of surface cracks and other imperfections in materials. This theory imposes a lower limit of about 1 µm for the ground particles. The beneficial effects of liquids and additives on the rate of grinding are well known.8 Mechanisms by which liquids and additives may aid in grinding were reviewed by Rose and Sullivan.&apos; One aspect of these effects is based on Rehbinder&apos;s theory of crack propagation in materials under stress.9 According to Reh-binder&apos;s theory, liquids or additives may promote the spread of cracks in stressed materials by lowering the surface tension at the crack tip. Rose and Sullivan surmise that the same mechanism may be operative during grinding, thereby facilitating comminution of the particles. In addition, Rose and Sullivan reviewed how additives may act as dispersants as a result of their being adsorbed on the surface of the particles being ground. This concept has been suggested by Quatinetz, Schafer, and smea15 to explain from their experiments the major role of additives that enabled them to grind metal down to 0.1 µm. Discussions of other comminution theories and additional sources of material on the subject will be found in Ref. 10. None of these previous suggestions and theories, however, can account for all phenomena encountered during ball milling of metals to submicron size in this and in a previous investigation by the author.6 The objectives of this investigation were to determine the behavior of metal powders during ball milling either in pure water or in oxygenated water and to gain an insight into the grinding mechanism. Zirconium, tantalum, iron, nickel, copper, and silver powders were ball-milled in pure water. These metals were selected because their oxides cover a wide range of free energies of formation. For comparison purposes, an alloy-type 430 stainless steel-was also ball-milled in pure water. The pressure of the hydrogen released during ball milling was monitored in order to determine the oxygen that combined with the metal or alloy. In order to obtain more information on the nature of the grinding process, nickel, copper, and silver powders were also ball-milled in oxygenated water (water pressurized with oxygen). The oxygen that reacted with the powders was determined from the pressure decrease in the mills. The powders resulting from ball milling in pure water and in oxygenated water were subjected to surface area, optical microscopy, and X-ray diffraction analyses. With these data, the oxygen calculated to be combined with the metals during ball milling, and comparison of the free energies of formation of the oxides of the milled powders with that of water, a comminution mechanism was postulated. MATERIALS, EQUIPMENT, AND PROCEDURES The materials used in this investigation were powdered metals, deaerated distilled water, high-purity helium, and commercial grade (99.5 pct purity) oxygen. The powdered metals used were zirconium, tantalum, iron, nickel, copper, and silver. A 16 pct Cr, ferritic stainless steel, type 430, was also used. The purities (or nominal compositions) and the surface areas of these metals and the alloy are given in Table I.

Jan 1, 1970

By Å. Sterten, R. Tunold, J. Brun, K. Dalatun

c axis compression behavior of beryllium single crystals at three purity levels under hydrostatic pressures up to 27 kbars was determined. Extensive non-basal slip, observed by two-surface trace analysis and transmission electron microscopy, occurred under a hydrostatic pressure of about 12 kbars (175 ksi) for the high-purity (twelve-zone pass) material and at about 19 kbar (275 ksi) for the lower-purity (zone-leveled) material. Prismatic loops with a (c + a) Burgers vector were observed in association with second-phase parti- A principal factor limiting the use of beryllium is its brittle behavior when tested at the usual strain rates (E = 10-4 sec- 1) at temperatures below about 200°C and under impact conditions at temperatures in excess of 200°C. It has been proposed&apos; that the brittle-ness of beryllium is associated with the lack of a sufficient number of independent slip modes and the absence of a slip mode with a Burgers vector out of the basal plane [presumably (c + a) pyramidal slip mode] and to the ease with which beryllium cleaves on the basal and second-order prism planes. The absence of pyramidal slip has been attributed to a high Peierls-Nabarro stress associated with the motion of dislocations with a (c + a) vector and the ease with which cleavage occurs on the basal and second-order prism planes. The experimental evidence in support of the proposed explanation for the brittleness of beryllium is far from complete; for example, that the ductile-to-brittle transition in polycrystalline beryllium is associated with the operation of profuse (C + a) slip has not been unequivocally established. The occurrence of (c + a) slip ({1122}(1123)) has been experimentally established2-5 under conditions where basal and prism cleavage are restricted in a c axis compression test. In these investigations (C +a) slip was found in high-purity beryllium single crystals tested in c axis compression* at 200°C and in Be-4.4 pct Cu and Be-5.2 pct cles in the lower-purity materials tested. The loops were related to surface "extrusions" observed on many of these same specimens. Nonbasal dislocations operating on (1122) planes with a (c + a) Burgers vector were observed. The presence of c and a dislocations together with (c + a) dislocations suggests that the (c + a) dislocations dissociate presumably on unloading or after failure of the test crystals to c and a dislocations. terial, (c + a) slip has only been observed near the the fracture4 surface in room-temperature c axis compression tests. Fracture in these tests occurs without measurable plastic flow, as determined with a strain sensitivity of 10"6. Since it has been shown for many metals that the application of hydrostatic pressure suppresses fracture,?-&apos; it was felt that studying the behavior of unalloyed beryllium single crystals stressed in c axis compression under a hydrostatic pressure would reveal whether (c + a) can occur if fracture was prevented, and that it might elucidate the role of (c + a) slip in the ductile to brittle transition. Evidence that (C + a) slip is associated with increased ductility in a high-pressure environment has been found in stress-strain tests on poly crystalline beryllium.10 The present paper describes a study on the influence of a hydrostatic pressure environment on the occurrence of (c + a) slip in beryllium single crystals. Material of two purity levels was tested in c axis compression over the pressure range ambient to about 27.5 kbars.* 1) MATERIAL PREPARATION AND CHARACTERIZATION Two lots of low-purity single-crystal beryllium were used. The first lot was "ingot secondary refined grade" and designated lot A. The second lot (lot B) was produced by a two floating zone pass zone-leveling operation in an argon-filled sealed quartz apparatus on a 1-in.-diam by 12-in.-long bar of Pechiney secondary refined-grade vacuum-cast and hot-extruded material. The high-purity material (lot C) was made by traversing twelve floating zone passes through a similar bar of Pechiney secondary refined-grade vacuum-cast and hot-extruded material. In a series of spark cutting and lapping procedures,11 single-. crystal specimens some 0.12 in. sq by 0.30 in. high were made with the sides Parallel to the first- and second-order prism planes and the basal plane within 3&apos; of arc of the top and bottom surfaces of the specimen. Such an accurate orientation is necessary because of the large difference in resolved shear stress

Jan 1, 1970

By W. G. Davenport, M. Paneni

Experiments in sheets of mercury and water have confirmed that the dynamics of bubble rise are simi-lur in Liquid metal and aqueous systems. RECOGNITION that bubble/metal interactions are important in pyrometallurgical refining processes (e.g., steelmaking, vacuum degassing, deoxidation of copper) has led to a number of investigations into the dynamic behavior of bubbles rising in liquid metals. Terminal velocities of bubbles rising in mercury (20°C)&apos; and silver (1000" to 1050"c)1,2 have been measured. Good agreement between aqueous and liquid metal systems has been obtained. Shapes of air bubbles rising in mercury have been determined experimentally by Davenport, Bradshaw, and Richardsonl who used a system of electric probes to measure maximum horizontal and vertical dimensions, Fig. 5. In the present work the motion and shape of nitrogen bubbles rising in a thin sheet of mercury (i.e., a "two-dimensional" environment) have been studied with the view of confirming dynamic similarity between aqueous and metallic systems. The use of thin sheets of liquid permitted direct visual and photographic comparisons of bubble behavior in clear liquids and in mercury. EXPERIMENTAL METHOD The thin sheets of mercury or water were held in a clear acrylic plastic container, Fig. 1. The sheets of liquid were 80 cm high by 30 cm wide and 0.47 cm thick. Single bubbles were formed by injecting the desired volume of nitrogen into an inverted semicircular cup (stainless steel). The cup was then rotated into the upright position using an external handle thus releasing a single bubble of measured volume. Nitrogen was injected into the cup from a nitrogen cylinder through the external handle and a hollow rotating axle. The volume of a bubble (volume range 0.01 to 10 cu cm) at any instant was obtained by connecting a soap film meter to a small gas space above the liquid. By this technique any change in the volume of a gas bubble in the liquid (i.e., during cup filling or during rise) produces an equivalent displacement from gas space A into the buret B. This displacement is determined by measuring the movement of a horizontal soap film in the buret. Bubble shapes were determined using single-frame photography. Rising velocities were determined by cinephotography with a 1/100-sec stopwatch in the field of view. In both cases dimensions were referenced to a rule attached to the container surface, also in the field of view. Linear dimensions were measured on projections of the single-frame negatives. Rising velocities were calculated from measured heights of rise (10 to 20 cm) and times of rise. The effects of disturbances caused by cup rotation and by bubble release were minimized by photographing bubbles at a point not less than 25 cm above the cup. MATERIALS Mercury: Technical-grade purified: minimum 99 pct Hg; Johnson Matthey and Mallory Limited, Toronto. Nitrogen: Technical-grade purified nitrogen: 99.99 pct Nz, 0.01 pct 02, trace argon; Air Liquide Canada Limitee, Montreal.

Jan 1, 1970

By Pervez Patel, Martin G. Frohberg, Demetrios Papamantellos

Model experiments were carried out on the mass transfer of iodine between n-hexane and a 4 2 X 10-3 M aqueous sodium hydroxide solution. The influence of mechanical stirring by means of small gas bubbles (ca. 0.01 ml in volume) on mass transfer was examined in independent experiments. The results were described with the help of the penetration theory and in the case of bubble stirring with the help of a mathematical model based on this theory. S LAG-METAL reactions, especially those in steel-making processes, often occur under conditions where gas is evolved in the form of bubbles that break through the slag-metal interface. It is known that stirring by means of bubbles increases the mass transfer coefficient, but there has been very little quantitative work done in this field. Because of the difficulties that attach to experimental work at high temperatures this study was confined to model experiments at room temperature. EXPERIMENTAL APPARATUS The transport of iodine from an n-hexane layer into a weak sodium hydroxide solution was studied a) with stirring by means of bubbles and b) with mechanical stirring. The apparatus for the experiments with bubbles consisted of the cell, two Dreschel bottles filled with hexane in order to saturate the nitrogen flow, one flow meter, and a nitrogen gas cylinder fitted with a fine-regulating valve. Fig. 1 shows the cell consisting of a glass cylinder 33 cm in height with an overall diameter of 7.5 cm embedded in a copper flange. The flange is connected to a Plexiglas plate and to the gas inlet. The Plexiglas plate has holes of 0.2 mm diam bored into it which are spaced 1 cm from each other. For the experiments with mechanical stirring the bottom of the cell was sealed off by a Plexiglas plate. The upper layer was stirred by means of a glass stir-rer and the bottom layer with the help of a magnetic stirrer in a counter rotatory direction. Both stirrer speeds could be varied continuously. EXPERIMENTAL PROCEDURE Iodine dissolves in n-hexane producing a stable violet color thus enabling the iodine concentration in

Jan 1, 1970