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By Charles E. Lundin

The character of the Er-D system was established by determining pressure-temperature-composition relationships. A Sieuerts&apos; apparatus was employed to make measurements in the temperature range, 473" to 1223"K, the composition range of erbium to ErD3, and the pressure range of 10~s to 760 Torr. The system is characterized by three homogeneous phase regions: the nzetal-rich, the dideuteride, and the trideuteride phases. These phases and their solubility boundaries were deduced from the family of isotherms of the system zchich relate the pressure-temperature-composition variables. The equilibrium plateau decomposition relationships in the two-phase regions were determined from can&apos;t Hoff plots to be: The differential heats of reaction in these two regions are AH = - 53.0 * 0.2 and -20.0 *0.1 kcal per mole of D2, respecticely. The differential entropies of reaction are AS = - 36.3 * 0.2 and - 31.0 * 0.2 cal per mole D2. deg, respectively. Relative partial molal and intepal thermodynamic quantities were calculated from the pure metal to the dideuteride phase. The study of the Er-D system was undertaken as a logical complement to an earlier study of the Er-H system.&apos; The primary interest was to compare the characteristics of the two systems and relate the difference to the isotopic effect. Studies of rare earth-deuterium systems by other investigators have been very limited in number and scope. Furthermore, there is even less information available wherein an investigator has systematically compared a binary rare earth-hydrogen system with the corresponding rare earth-deuterium system. The available information consists primarily of dissociation pressure measurements in the plateau pressure region of a few rare earths. Warf and Korst&apos; determined dissociation pressure relationships for the La- and Ce-D systems in the plateau region and several isotherms for each system in the dideuteride region. They compared these data with those of the corresponding hydrided systems. The study of these systems as a whole was very cursory and did not give sufficient data for a thorough comparison of the effect of the hydrogen vs the deuterium in the respective rare earths. The heat capacities and related thermodynamic functions of the intermediate phases, YH, and YD2, were determined by Flotow, Osborne, and Otto,~ and the investigation was again repeated for YH3 and YD3 by Flotow, Osborne, Otto, and Abraham.4 This investigation studied only these specific phases. Jones, Southall, and Goodhead5 surveyed the hydrides and deu-terides of a series of rare earths for thermal stability including erbium. They experimentally determined isotherms of selected hydrides and plateau dissociation pressures for deuterides. These data allowed comparison of the enthalpy and entropies of formation of the dihydrides and dideuterides. To date, no one rare earth has been selected to thoroughly establish the complete pressure-temperature-composition (PTC) relationships of binary solute additions of hydrogen and deuterium, respectively. The objective in this investigation was to provide the first comparison of a complete family of isotherms of a rare earth-deuterium system with those of a rare earth-hydrogen system. This would allow one to determine what differences exist, if any, in the various phase boundaries and the thermodynamic relationships in various regions of the systems. I) EXPERIMENTAL PROCEDURE A Sieverts&apos; apparatus was employed to conduct the experimental measurements. Briefly, it consisted of a source of pure deuterium, a precision gas-measuring buret, a heated reaction chamber, a mercury manometer, and two McLeod gages (a CVC, GMl00A and a CVC, GM110). Pure deuterium was obtained by passing deuterium through a heated Pd-Ag thimble. A 100-ml precision gas buret graduated to 0.1-ml divisions was used to measure and admit deuterium to the reaction chamber. The reaction unit consisted of a quartz tube surrounded by a nichrome-wound furnace. The furnace temperature was controlled by a recorder-controller to . An independent measurement of the sample temperature in the quartz tube was made by means of a chromel-alumel thermocouple situated outside, but adjacent to, the quartz tube near the specimen. Pressure in the manometer range was measured to k0.5 Torr and in the McLeod range (10~4 to 10 Torr) to *3 pct. The deuterium compositions in erbium were calculated in terms of deuterium-to-erbium atomic ratio. These compositions were estimated to be *0.01 D/Er ratio. The erbium metal was obtained from the Lunex Co. in the form of sponge. The metal was nuclear grade with a purity of 99.9+ pct. The oxygen content was reported to be 340 ppm and the nitrogen not detectable. Metallographically the structure was almost free of second phase (<i vol pct). A quantity of sponge was arc-melted for use as charge material. The solid material was compared with the sponge in the PTC relationships. They were found to be identical. Therefore, sponge material was used henceforth, so that equilibrium could be attained more rapidly. The specimen size was about 0.2 gr for each loading of the reaction chamber. The procedure employed to obtain the PTC data was to develop experimentally a family of isothermal curves of composition vs pressure. First, a specimen

Jan 1, 1969

By S. Luszcz, R. W. Vook, Fred Witt

In situ X-ray investigations were made on polycrys-talline silver films deposited by vacuum evaporation on (111) CaF2 and (100) NaCl single-crystal substrates at 80°K. The films were evaporated and annealed in an X-ray diffractometer attachment having a residual gas pressure of 2 x lo-&apos; Torr. All measurements were made without exposing the films to the atmosphere. Measurements were made on the films in the as-deposited state and after various annealing treatments. The intrinsic stacking and twin fault densities, the magnitudes of the uniform and nonuniform strains, and the crystallite sizes were determined. In addition the textures in the films were measured qualitatively. The results obtained for the as-deposited films on single-crystal substrates are in substantial agreement with previously reported results for silver films deposited on glass. Intrinsic stacking and twin faults, as well as uniform and nonuniform strains, were present in these films. During the various annealing treatments (up to 350°C) the faults and nonuniform strains annealed out. Considerable grain growth and texture changes occurred also. The effects were much greater for the NaCl substrate than for the CaFz substrate. The relative magnitudes of the grain growth in the variously oriented grains could be explained qualitatively in terms of the thermal strains and strain energies introduced into the differently oriented grains during the initial, irreversible anneal. These strains were due to the different thermal expansion coefficients of the film and substrate. X-RAY diffraction measurements on evaporated films deposited on substrates at low temperature have the advantage that many of the imperfections introduced into the film during deposition are "frozen in". Thus, the influence of a very important experimental variable, substrate temperature, on the imperfection structure of evaporated metal films may be studied. Moreover, the effects of annealing such films makes possible the study of thermally activated recovery processes in these films. The present study was designed to determine the influence of single-crystal substrates on the resultant film structure relative to the previous results obtained using glass substrates.&apos; To this end great care was taken to keep the experimental variables the same in the two cases. Different experimental conditions would, of course, result in films having different physical properties. Again the initial substrate temperature was in the neighborhood of 80°K and the films were subsequently annealed to 350°C. The pure metal silver was chosen for evaporation, primarily because of its relatively low stacking fault energy and consequent high fault density in the as-deposited state. The silver films were formed by evaporation onto air-cleaved {ill} CaF, and (100) NaCl surfaces cooled to 80°K in an X-ray diffractometer attachment2 having a base residual gas pressure of 2 X l0-&apos; Torr. The films were not exposed to the atmosphere until all of the X-ray data had been recorded. In this way one of the most important experimental variables, environment, could be well-controlled and reproduced. X-ray measurements were made at the temperature of deposition and included determinations of the diffraction line peak positions, line shapes, and integrated intensities. The peak position measurements were used to determine the intrinsic stacking fault densities and the average uniform strain in the film. The shapes of the diffraction lines provided information on the twin fault density, true crystallite size, and average nonuniform strain. The preferred orientation in the film was determined qualitatively from the integrated intensities. I) EXPERIMENTAL PROCEDURE The evaporator attachmentZ was charged with 99.999 pct Ag pellets positioned in a tantalum filament which had been outgassed previously at l0-8 Torr. The CaFz and NaCl single crystals were cleaved in air and then placed in position in the chamber so that their cleavage surfaces were on the diffractometer axis. The chamber was prepumped using a sorption pump, sealed off, and then baked at 150°C for 24 hr. The ion pump operated during the bakeout cycle. The substrate was then heated to 500° C by means of an auxiliary heater and kept hot until the rest of the chamber was cooled slowly to room temperature. This bakeout procedure consistently resulted in an ultimate pressure in the low lo-&apos; Torr range. The substrate was then cooled down on 80°K. Its temperature was monitored by a thermocouple wedged into the rear of the copper substrate holder. The diffracted intensity and peak position of the 111, 222, and 333 CaF, lines were measured prior to evaporation. Nickel-filtered, pulse-height-discriminated copper radiation was used. Similar measurements were made for the 200 and 400 lines from NaC1. These measurements were used as a lattice parameter check and to determine the thickness of the evaporated silver films from the attenuation of the substrate lines. The evaporation rates were approximately 3A per sec for both films while the maximum pressures during evaporation were 3 x lo- &apos; and 7 x 10"8 Torr for the CaF, and NaCl cases, respectively. The film thickness was measured by the attenuation of the CaF, and NaCl substrate lines and by at optical interference method. Values of 1700 and 1500A, respectively, were obtained for the silver

Jan 1, 1969

By M. H. Richman, W. D. Sproul

THE indexing of a field-ion micrograph is in many cases similar to the indexing of the poles of a stereo-graphic projection.* If one has a standard projection as a guide, the task of indexing is relatively simple. In this way, field-ion micrographs of bcc and fcc wires have been oriented from projections available in most any text on X-ray diffraction. These books always illustrate the projections for the bcc and fcc structures and even for a particular hcp structure. There is no "standard" hcp projection because the angles between planes depend on the c/o ratio of the unit cell. In order to index field-ion micrographs of hcp metals, one must compute the angles between various planes—a rather laborious task. In recent work,, we have had to identify the structure and index the orientation of field-ion micrographs of the compounds WC and W,C. To overcome the problem of computing angles by hand or calculator, a program was written in Fortran lV for use on the IBM 360, Model 50 which will compute all angles between planeswith indices (hlklll) and (17,k,l2) in the range from (447) to (444) for any given c/o ratio. The program, which is more specific than that of Camp and clum, was so written as to take advantage of the multiplicity of the hexagonal system, and, for the limits defined above, only 26,649 values out of a possible 265,721 values had to be computed. This is a 90 pct saving in time (and cost) and was accomplished in the following way. Because the lattice parameters of the hexagonal crystal system are defined such that: a = b * c [1] the I index does not enter into the multiplicity (i.e., it cannot be transposed with the h or k indices), and all possible values of h and I, from 4 to 4 must be used—a total of eighty-one combinations. The index h2 varies the most—it must assume values from - Ill,: to 3Ihl. The indices kl and k2 are always equal and vary

Jan 1, 1969

By Edward A. Loria

Vanadium is one of the materials being considered for use in fast-breeder nuclear reactors. It is a good heat conductor and nearly transparent to radiation, properties regarded as ideal for nuclear fuel containers. The problem has been the metal lacks ductility needed for forming. In our program at Reno Laboratory the effects of the U.S. Bureau of Mines (previous affiliation) the effects of oxygen, nitrogen, and carbon have been reported21,22 nd most recently the effects of boron and boron with carbon.23 The continuing study by the Ames Laboratory of the effects of hydrogen, the most subtle and fugitive element in its relationship to structure and properties, is commendable. The solubility of oxygen, nitrogen, and hydrogen in vanadium (and other bcc metals as well) is related to material purity. With the available limit on metal purity and the interaction (clustering) of hydrogen with the other interstitial and substitutional elements present in their base metal, would the authors comment on the tolerance of hydrogen in relation to oxy- gen that will avoid intergranular fracture and above what content or how seriously harmful in oxygen at various carbon levels? Understanding the true effects of hydrogen (and other interstitial impurities) in vanadium depends on knowledge of the amount present under various conditions and its distribution between various possible positions. Significant amounts of hydrogen may be in "thermo-dynamic" solution, but not as individual atoms in interstitial positions. It may be in grain boundaries, in subgrain boundaries, in vacancies, or in dislocations. Some of the hydrogen may be paired or in complexes with impurities in substitutional solid solution in vanadium. This makes the analysis of the effect of hydrogen vastly more difficult. It should not be expected that hydrogen in grain boundaries will affect slip in the same manner as hydrogen in interstitial positions, nor will the complexes react to heat treatment the same way as hydrogen in dislocations. It is gratifying that the authors have quoted results from two of our inestiations,&apos; but neglected one2&apos; which applied the Petch theory to electrorefined vanadium of high purity and analyzed the effects of oxygen, nitrogen, and carbon on the ai and ky values obtained by the extrapolation and grain size methods. Details can be found in the paper. Briefly, we obtained a ui value of 10,300 psi (t = 5150 psi) at 273°K. The metal also exhibited profuse twinning with high ductility at 123" and 77°K. Nitrogen was more effective than oxygen in increasing ui, the increase being about twice as much for the same impurity level. Carbon additions resulted in only minimal changes in Ui. The corresponding ky increased as oxygen and nitrogen increased, with nitrogen being more effective. Carbon additions resulted in little change. Although the difference in yield strength of vanadium was explained on the basis of the Cottrell-Bilby model of the tearing away of dislocations from the impurity atoms,

Jan 1, 1970

By Thomas C. Wilder

ALTHOUGH there is an abundance of thermodynamic information available in the literature today, no experimental data on the standard molar properties of formation of zinc oxide have ever been reported for the temperature range 419.5" to 908°C. This is the temperature range for which zinc metal is a liquid. so that the formation reaction is: Data have been reported for Reaction [I] where zinc is a solid or a gas. Maier. Parks, and nderson&apos; measured the electromotive force of the cell: in order to obtain the free energy of the reaction: over a limited temperature range. With the known properties of H2O, those of ZnO were calculated. These workers found that t;j of ZnO is -75,930 ± 150 cal per gfw and Mtf is-83,000 300 cal per gfw at 25°C. akolkin made similar measurements to obtain ilF7 and AH; of -75,870 and -82,540 cal per dw respectively at the same temperature. Truesdale and arirlg and Kitchener and 1gnatowicz4 measured the equilibrium of the reactions: respectively, in order to determine Ff of ZnO from the known values of CO, CO2, and H/X The thermodynamic properties of ZnO, presumably taken from the above investigations, are tabulated by Kubaschewski and vans&apos; and Wicks and Block" The latter reference also gives values for the range where zinc is a liquid. These numbers are probably acquired from the known extremities of the range plus Cp data of the components. In the present work the electromotive force of the cell: was measured in order to calculate the properties of ZnO from the well known values of NiO for the temperature range where zinc is a liquid. In Cell [B] the cathode reaction is: and at the anode: The overall reaction is: for which the free energy change is: where is the electromotive force in volts of cell [B] after correction for the Pt-Ta thermocouple. The standard molar free energy of formation of NiO is very well known as a function of temperature. Kiukkola and agner&apos;s data are used in most tabulations. The error is probably less than I00 cal per gfw, The latter authors were the first to use solid mixed oxide electrolytes in galvanic cells. Their work with NiO and other oxides was performed with cells similar to Cell [B] using a variety of reference electrodes of fixed oxygen potential. There may be a small composition range for which either NiO, ZnO or both exist. If nonstoichiometry does exist for these oxides, then the properties described herein are those for the metal saturated compositions. The simple cell design was the same as described elsewheres except that the molten metal electrode was about 125 g of pure zinc including a small (0.25 to 0.50 g) amount of ZnO. The mixture of pure nickel and NiO powders making up the reference electrode was contained in a tube of the ZrOl + CaO electrolyte which was immersed in the molten zinc. A platinum contact was made to the reference electrode; tantalum to the zinc. A refractory cement plug was necessary at the open end of the electrolyte tube in order to prevent zinc vapor from contaminating the reference electrode. Experimental details such as the vessel containing the cell, furnace, temperature control and measurement, and electromotive force determinations of the cell have all been outlined previously. The temperature was varied randomly in the range between the freezing and boiling points of zinc. After each temperature change, time was allowed for the drift at the new temperature to be less than +;"c whereupon a series of three cell readings was taken at least 10

Jan 1, 1970

By L. W. Sink, B. H. Kear, S. M. Copley, M. F. Hornbecker

IT is well known that under appropriate conditions of chill casting, many metals and alloys develop a columnar grain structure having a high degree of preferred orientation.&apos; In cubic materials, this principle has been employed to produce shaped castings with a unidirectional cube oriented columnar grain structure as well as shaped single crystals with a cube axis parallel to the growth ais.&apos; "he purpose of this note is to describe how this same principle can be enlployed to produce wires, rods, or more complex crystal Shapes with a variety of conlpositions in cube or random orientations by a single casting operation. The procedure for directional solidification of in Fig. 2. Examination under crossed polarizers indicated that the striations were a second phase (presumably the B19 phase) which frequently forms initially at primary grain boundaries and then grows toward the interior as thin ribbons. X-ray studies on filings of this alloy reannealed at 1300°C for 15 min indicated a strong pattern of the fcc phase together with a substantial pattern of the B19 phase. The un-annealed filings, however, showed only the pattern of a cold-worked, fcc structure. In these two-phase alloys it appears, therefore, that the B19 phase is transformed into a disordered fcc structure by cold-working. The mechanical properties of these alloys would be of some interest since a fcc structure might be expected to possess greater ductility than the or-thorhombic B19 structure from which it forms. The author wishes to thank Mr. Harvey Yakowitz and Mr. Richard Manuszewski for obtaining the photomicrograph in Fig. 2. This research was partially supported by Research Grant DE-02455 from the National Institute of Dental Research. columnar grained ingots involves pouring a melt into a preheated mold having a water-cooled copper chill at its base? and controlling the temperature gradient in the mold so as to maintain columnar grain growth from end to end of the ingot.4 This is illustrated schematically in Fig. l(a). It should be noted that the longitudinal temperature gradient promoting unidirectional columnar grain growth is maintained by heat loss primarily through the solidifying ingot to the chill, so that both the temperature gradient and the growth rate diminish with increasing distance from the chill. As the distance from the chill increases? the columnar structure coarsens, and the perfection in preferred orientation tends to increase progressively. Using this same procedure of gradient solidification, Fig. l(b) shows a simple scheme for producing a batch of crystals having the same cube orientation rather than an assemblage of columnar grains. Wires and rods are produced by cementing ceramic tubes to the wall of the mold. The steep temperature gradient produced in the central chamber is controlled by the mold heater to produce favorable conditions for crystal growth. Clearly the success of the operation lies in the selection of individual growing columnar grains by some kind of orifice at the base of each channel in the mold. Using simple constrictions, experience with nickel-base alloys has shown that 4 -in.-diam holes at a distance of -1 in. from the chill yield crystals within 7 deg of (100); closer control of orientation is possible by placing the crystal selector farther away from the chill.

Jan 1, 1970

By B. C. Wonsiewicz, W. W. Wilkening

Following a procedure proposed by Wheeler and Ireland, Plane stress yield loci were constructed from Knoob hardness numbers. Basically, six differently oriented hardness measurements were made on three orthogml surfaces through pure poly crystalline magnesium sheet, a magnesium single crystal, and sheet of the magnesium alloys: Mg + 0.5 pct Th, Mg + 4 pct Li, AZ31B, and EKOO. Hardness loci were found to be in poor agreement at small strains (E < 0.05) with loci established by a more rigorous technique. At larger strains (E - 0.10) the agreement is fair, but at this stage in deformation the conventional locus has lost much of the asymmetry that characterizes these anisotropic materials. Two effects which will lead to distortions in the Khn locus are discussed with reference to the geometry of plastic flow during a hardness test. DETERMINING a material&apos;s resistance to multiaxial loading is of interest not only from a structural design viewpoint but also from that of deformation processing. Unfortunately, the determination of the yield locus, although simple in principle, involves tedious procedures if the results are to be at all rigorous.&apos; The idea, first proposed by Wheeler and 1reland2 of determining the yield locus by means of six Knoop hardness impressions along the principal directions in a material has obvious appeal. It is simple, quick, and should be applicable to very thin sheets. If such a technique could be demonstrated to produce consistently reliable results, it would be of interest to both researcher and designer. Lee, Jabara, and ackofen have compared the yield locus determined by Knoop hardness measurements (the Khn locus) to a locus determined by more rigorous techniques. They found good agreement for two titanium alloys at a plastic strain of about 1 pct. The purpose of this paper is to investigate if the Khn locus construction is a reasonable approximation to the locus of a highly anisotropic material. Examples of such materials are magnesium and magnesium alloys which have severely distorted yield loci which in turn reflect markedly dfferent yield strength in different directions.&apos; In pure magnesium, for example, the yield stress in tension along the transverse direction may be four times the yield stress in compression in the same direction and twice the tensile yield stress in the rolling direction. Predicting such large differences ought to serve as a severe test of the Khn locus construction. EXPERIMENTAL PROCEDURES Samples of rolled sheet, 0.250 in. (6.35 mm) thick, of pure magnesium and four magnesium alloys (Mg experimental materials. The pure magnesium together with the lithium and thorium alloys were those used in the study of Kelley and Hosford. The grain size was ASTM number 4 for the pure magnesium and number 6 for the alloys. HARDNESS TESTING The materials were sectioned along the rolling and transverse planes, mounted in a quick setting resin, and mechanically polished. Most of the hardness tests were performed on a surface prepared by electro-polishing (30 pct nitric acid in methanol at 0°C and 20 v) with the exception of the AZ31B and EK00 alloys which were made directly on a metallographically polished surface. However, subsequent hardness tests on the same sample after heavily electropolishing, revealed essentially the same hardness as before. At least twenty Knoop hardness impressions under a 100-g load were made in each of the six orientations shown in Fig. 1. The average hardness number and standard deviation were then calculated for each orientation. CONVENTIONAL LOCUS CONSTRUCTION Yield loci were constructed using a technique described in detail by Lee and ackofen,&apos; in which the flow stress (stress at a given plastic strain) fixes the coordinates of a point on the locus and measurements of the strain ratio serve to establish the slope of the locus at that point. The loading paths which correspond to uniaxial tension or compression tests establish the four intercepts of the locus with the coordinate axes plus one point on the balanced biaxial tension line Tensile testing was performed along the rolling and transverse (r, t) directions. Samples had a uniform rectangular gage length 1 by 4 by 4 in. (25.4 by 6.35 by 6.35 mm) and were deformed at a strain rate of 3.33 x 104 sec-&apos;. The tests were interrupted periodically to unload the sample and measure the plastic strains by means of X-Y post yield strain gages. Compression tests in the rolling, transverse, and through-thickness (r, f, z) directions were performed on 1/4 in. (6.35 mm) cubes at an initial strain rate of 8.33 x sec-&apos;. Lubrication was provided by 0.002 in. (51 pm) Teflon sheet which was renewed after unloading for micrometer measurements used to calculate the strains.

Jan 1, 1970

By R. J. Fruehan

The oxygen activity and concentration were measured in Fe-Cr-0 melts in equilibrium with an oxide phase at 1600°C (2912°F). The activity was determined by ,use of the following solid oxide -electrolyte galvanic cell CY-Cr8,(s) I ZrOz(CaO) I Fe-CY-G(saturated)(l) The oxygen concentration decreases with increasing Cr concentration to about 270 ppm 0 at about 7pct CY and then increases gradually. The activity coefficient of oxygen (fo) decreases with increasing Cr. In melts containing up to about 20 pct Cr, log f is approximately a linear function of wt pct Cr with a slope (e q 2) of —0.037. The activity of chromium was calculated and found to exhibit a small negative deviation from Raoult&apos;s law. From the activity and solubility data for low chromium melts, the free energy of formation of chromite, FeCr204, was found to be -79.8kcal per mole where pure liquid chromium and oxygen at I wt pct in Fe are the standard states. ThE effect of chromium on the chemical behavior of dissolved oxygen in liquid iron is of great importance in controlling the deoxidation of steels containing a significant amount of chromium. Chen and chipman&apos; equilibrated Fe-Cr melts in the presence and absence of slag with hydrogen-water vapor mixtures. They concluded that at 1595°C chromite was the oxide phase in equilibrium with Fe-Cr alloys containing less than 5.5 pct Cr while at higher chromium concentration Cr,O, was the stable phase. In the composition range 0 to 10 pct Cr they found that the interaction coefficient, was equal to -0.041. Turk-dogan,&apos; Schenck and Steinmetz,, and pargeter4 measured egr) in a similar manner and found the value to be -0.064,-0.04, and -0.052, respectively. McLean and Be11 evaluated egr) from their data on the equilibrium of Fe-Cr-Al-0 alloys with H2/H20 mixtures and found it to be -0.058. However, McLean and Bell&apos;s value should only be considered an estimate because the effect of aluminum on the activity coefficient of oxygen is about a hundred times greater than that of chromium. Consequently, an error in the value of egl) used, which at the present time is not well-known, or an error in aluminum analysis, which is present in very small quantities, will result in a significant error in egr). Fischer et a1.6 determined the interaction coefficient (eEr) in Fe-Cr-0 melts not in equilibrium with an oxide phase and containing less than 18 wt pct Cr at 1600°C electrochemically. They determined a value of -0.031 for egr). Hilty et aL7 measured the oxygen content of Fe-Cr melts in equilibrium with an oxide phase containing up to 50 pct Cr. They found that the solubility of oxygen decreased as the chromium content increased to about 6 pct Cr and then increased gradually. They concluded that the equilibrium oxide phase was chromite below 3 pct Cr, distorted spinel from 3 to 9 pct Cr, and Cr,04 above 9 pct Cr. Adachi and lwamotoa also investigated this system, but did not find Cr30,. They X-rayed the equilibrium oxide phases and did not find the presence of Cr,O,. They also X-rayed the oxide phase extracted from a 65 pct Cr melt which was heat treated and did not find metallic chromium as would be expected if Cr3O4 were the equilibrium oxide phase as indicated by the reaction : 3Cr3O4 — 4Cr2O:, + Cr [lj It was the purpose of the present investigation to determine the effect of chromium on the activity coefficient of oxygen in Fe-Cr melts by measuring the activity and solubility of oxygen equilibrated with an oxide phase in the composition range 0.18 to 50.5 wt pct Cr at 1600°C (2912°F). The activity of oxygen in the melts was determined by use of the following galvanic cell: The relationship between the partial pressure of oxygen in equilibrium with the melt and the reversible electromotive force of the cell (E) is where 11 = 4, F is the Faraday constant, pb, is the oxygen pressure in equilibrium with the meit and is the oxygen pressure in equilibrium with Cr203 as determined from the free energy data compiled by Elliott et al? The oxide phase in equilibrium with pure chromium was assumed to be Cr If Cr30, were the equilibrium phase the activities derived would be approximately the same, since the best estimated free energy of formation of Cr,O,, if it does exist, is approximately % the free energy of formation of The activity of chromium in Fe-Cr alloys at 1600° C was also determined from the measured electromotive force. The activity of chromium (aCr) is related to the electromotive force as follows: , The oxide phase in equilibrium with pure chromium and Fe-Cr melts from 10 to 52 pct is assumed to be Cr203 so that n equals three. If future work proves the existence of Crs04 in equilibrium with Fe-Cr melts and pure chromium, the experimental results can be reevaluated using a value of $ for n in Eq. 141. A value of ^ for n will make the activities about 10 pct higher. In order for Eqs. 131 and [4J to be valid the electrolyte, ZrOa(CaO), must exhibit predominantly ionic conduction at the temperature and oxygen partial pressure of its use. Previous work1&apos; has demonstrated that ZrOz(Ca0) is predonlinantly an ionic conductor

Jan 1, 1970

By C. D. Williams, C. E. Ells

The susceptibility of quenched and aged Zr-2.5 wt pct Nb alloy to embritt2ement during irradiation has been examined for a number of solution temperatures and aging times. Material quenched from temperatures approximately 40°C below the transus has high tensile ductility, and this ductility is insensitive to aging at 500°C or to irradiation. If, however, the material is quenched from temperatures above the transus it becomes highly susceptible to loss of ductility either from aging at 500 or from irradiation. Inter granular failure is characteristic of the materials having low ductility. The distribution of the equilibrium phase is found to control the susceptibility to embrittlement by restricting 6 grain growth during heat treatment and thus influencing crack propagation. IN zirconium, as in titanium, -stabilizing alloy additions are used to obtain high strengths via quench and age heat treatments, and the Zr-2.5 pct Nb alloy has been developed1 because of its strength advantage over the Zircaloys. Early in the development of the Zr-2.5 pct Nb alloy the problem of 13 embrittlement was appreciated, and for this reason the solution temperature was chosen below the p transus.&apos; In the course of irradiation studies on quenched and aged Zr-2.5 wt pct Nb alloy it was found&apos; that irradiation introduced an important aspect of p embrittlement, riz., material quenched from the phase and aged 24 hr at 500°C was severely embrittled by moderate doses of neutron irradiation. This effect had not been studied in titanium alloys. In titanium the metallurgical features leading to 0 ernbrittlement were found to be structures with: a) coarse a platelets at the grain bondaries, b) finely dispersed a uniformly distributed throughout the (0) matrix,6 c) Widmanstatten a-13 with more than 50 pct P, d) the presence of some metastable p transformation products,3 and e) large prior -phase grain size.5 Alternatively, the presence of a uniform distribution of coarse a was conducive to high ductility and a structure largely of equiaxed a was very dctile. The detailed mechanisms of the embrittlement have not been worked out for all of these conditions, although weakness at either a-matrix boundaries or prior p grain boundaries have been prominent in the eculation. It was proposed that acicular a might act as a mild notch, and low ductility has been associated with easy fracture along its boundary.&apos; There have been two opposing suggestions for the source of the high ductility associated with equiaxed a phase. JaffeeB proposed that this a would accept a large por- tion of the oxygen, thus increasing the ductility of the matrix, whereas after study of a Zr-Nb-Cu alloy Weinstein and oltz proposed that the a phase, softer than the martensitic matrix, acted to blunt cracks formed in the matrix. In the present work we have studied the effect of neutron irradiation on the ductility, particularly the P embrittlement, of the Zr-2.5 wt pct Nb alloy. By a variation of solution temperature and aging time a variety of metallurgical conditions have been examined, and a range of resultant ductilities obtained. The ductility has been related to the material microstructure and mode of fracture. EXPERIMENTAL The alloy used in the present work came from two separate ingots fabricated into rod of 3/8 or i in. diam, Table I. For both batches the P transus temperature was approximately 890° C. Most of the heat treatments were done directly on lengths of the j} in. diam rod, after which the tensile test specimens were machined. Quenching was achieved by dropping rods from a dynamic vacuum into water, the cooling rate estimated to be 2 100°C per sec. For aging the rods were encapsulated in evacuated silica tubes. Round tensile test specimens, with gage diam and length 0.160 and 1.0 in., respectively, were used throughout and pulled at room temperature or 300°C on Instron tensile machines, at a crosshead speed of 0.05 ipm. Specimens were irradiated in the NRX and NRU reactors, in facilities described in previous publications.&apos;0 The metallurgical conditions examined have been: All tensile test specimens were machined with axes in the axial direction of the swaged rod. Although the specimen had a degree of preferred crystallo-graphic orientation with basal plane normals both parallel with and perpendicular to the tensile axis, the material was comparatively isotropic." The techniques of thin foil examination in the electron micro-

Jan 1, 1970

By Mark J. Klein

The creep of 043 was studied over the temperature range 1650" to 3200°F and over the stress range 3000 to 44,000 psi. The steady-state creep rate over this range of stress and temperature can be expressed by the equation where A is a constant, is the stress, and is -0.8 x 103 psi-&apos;. Over a narrow range of stress variations c0 a and for this proportionality n varies from 3 to 30 in accordance with the relation n = aB. Above about 2400° F, H, the apparent activation energy for creep, is 110,000 cal per mole, a value about equal to that estimated for self-diffusion in this alloy. Below 2400°F, H increases with decreasing temperature reaching a value of -125,000 cal per mole at 1700° F. In this temperature region, H appears to be a function of the interstitial concentration of the alloy. MOST of the detailed creep studies of dispersion strengthened metals have been concerned with metals having fcc structures. However, there are a number of important refractory alloys with bcc structures that derive part of their high temperature strength from an interstitial phase and whose creep behavior has not been well defined. This paper describes the creep behavior of the bcc alloy, D43, over the temperature range 1650" to 3200°F (0.4 to 0.7 Thm) and over the stress range 3000 to 44,000 psi. In addition to colum-bium, this alloy contains 10 pct W. 1 pct Zr, and sufficient carbon (-0.1 pct) to form a carbide dispersion throughout the matrix of the alloy. The effects of variations in temperature and stress on the steady-state creep rate of this alloy are presented in this paper. EXPERIMENTAL PROCEDURES Creep tests were made in a vacuum of 106 torr under constant tensile stress conditions using a Full-man-type lever arm.&apos; Creep specimens were machined from 0.020-in. D43 sheet (grain size -5 x l0-4 in.) processed in a duplex condition (solution annealed -2900°F, 40 pct reduction in area, aged 2600°F). The specimens were tested in this condition without further heat treatment. Specimen extensions over 1-in. gage lengths were continuously recorded using a high temperature strain gage extensometer. Differential temperature and stress measurements were used to determine temperature and stress dependencies of the creep rate. Activation energies were calculated from the changes in strain rate induced by abrupt shifts in the temperature during constant stress creep tests. The 100°F temperature shifts used in most of the activation energy determinations required 15 to 90 sec depending upon the temperature at which the shift was made. The dependence of strain rate on stress was determined by measuring the change in strain rate for incremental stress reductions during constant temperature tests. It has been shown that columbium-base alloys such as D43 are susceptible to contamination by gaseous interstitial elements during vacuum heat treatments.&apos; In this regard, it is unlikely that these alloys can be heat treated without some loss or gain of interstitial elements despite the precautions taken to control the heat treating environment. However, several factors suggest that changes in interstitial concentrations of the specimens during testing did not affect the results presented in this paper. First, the dependence of the creep rate on the stress or temperature determined during the course of a single creep test showed no variations with the duration of the test. A variation would be expected if a loss or gain in interstitial concentration during the course of the test affected results. In addition, precautions taken during this investigation to minimize interstitial contamination by wrapping the gage lengths of the specimens with various foils2 (Mo, Ta, W) did not produce a detectable change in the stress and temperature dependencies relative to the unwrapped specimens. The averages of duplicate analyses for carbon and oxygen in several specimens determined before and after creep testing are listed in Table I. The combined nitrogen and hydrogen concentrations which were ordinarily less than 50 ppm did not change in a detectable way with creep testing. The analyses show that only minor changes in carbon concentration occurred during creep testing except for specimen 4. This specimen which was tested at 3100°F lost a significant amount of its carbon concentration to the vacuum environment. Specimen 1 gained 100 ppm of O, while specimens 2, 3, and 4, which were tested at progressively higher temperatures, lost increasing portions of their initial oxygen concentrations during testing. RESULTS AND DISCUSSION The Temperature Dependence of the Creep Rate. The apparent activation energy for creep, H, was de-rived from creep curves similar to that shown in Fig. 1. Steady-state creep was rapidly attained at the beginning of the test and with each change in temperature. This behavior suggests that the alloy rapidly attains a stable structure with each shift in temperature or that the structure is constant throughout the test. Since the dispersion will tend to stabilize the structure, the latter is probably the case. The activation energy was found to be independent of the direction of the temperature shift and the magnitude of the shift (50" or 100°F). Although H was approximately independent of the strain, there was a tendency for it

Jan 1, 1970

By Neils Hansen

The microstructure of dispersion-strengthened aluminum aluminum-oxide products containing from 0.2 to 4.7 wt pct of aluminum oxide has been examined by optical and transmission electron microscopy, and the flow stress has been determined at room temperature and at 400C by tensile testing. Products were examined as recrystallized and as high-temperature extruded, and the microstructures consisted of a fine dispersion of oxide particles in a matrix divided by respectively recrystallized grain boundaries and subgrain boundaries. The flow stress (0.2 pct offset) at room temperature of recrystallized dispersion strengthened aluminum aluminum-oxide products is the superposition of dispersion strengthening and grain boundary strengthening. This superposition has been found to be linear. The flow stress (a) can be related to the grain size (t) by the Petch equation: ing content of oxide and k is a constant independent of the oxide content. For extruded products a similar relation has been found by replacing the grain size by the subgrain size. The k-value is of the same order for the two types of structure, which shows that the subgrain boundaries are as effective slip barriers as grain boundaries. Tensile testing at 400C of re-crystallized and extruded products shows that oxide dispersion strengthening is very effective, whereas the strengthening effect of grain boundaries and subgrain boundaries is small. THE microstructure of dispersion-strengthened products consists of hard particles finely distributed in a metal matrix. The strengthening effect of the dispersed phase has been fairly well established,1 and it has been found that the size and volume fractions of the dispersed particles are important structural parameters. However, in many dispersion-strengthened products which have been worked and heat-treated during manufacture the matrix is divided into well-defined grains or sub-grains, which may also have a strengthening effect. A model of the matrix strengthening in dispersed products worked during manufacture has been proposed,2 introducing the energy of the structure as a strengthening factor, especially at low temperatures. A difficulty in this model is, however, to relate this (stored) energy to the structural parameters directly observable as for instance grain size. The strengthening effect of the matrix grain size after recrystallization has been in- vestigated for nickel-thoria (TD-Nickel) products3 and for copper aluminum-oxide products. Conclusive results were, however, not obtained as the grain size of TD-nickel was constant. 5 to II , after recrystallization at temperatures from 700 to 1200°C and as the copper products containing 5 to 1 wt pct of aluminum oxide could not be recrystallized even after severe cold reduction and heat treatment at 1050C. For aluminum aluminum-oxide products containing from 1 to 5 wt pct of aluminum oxide it has been shown that the tensile strength at room temperature decreases when an extruded product is cold-worked and recrystallized. The matrix in the extruded products is divided into well-defined subgrains of micron size, and as the grain size of the recrystallized products is about two orders of magnitude higher, it is obvious that grain boundary strengthening occurs. Preliminary results8 have indicated that the flow stress containing no grain boundaries, A is a constant and t is the subgrain size. At elevated temperatures the effect of boundaries is more complex; it has been shown11 that recrystallized products having an oxide content of about 3 wt pct are more creep resistant than extruded material in the temperature range 400° to 600°C, whereas on application of a higher strain rate the tensile flow stress (0.2 pct offset) is higher in extruded than in recrystallized aluminum—5 wt pct aluminum oxide products at temperatures from room temperature to 427°C (800), Finally it has been shown12 that the Brinell hardness at 350°C of extruded products having about the same content of aluminum oxide increases with decreasing grain size, determined by X-ray line-width measurements. The present study was undertaken to obtain a quantitative relationship between the tensile strength and the grain size of aluminum aluminum-oxide products in the recrystallized as well as in the extruded state. The tensile testing was performed at room temperature and at 400uC. The grain size of the recrystallized products was varied by changing the degree of cold-work preceding the recrystallization heat treatment. In extruded products grain (or subgrain) size variations were obtained by high-temperature heat treatment after extrusion. EXPERIMENTAL a) Materials. Aluminum aluminum-oxide products have been manufactured by consolidation of aluminum powder covered with a layer of aluminum oxide formed during powder manufacturing. The products examined were manufactured from atomized powder containing

Jan 1, 1970

By Thomas C. Wilder, Walter E. Galin

Measurements of the electromotive force of the cell at 995°C have shown that the cell may be used to detennine the zinc content of molten Cu-Zn alloys to the nearest 0.05 wt pct. The cell is used for brass melts to which no ZnO is added intentionally, because essentially all oxygen in the Cu-Zn-O system is present in the form of ZnO. The cell is also used to detev,nine the thernodynamic activity of zinc for Cu-Zn al1oys of the brass composition range, from which the equilibrium partial pressure of zinc for such alloy may also be calculated. The electrochemical measurement of the concentration of the most active metal in other engneering- alloy systems by a similar technique is also considered. IN recent years there have been many studies concerning the direct electrochemical measurement of oxygen content in molten engineering metals.&apos; These investigations employ galvanic cells in which the electrolytes are various solid mixed-oxides which conduct current by movement of oxide ion vacancies under the influence of an oxygen potential gradient. Similar cells have also been used to measure the thermodynamic properties of mixing of solid alloys,&apos; one component of which has a much greater affinity for oxygen than the other component(s). One alloy system of obvious engineering importance comprised of two metals of greatly differing affinity for oxygen is copper-zinc. At the casting temperature of molten brasses, the zinc concentration is very difficult to control not only because its vapor pressure is very high, but also because of its high affinity for oxygen to form ZnO. Inasmuch as brass castings may be required to have a zinc concentration within very narrow limits, it would be advantageous for the brass industry to have a means for quickly measuring the zinc content in the molten alloy just prior to pouring. A calculation based on earlier thermodynamic studies of oxygen and zinc in dilute solution in binary Cu-O and CU-Z&apos; alloys indicates that the oxygen concentration of molten 70 wt pct Cu-30 wt pct Zn brass must be much less than 1 ppm for the separate phase ZnO not to exist in the ternary Cu-Zn-O system. Thus it can be assumed that all molten brasses, both in the laboratory and in the foundry, are sufficiently saturated with oxygen for the separate phase ZnO to be present. In view of the foregoing reasons, the cell investigated at 995°C for the purpose of measuring the zinc concentration in molten copper-zinc alloys, particularly those of brass compositions. In this cell the left hand electrode is a reference electrode comprised of a mixture of the pure powders of nickel and nickel oxide, to which a platinum contact is made. The electrode of interest is a molten Cu-Zn alloy of unknown zinc concentration with a tantalum wire contact. The electrolyte is the commercially available (Zircoa) calcia stabilized zirconia. The ZnO in the Cu-Zn electrode compartment is not intentionally added, but is naturally present as a result of the infinitesimal amount of oxygen required for its formation in this alloy system. In cell [A] nickel oxide is reduced at the cathode where O= represents the electrolyte. At the anode dissolved zinc of the Cu-Zn alloy is oxidized to form ZnO(s) in its standard state Thus the overall reaction is for which the molar free energy change is where 5 is the Faraday equivalent (23,063 cal per v equivalent), is the electromotive force of the cell in volts after correction for the Pt-Ta thermocouple, FOf (ZO) and AFF (NO) are the standard molar free energies of formation of the respective oxides, and a2n is the thermodynamic activity of zinc in the molten Cu-Zn alloy. The reference state of zinc in this case is taken to be the pure liquid. At constant temperature all the terms of Eq. [I] are constant except and azn Thus since all other participants in the electrochemical reaction are in their standard states, the change in the electromotive force of Cell [A] is represented only by a change of the thermodynamic activity of zinc at constant temperature. The zinc concentration is related to the activity by sampling the alloy for chemical analysis after cell measurements are taken. EXPERIMENTAL A typical cell construction is shown in Fig. 1. This and other experimental details were similar to that described earlier with the exception of some minor modifications which are detailed below.

Jan 1, 1970

By J. Szekely, V. Stanek

Heat and gas flow interactions in nonisothermal packed beds were examined in the paper. A study was made of the response of the system to a step or pulse change in the inlet gas temperature, with a fixed value of the overall pressure drop across the bed. Under these conditions the heat flow and fluid flow transients are coupled because the convective heat transfer rate is alfected by the mass flow rate of the gas which, in turn, for a fixed pressure drop will depend on the temperature profile in the system. The Problem was stated in terms of differential equations, which were integrated numerically. The parameters used in the computation were so chosen to reflect the operation of the stack region of blast furnaces and sinterbeds. The computed results showed that the pressure drop-flou1 relationships and the rate of progress of the temperature fronts (waves) were strongly interrelated and were also markedly affected by variations in particle size and in void fraction. MANY metallurgical processes involve the contact of solid particles with a gas stream in a packed bed arrangement under nonisothermal conditions. Examples are shaft furnaces, regenerators, pebble heaters, and the sinter process. Although the heat and mass transfer aspects of packed bed processing have been extensively studied, the associated flow phenomena and in particular the effects of nonisothermality on the flow, have received little attention to date. :In fact, most of the previous workers considered the flow rate of the carrier gas time invariant and independent of the heat or mass transfer processSzekely and carr2 have shown that the pressure drop-flow relationships in nonisothermal packed beds may be strongly affected by the temperature profile within the system. These authors have devised a method for relating the pressure drop to flow for a given temperature profile. thus providing an extension of the Ergun equation to steady nonisothermal systems. Finally, in a recent paper willmot3 examined the response of regenerators to externally imposed changes in the gas flow rate. Both these investigations were somewhat specialized since, in general, neither the temperature distribution nor the gas flow rate may be regarded as truly independent variables, thus the heat transfer and the fluid flow problems have to be solved simultaneously. The work reported here is divided into two parts: in Part 1 the general formulation is given and a detailed discussion is presented of fixed bed systems. Part is devoted to systems in which a gas and particulate solids are contacted in a counter flow arrangement. FORMULATION Consider a nonisothermal fixed packed bed, extending from through which a gas is flowing. The problem is to develop a relationship between pressure drop and gas flow (with the gas flow being the dependent variable) and to evaluate the rate at which a change in the gas temperature, imposed at the inlet, travels through the system. Let us consider an infinitesimal section of the bed and establish differential heat, mass, and mechanical energy balances over this region. In addition, relationships are needed describing the temperature dependence of gas density (equation of state) and the temperature dependence of viscosity. These may be written as: where c, and are constants. It is noted that in this formulation we neglected any temperature gradients within the solid particles, thus assuming these to be either sufficiently small or of sufficiently high thermal conductivity to meet this condition. The justification and consequences of this assumption will be mentioned in the discussion.

Jan 1, 1970

By J. Szekely, V. Stanek

Steady state heat transfer and fluid flow rates are predicted for the counter-current contacting of gas and solid streams. For a .fixed pressure drop across the bed the convective heat transfer rate is affected by the gas flow rate, which in turn is affected by the temperature distribution in the system. Variations in bed voidage and in particle size were found to change the predicted temperature profiles quite markedly, in particular when the thermal capacities of the gas and solid streams were equal. IN Part I we examined the effect of step and pulse changes in the inlet gas temperature on transients in the flow of gases through fixed packed beds. Here we shall discuss heat flow and fluid flow interactions in the counter-current heat exchange between packed solids and a gas stream. Counter-current heat exchange in packed beds has been extensively studied. In his now classical paper, Furnas&apos; assumed constant property values and developed a simple analytical. expression for the temperature profiles in the gas and in the solid stream. A highly sophisticated analysis of counterflow systems, allowing for transients and interparticle conduction was developed by Amundsen. In preparing a very readable review of this topic Elliott&apos; also presented a number of elegant solutions, including the case involving heat generation. A great deal of work has been done on the application of these concepts to the modelling of the blast furnace. However, in all this previous work the mass flow of the gas was regarded as a quantity independent of the gas-solid heat exchange process. In the following we shall consider counter-current heat exchange between a gas stream and a packed bed of solids, under conditions where the pressure drop across the bed is specified, rather than the mass flow of the gas. FORMULATION Consider a packed bed, extending from y&apos; = 0 to y&apos; = L, in which moving solids and a moving gas are contacted in a counterflow arrangement. Let cold solids, with an initial temperature Tsi and mass flow GA, enter at y&apos; = L, and hot gas with an initial temperature of Tgi and mass flow Gg enter at y&apos; = 0. In contrast to the problems considered in Part I this system has a nontrivial steady state and we shall restrict our attention to the study of the pressure drop flow relationships and the resultant temperature distributions for steady state conditions. Under these conditions the formulation of the problem requires the statement of the following: 1) heat balance on the gas 2) heat balance on the solids 3) the mechanical energy balance equation and in addition the knowledge of the equation of state, and a functional relationship between viscosity and temperature and an expression for the heat transfer coefficient in terms of the other parameters. On neglecting temperature gradients within the solids and for a constant mass flow of the gas, the solution to this problem is well known.&apos;" The expressions for the temperature profiles of the gas and solid streams may be put in the following form: where Tg and T, denote the dimensionless gas and solid temperatures respectively, y is a dimensionless distance, and R is the ratio of the thermal masses. These quantities are defined as follows: For the particular case of equal thermal mass capacities by the gas and the solid, that is R = 1, the Eqs. [1]and[2] simplify to Eqs. (71 and[8], respectively: i.e., the resultant profiles are linear. As for most practical cases , solutions given by Eqs. [7] and [8] will contain only the inlet temperatures as parameters. For a fixed, predetermined value of C;b the above

Jan 1, 1970

By M. Marciano, K. Kreider

Silicon carbide coated boron fiber (Borsic) reinforced aluminum composites were made which exhibit strength and modulus values predicted by the rule of mixtures. A successful technique for fabricating these composites consists of plasma spraying of monolayer filament reinforced tapes and subsequent diffusion bonding of these tapes into composites. The elastic properties of the composites were given by the rule of mixtures in reinforced directions. (The modulus parallel to the fibers was 32 x 106 psi for a composite with 50 pct by volume fibers. The other elastic properties for a composite with 50 pct fiber were transverse modulus = 14 to 16 X 10 psi, G = 9.5 x 106psi, and v,, = 0.23. Tensile strengths of over 190,000 psi were measured for the composites with 50 pct fiber and the average strength for these composites was 162,000 psi. These same composites had interlaminar shear strengths of 12 to 15,000 psi. Multidirectionally reinforced composites demonstrated the same reinforcement efficiency in strength and modulus as the uni-directimlly reinforced composites. Comrpessive strengths of composites with 50 pct by volume fiber were found to be greater than 250,000 psi. STRUCTURAL filamentary composite materials are being considered for use where high strength, high modulus, and lightweight materials are required. Metal matrix composites have been considered for use at temperatures ranging from cryogenic (aluminum stainless steel) to 6000°F (plasma sprayed tungsten plus tungsten fiber) and other requirements have been similarly diverse. Metal matrix composites offer high modulus and high strength in unreinforced directions, the ability to be joined by brazing or welding, and greater fracture toughness compared with the more easily fabricated resin matrix composites. This paper deals with the mechanical properties of aluminum boron composites, a system which is being considered for applications where lightweight, high modulus, and high strength are required at temperatures including those above the normal operating temperatures of resin composites. The boron fiber used in this investigation had an average room temperature strength of greater than 400,000 psi, a modulus of 58 x 106 psi, and a density of 0.09 lb per cu in. (2.6 g per cu cm).&apos; The fiber used in the composites had a coating of silicon carbide to retard degrading reactions with the matrix at elevated temperatures. The coated fiber, which is sold commercially as Borsic: has been reported to have excellent resistance to degradation at temperatures up to 600°C in air and in aluminum2 and the same mechanical properties as the standard bare boron fiber. Fabrication of Composites. The composites were diffusion bonded from plasma sprayed monolayer tapes. These tapes were fabricated on a substrate of aluminum alloy foil which is incorporated in the composite matrix. The technique consists of winding a layer of fiber on the foil, and plasma spraying the balance of the matrix alloy over the windings which bonds all of the components together. Advantages of this process include the ability to achieve good fiber spacing, Fig. 1, the formation of a good bond between fiber and matrix (achieved during the plasma spraying), and the adaptability of the process for forming large and complex parts. This is particularly true when used with brazing foil backing which alloys low pressure bonding of the structural shape. PROPERTY EVALUATION TECHNIQUES Tensile test specimens, 5 to 7 in. long with a reduced gage section and ,030 to ,040 in. thick, were cut from composite sheets with a diamond saw. The gage section was .200 in. wide, while the gripped section was .240 in. wide and over 2 in. long in each grip. The standard gage length used was 1 in., however, transverse specimens (i.e., specimens with the fibers oriented 90 deg to the loading axis) were tested without a reduced section. Twelve specimens were also made with a 3-in. gage length but no significant effect of gage length was measured. The tensile test specimens were mounted in a fixture and aligned in friction grips using a 10X microscope. Unbonded foil doublers (0.010-in.-thick aluminum) were used to insure firm gripping with a minimum of damage to the fibers. The test specimens were transferred to the testing machine in a rigid fixture and mounted in the self aligning loading train. Testing was performed using a Tinius Olsen four screw testing machine with a torsion bar LVDT load

Jan 1, 1970

By R. Wang, N. J. Grant, B. C. Giessen

By crystallographic and X-ray methods, the existence and isonzorphism of Ti3Rh5 and Hf3Rhs were confirmed. Both phases are of the orthorhombic Ge3Rh5 type; lattice parameters and refined positional parameters are given. The structure is related both to the filled-up NiAs-B8 and Cu-AI types. An analogous phase with zirconium does not exist; the effect of ternary substitutions for titanium ad hafnium suggests a size factor limit to be active. A recent survey of phase diagrams of the T4 metals titanium, zirconium, and hafnium with the T, noble metals rhodium and iridium indicated the existence of the A3B5 phases Ti3Rhs, ZrsRhs, and HfsRhs. Ti3Rhs and Hf3Rh5 were found to be isostructural, based on the line-rich powder patterns which had not been analyzed. Zr3Rh5 was considered to have a substructure of the NbRu type (orthorhombically distorted B2-CsCl type).&apos; Because, in combinations with other transition metals, hafnium and zirconium are generally more likely to form isostructural phases than hafnium and titanium (with the significant exception of the Ti2Ni-"E93" type phases based on T4 metals2), the reversal of this relation for the A3B5 phases was of interest. As shown in the following, the nonexistence of Zr3Rhs has been established, the structures of Ti3Rh5 and Hf3Rh5 have been worked out, and crystal chemical relationships and stability criteria are reported. EXPERIMENTAL METHODS AND RESULTS Alloy Preparation and Phase Diagram Work. Alloys were prepared from high-purity (99.99+ pct) elements by arc-meltin3,4.Heat-treated alloys were annealed in a vacuum of 3 x X torr for 24 hr at 1300DC. Metal-lographic samples were etched electrolytically in concentrated HCl with 5 v ac for 5 min.3 X-ray diffraction powder patterns were taken on a GE XRD-5 dif-fractometer with Cum radiation at low scanning rates (0.2 deg per min for 28). It was confirmed that Ti3Rhs and Hf3Rh5 have similar diffraction patterns, and that an alloy with the composition Zr3Rhs has a different pattern. Six Zr-Fh alloys with 59 to 69 at. pct Rh were therefore prepared and investigated in the as-cast state by X-ray diffraction and metallography. Alloys at 59 and 61 at. pct Rh were found to be a single phase, with the distorted B2-CsC1 type structure typical for the off-stoichio- metric region of the phase (Zr,-,Rh,)Rh. This phase forms a eutectic with ZrRhs at about 66 at. pct Rh: accordingly, alloys between 61 and 69 pct at. pct Rh consisted of two phases. There is no evidence for the existence of Zr3Rh5. Based on the results in Rafs. 1 and 5, on the present work on Zr-Rh, and on several additional alloys investigated, the portions between the AB and AB3 stoi-chiometry for Ti-Rh, Zr-Rh, and Hf-Rh are as follows: Further, several ternary alloys near Ti3Rhs and Hf,Rhs were prepared in which it was attempted to replace titanium and hafnium partly by zirconium, niobium, tantalum, and germanium. The results will be discussed in a later section. Structure Determination of Ti3Rh5. Since Ti3Rh5 and Hf3Rh5 are isostructural, the following discussion will largely deal with the former. Although the powder pattern of TisRhs is complex, as found previously,1 it could be indexed by comparison with other structures of A3Bs stoichiometry. Ti3Rh5 was found to be isostructural with Ge3Rh5, whose orthorhombic structure had been elucidated by Geller.9 As both the sizes and atomic numbers of germanium and titanium are comparable, the unit cell volume and the peak intensities could be expected to be similar; however, significant differences exist in the atomic positions, as will be shown. All lines in the powder patterns of Ti3Rh5 and Hf3Rhj could be indexed with primitive orthorhombic unit cells with the lattice constants: The fractional errors are 10 The low-angle portion of the indexed powder pattern of Ti3Rh with sin2 8 < 0.30 is listed in Table I. The extinction laws Okl only with k = 2n and hOl only with h - 2n are compatible with the space group Pbo2 and the more symmetrical space group Phnm of Ge3Rh5. Finally, the positional parameters of Ti3Rh5 and HfsRhs were refined under the assumption that titanium and hafnium occupy the germanium positions in Ge3Rh5. Integrated intensities were obtained from the diffraction patterns by planimetry. Intensities of overlapping reflections were separated by an iteration process incorporated into the least-squares positional refinement program according to a method described previously. The intensities of Ge3Rh5 were used in the first separation cycle, while the atomic parameters of Ge3Rh5 were used as starting values in the first refinement cycle. Absorption due to specimen

Jan 1, 1970

By Yuri D. Tretyakov, Robert A. Rapp

The stoichiometry ranges ofNiOl+y and LiFe,O,-d were established by high-temperatwe electrochemical meas7rements in a stabilized-zirconia electrolyte cell. The results were consistent with doubly ionized cation vacancies in NiO,+y and interstitial lithium or iron ions in . The defect structure of the ternary ferrite was derived from the consideration of equilibration with respect to oxygen between the solid and the gas phase. The absolute magnitudes of defect concentrations were calculated. Pavtial molar enthalpies of oxygen in the compounds were calculated and interpreted in terms of the enthalpy of defect formation in these crystals. NICKEL oxide (NiO,,?) is a metal-deficient, p-type, extrinsic semiconductor whose properties are consistent with a structural model based on the presence of cation vacancies as the predominant ionic defect at sufficiently high temperatures and oxygen activities. A survey of previously reported conductivity studies and the presentation of some more recent conductivity measurements will be given in a later paper.&apos; The absolute magnitude of the equilibrium vacancy concentration in NiO has been reported from combined conductivity and thermogravimetric data of &apos; However, disagreement exists concerning the state of ionization of the nickel vacancies in NiO. Some authors3-&apos; have proposed that the predominant defects in NiO are singly ionized nickel vacancies and positive holes (h&apos;) formed by the reaction where, accorhng to the notation of Kroger and Vink,6&apos;7 Oq represents an oxygen ion on its normal lattice site. Other authors"27E have proposed that doubly ionized nickel vacancies (V{i ) and positive holes are predominant and are formed by the reaction One purpose of the present investigation was to establish the nature of the predominant defect in NiO,+? as well as its equilibrium concentration and thermo-dynamic properties at elevated temperatures and known oxygen activities. To accomplish this purpose, the coulometric titration of oxygen into and out of NiO was accomplished using a galvanic cell involving the calcia-stabilized zirconia electrolyte. The ternary oxide, lithium ferrite (LiFe50,-6) is ferromagnetic and has the inverse spinel structure LiFe,O,-d Thus, in the ideal stoichiometric XY204 lattice, Fe3 ions occupy one-eighth of the tetrahedral (A) sites, and i&apos; and Fe3&apos; ions share at random one-half of the octahedral (B) sites within the fcc sublattice of oxygen ions.g The structure and thermodynamics of the spinel structures have been comprehensively decribed.&apos;-&apos; A recent compilation of literature for lithium ferrite is also available. 15 Oxygen-excess LiFe50sis not expected to exist because both i&apos; and 17e3& ions are in their highest normal valency states (positive hole formation is not favorable). As will be discussed, equilibration of the ferrite crystal with oxygen of a surrounding gaseous phase will result in the introduction of equilibrium concentrations of ionic and electronic defects. In the Results and Discussion section a defect model for LiFe50,-d is proposed. This defect model is tested by high-temperature coulometric titration experiments. EXPERIMENT The oxygen activities in nonstoichiometric NiOl+, (and also LiFe,O,-d) were measured by means of coulometric titration with the high-temperature galvanic cell The critical characteristic of a suitable experimental cell is the complete isolation of the phase to be investigated in a minimum sized chamber which is free from extraneous sources and sinks for oxygen (leakage). Then oxygen is only admitted to the chamber or removed from it in known amounts by coulometric titration, which involves the passage of oxygen ions through the solid electrolyte with electrochemical oxidation and reduction reactions at the platinum contacts to the electrolyte. The experimental cell is shown in Fig. 1. The cap of the cell was the tip from a closed-end alumina tube, which was found to be leak-free from a helium leak-detector test. This alumina cap was about 1.2 cm OD and about 1.5 cm high. A Zircoa calcia-stabilized zirconia tablet (crucible lid) of 1.5 cm diam and 0.3 cm thickness served as the solid electrolyte. A Pyrex ring of about 0.1 cm thickness was placed between the electrolyte and the cap. The electrolyte tablet was painted with platinum paste on the entire outer face and on that part of the inner face which would be within the enclosure; these electrodes were further prepared by heating in air at 1000°C for 2 hr. Nickel oxide powder, listed as 99.999 pct pure, was purchased from Leico Industries, Inc. Pills of the NiO were cold-pressed and sintered at 1050°C for 3 hr in a platinum crucible. This poorly sintered NiO was crushed, and chunks were wrapped in 52-mesh Pt gauze (to catalyze the solid-gas exchange) for place-

Jan 1, 1970

By B. S. Hickman

Using primarily quantitative single crystal X-ray techniques studies have been made of the precipitation of the metastable w phase in alloys of titanium with Mo, Mn, Fe, Cr, and Nb. It is shown that, in agreement with earlier work on the Ti-V system, the composition of the w phase during aging approaches a constant; the value varies from one system to another in a systematic manner with the electron concentmtion. Results relevant to the mechanism of w precipitation during quenching and during aging are presented and discussed. Results are presented to confirm that the w phase is coherent and the nzorphology of the Necipitates is described and discussed. PRECIPITATION of the metastable w phase in alloys of titanium is of interest because of its severe embrittling effect and also because of the enhancement of the superconducting properties which accompanies precipitation.1 The general features of w phase precipitation have been discussed by cuillan and by Bagaratskii et d3 The work described in this paper forms part of a general study of w phase precipitation which was undertaken to elucidate various unknown or controversial features of the process, namely: i) What is the mechanism of formation of the phase ii) How does its composition, morphology, volume fraction, and particle size vary with the type and content of the alloying element and with the heat treatment conditions? iii) What is the relation between the w phase and the equilibrium a phase? In a previous paper4 the results of a detailed study of w phase in Ti-V alloys were presented. It was shown that: U During aging the composition of the w phase approached a saturation value of 13.5 to 14.0 at. pct V. ii) w precipitated as cube shaped particles and once precipitation was complete, particle coarsening did not occur. iii) a phase formed initially by direct conversion of w-phase precipitates. In this paper similar measurements are reported on Ti-Mo alloys and some more limited data given on the Ti-Cr, Ti-Fe, Ti-Mn, and Ti-Nb systems. The results i.1 the six systems are then compared. It has been shown previously by electron microscopy that the w precipitates as ellipsoids in both the Ti-Nb system,&apos; and in i-o.&apos; Blackburn and williamsz3 have reported the precipitates in Ti-Fe, Ti-Cr, and Ti-Mn are cubic, i.e., similar to the Ti-V system.476 EXPERIMENTAL METHODS 1) Specimen Preparation. Samples of the compositions shown in Table I were prepared by arc melting. High purity electron beam zone refined titanium (Materials Research Corporation Grade I—total impurity content =250 ppm) was used; the other source materials are given in the table. A commercial Ti-8 wt pct Mn alloy was used for studies of this system. After homogenization the materials were rolled or cut into strip approximately 1/4 in. by 2 in.; the strip was then heated at temperatures of 1200" to 1300°C for several hours by passage of a direct current through the strip in an ultra high vacuum system (<10-torr) and then quenched by admitting high purity helium and switching off the current. This treatment resulted in crystals up to 2 in. by 2 in. by A in., almost all of which had a [100] direction nearly perpendicular to the plane of the strip. The crystals were cut from the strip by spark machining and their composition checked by lattice parameter measurements supported by chemical analysis. 2) X-Ray Diffraction Methods. Lattice parameter measurements were made using a modification of the single crystal spectrometer described by ond. In this technique the crystal angles for reflection on either side of a highly collimated primary beam are measured. Generally the 400 CuK, or the 400 CuKp reflections of the bcc lattice were used. In the as-quenched crystals a precision of 1 in 105 was obtained, but due to line broadening the parameters in aged materials could only be measured to 1 in lo4. This accuracy was, however, quite adequate for the purposes of these experiments and enabled the composition of the p phase to be obtained to 10.1 pct in all systems except Ti-Nb where the lattice parameter change with composition is very small. The precipitation and growth of the w phase was examined using a single crystal goniometer attached to a Phillips diffractometer. The relative intensities and line shapes of b-phase and w-phase reflections

Jan 1, 1970

By J. Szekely, J. W. Evans

A formulation is given for radiant heat loss from the surface of molten steel held in a ladle and numerical solutions are presented for the resultant integrodif-ferential equations. The results are shown on plots of the net rate of heat loss against time, with the geometlry, emissivities, thermal properties of the walls, and Preheat as parameters. It was found that the rate of the vadiant heat loss from the exposed metal surface was comparable to the conductive losses from the Portion of the metal in direct contact with the walls. The radiant heat loss was markedly affected by the geometry and by the thermal properties of the walls the net radiant heat loss from the steel could be drastically reduced by providing a top cover and also by the provision of preheat. The results of this work are thought to be relevant to heat loss calculatious concerned with vacuum degassing and with operatzons in which molten metals are held in ladles. WhEN molten steel is held or transferred in ladles heat losses inevitably occur. The accurate assessment of these heat losses is of importance in process design considerations, as the temperature of the metal may have to meet rigid specifications in certain stages of the processing sequence. As shown in Fig. 1, the principal mechanisms responsible for heat loss are: i) Conduction into the walls and ii) Thermal radation from the top metal surface (possibly augmented by natural convection) The conductive heat loss may be readily calculated as much information is available in the literature on both the general theory1&apos;&apos; and on its particular applications to ladle heat loss calculations. Evaluation of the radiant component of the heat loss is less straight-forward. In two recent papers it has been shown that the presence of a slag layer 2 to 3 in. thick would prevent any significant heat loss for periods up to an hour,5&apos;" however, should bare metal be exposed, constituting the radiant surface, substantial heat losses could occur. This latter case is of considerable practical interest in various degassing operations and in the vacuum processing of certain specialty steels, where the top surface of the metal is not covered by a protective slag layer. Under these conditions a substantial radiant flux is emitted by the top metal surface; part of this flux leaves the system, part of it falls in the ladle walls. The radiant flux incident on the ladle walls is in part absorbed and conducted away, and in part reflected and re radiated. The purpose of the paper is to present a general formulation of the problem, the resultant solution of which should enable computation of the net radiant heat loss from the metal surface. The principal variables, the effect of which will be investigated, include initial metal temperature, geometry, thermal and radiative properties of the wall, and various modes of preheat. FORMULATION Consider a cylindrical enclosure, shown in Fig. 2, with a diameter D, height H, and wall thickness W, the cover at the top of the cylinder having a wall thickness L. If the temperature of the base (representing the top surface of the metal) is kept constant, say, Tb, the problem is to calculate the net rate of heat loss from the base as a function of time, geometry, the initial temperature distribution, and the thermal properties of the system. As customary, the formulation will consist of the mathematical statement of the conservation of thermal energy within the system.

Jan 1, 1970

By K. M. Kolobova, E. Z. Kurmayev, A. Z. Menshikov, S. A. Nemnonov, V. A. Trapeznikov

The present work represents a review of our intestigations into the X-ray spectra of solids, which have been carried out by us along two lines of upproach: 1) St@ of the electronic structure of high -rnelting compounds of transition metals; 2) stzddy of the interatomic binding forces by the fine structure of absorption spectra under varying temperature and concentration conditions. With this in mind, int>estigations hare been made into the emission and absorption spectra of the K-series of atoms of transition elements in alloys and compounds. In the problem concerned with the electronic structure of high-melting compounds of transition metals, considerable importance is attached to the question of ascertaining the role played by 3d4sp electrons of metal and 2p electrons of metalloid in the formation of chemical bonds and of giving, on the strength of this data, an explanation of the properties which seem to be contradictory at first sight: good conductivity inherent in metals and high brittleness and high-melting characteristics typical of covalent compounds. Over the last few decades this question has attracted the attention of many investigators whose studies were devoted to this type of compound. Many works, both of theoretical1-3 and experimental4- l5 nature , have contributed much to the elucidation of the physical character of the phenomenon referred to. Alongside the investigations in the field of electric resistance, Hall effect, paramagnetic susceptibility, electronic heat capacity, and the temperature of transformation into a superconducting state, we should particularly emphasize the studies on X-ray spectra, which have played a significant role in the development of the zonal theory for high-melting compounds. This fact has become highly evident after the theoretical calculation of the energy spectrum has been made and the density of states curve in Tic, TiN, and Ti0 has been plotted by Ern and Switendick3 who used the method of augmented plane waves. Correlation of the results of the theoretical calculation of the band structure in the compounds listed above with the data obtained by X-ray spectroscopy studies showed that X-ray spectra can be a source of valuable information on the electronic structure of high-melting compounds within a wide range of ener- gies. It is well known that X-ray emission spectra basically give a direct picture of the occupied excited states above the Fermi level of a solid (valence band, conduction band) and that the absorption spectra (region of initial absorption) give a similar picture of states corresponding to the unoccupied levels near the Fermi surface. Full information on the density distribution of electronic states by energies can be obtained by combining the emission and absorption spectra. Such an opportunity arises from investigation

Jan 1, 1970