Search Documents

By L. D. Jaffe, D. C. Buffum

EARLIER the authors and coworkers had pre sented data on isothermal temper embrittlement of an SAE 3140 steel?' In that work, however, attention was concentrated on embrittlement at 575°C and below. Preliminary measurements showed that embrittlement at 675 °C was much mote rapid than would be expected from the data at the lower temperatures: and indicated the need for further work above 575°C. The possibility of an upper nose in the time-temperature embrittlement diagram above 575°C was pointed out. Procedure The same heat of SAE 3140 steel was used as in the previous work, and experimental methods were the same. All blanks were austenitized 1 hr at 900°C and water, quenched. Groups were then tempered 1 hr at 675"C, water quenched, and given embrittling treatments for times of 5 min to 240 hr at 675" to 600 °C, as shown in Table I. Other groups, after the temper and quench, were embrittled 48 to 1440 hr at 550°C. Another series was tempered 240 hr at 675"C, quenched, and held for various times at 650" to 375°C (Table 11). For comparison, a series was tempered for these same times at 700°, 675", and 650°C (Table 111). The Ae, temperature for this heat of steel had been determined as approximately 690°C;1*a the treatments at 700°C were introduced to determine the effects of tempering in a range where a little austenite would fonn. A large number of other treatments, involving holding at two temperatures in succession after the temper, and occasionally slow cooling or heating, were also studied. These treatments are listed in the tables. Austenitizing was done in air; all subsequent treatments were in salt. Blanks were always water quenched after the last treatment. The heat-treated blanks were machined to standard V-notch Charpy specimens. A transition curve for each group was obtained by breaking the specimens at various temperatures on a 217 ft-lb impact machine having a striking velocity of 16.8 fps, and plotting the observed percentage of fibrous fracture vs the test temperature. From this plot, the transition temperature was taken as the lowest temperature at which the fracture appeared 100 pct fibrous; its accuracy is estimated as +5"C. Generally, hardness was measured on every bar. Since a decrease in hardness itself seems to decrease the transition temperature somewhat, a correctiona of —0.28"C per Bhn was applied to permit comparison of the transition temperatures on a uniform basis of 24 Rc. The transition temperature and hardness results are given in Tables I through VI. Where overlap of earlier work1.' occurred, data were combined.

Jan 1, 1958

By Leslie Burris, J. C. Hesson

The diffusivity of uranium in liquid cadmium has been measured as a function of temperature by the capillary -bath technique. The diffusivity measurements were made at 4509 500°, 575º, and 650°C, over which temperature range the diffusivity of uranium varied from 1.3 to 2.3 x 10-5 sq cm per sec. The diffusivity of uranium in cadmium was less than that predicted by the Einstein-Stokes equation. VERY little information is available on the diffusivity of uranium in liquid metals. Bonillal has reported a single measurement of the diffusivity of uranium in bismuth. Measurement of the diffusivity of uranium in cadmium was undertaken, not only to supply data in an area in which a dearth of information exists, but also because of a need for this physical property measurement in connection with mass transfer studies now underway in the U-Cd system at Argonne National Laboratory. The capillary-bath method used in this work has been successfully used by other workers2"4 to obtain diffusivities in liquid metal systems. Although considerable experimental technique is required to obtain reliable diffusivity measurements, the basic information required is only that of temperature, capillary tube length, the initial and final compositions of the alloy in the capillary tube, and the composition of the bath into which the solute element is diffusing. From these measurements, diffusivities may be calculated using Fick's law. The diffusivity equations are presented in the Appendix. Equations which are often used to predict diffusivities in dilute solutions5 are the Einstein-Stokes equation: D= (kT)/(6pµr) and the Eyring equation: D = (?,kT/?y?zp), where D= diffusivity (sq cm per sec), k = Boltzman's constant, T= temperature ("K), p = viscosity (poises), r = radius of diffusing molecule or species (cm), and A,, Ay, and A, are the diameters of the diffusing molecule in the three axial directions. For round molecules, the Eyring equation gives values which are 37r times larger than the Einstein-Stokes equation. In this work, the experimental values were compared to the Einstein-Stokes equation. Experimental values of diffusivities in liquid metal systems usually lie between those predicted by the Einstein-Stokes and Eyring equations. The viscosity values for cadmium were obtained from the Liquid Metals Handbook.6 EXPERIMENTAL APPARATUS AND PROCEDURE The general procedure employed consisted of rapidly immersing in a molten cadmium bath containing a negligible amount of uranium a capillary

Jan 1, 1963

By R. J. Teitel

FOUR or five years ago an investigation was initiated to study and develop basic information on llquid metal fuels for future power nuclear reactors. The objective was to find fluid forms of uranium (solutions or dispersions) in low melting alloys. A summary of the different fuels was presented in a recent article by Teitel, Gurinsky, and Bryner.' The combination of uranium and bismuth, with its low thermal neutron cross section and low melting point, is of basic interest to the development of these fuels. Most of the previous effort on the U-Bi phase diagram was done under security regulation. Ahmann and Baldwin' reported their findings on the U-Bi system in 1945. Their results indicated that the phase diagram had at least two compounds and possibly one more. UBi, and UBi were established by chemical analyses of residues after high temperature centrifugation of alloys. UBi, decomposed peritectically at1980°C and there were signs of another reaction above that. UBi melted at a very high temperature, and there were also signs of a monotectic at high temperatures in high uranium alloys. The solubility of uranium in liquid bismuth has received attention by three groups of authors: Bareis,X ayes and Gordon,' and Ahmann and Baldwin.' The latest and most widely accepted data is that of Bareis, and it is included in Fig. 1 along with new data at higher temperatures. Brewer et al." and Ferro"'' have used X-ray diffraction to determine the number and crystal structure of the compounds. Brewer identified the structure of UBi to be a simple NaCl cubic structure having a cell edge of 6.364A. The density of the compound was calculated to be 11.52 g per cu cm. Two other unindexed patterns were presented as possibly belonging to two other compounds. Ferro agreed with Brewer on the crystal structure of UBi and identified the structure of UBi, and U,Bi,. The intermetallic compound UBi, has a tetragonal cell with a = 4.445A and c = 8.908A. The density of the compound was determined and the number of atoms per cell (six atoms) was calculated. The X-ray density reported was 12.35 g per cu cm. The compound UsBir was treated in a similar way and it was reported to have a cubic structure 9.35A on an edge and to contain 28 atoms per cell. The density was given as 12.59 g per cu cm. Teitel reported that there was a compound UaBi, and no UBi, compound. (The difference in the compositions is 3.3 pct U.) This conclusion was based upon the extrapolation of some 20 tie lines in the

Jan 1, 1958

By A. H. Daane, A. S. Wilson

The U-Cr system is of the simple eutectic type with some solid solubility of chromium in r and ß uranium. The eutectic occurs at 20 atomic pet Cr and melts at 859°C. The maximum solubility of chromium in y uranium is 4 atomic pet at the eutectic temperature, and in ß uranium the solubility is estimated to be 1 atomic pet. y-ß and ß-a transformations were found to occur at 737°estimatedt° and 612°C respectively. DURING 1944 and 1945, the U-Cr constitution diagram was studied in this laboratory as a part of a research program on uranium metallurgy in the Manhattan Project, and the work was described in a Manhattan Project report issued in December 1945. This paper is based on that report, which has been declassified. Prior to this study, it had been shown by other Manhattan Project workers that the low Cr-U alloys could be quenched to retain the form of uranium. Experimental The uranium used in this work was massive metal prepared in this laboratory and contained less than 0.1 pct of other elements. The chromium was 200 mesh powder obtained from the A. D. McKay Co. and was found on analysis to be 99.5 pct Cr with 0.3 pct Fe the major impurity. Alloys, weighing 400 to 600 g, were prepared by induction heating the components to 1700°C in slip-cast ZrO, crucibles in a vacuum of 3x103 mm Hg. TO prevent too violent agitation of the melt by the induction field with subsequent crucible breakage and sample loss, the ZrO2 crucible was placed in a graphite crucible, which was surrounded by a layer of powdered carbon insulation 2 to 3 cm thick. Polished vertical sections of the alloys were examined microscopically to confirm their homogeneity. Heating and cooling curves were taken on the alloys by reheating them in ZrO, crucibles to 1200°C and inserting a mullite-protected chromel-alumel thermocouple into the melt by means of a slip seal in the vacuum head of the furnace. A recording potentiometer traced the curves which had a slope of from 3" to 6" per min. Samples of the alloys were prepared for metallo-graphic examination by conventional mechanical polishing techniques followed by an electrolytic polish in an ethylene glycol-phosphoric acid-ethyl alcohol bath. The structure of the alloys was brought out clearly by this procedure so that no further etching was required. Samples for chemical analysis were taken from drillings from the top, center, and bottom sections of the alloys. The uranium was determined by titra-tion with Ce(SO1)2, while the chromium was titrated with FeSO,; the uranium and chromium totaled at least 99.6 pct in all of the alloys prepared. X-ray samples were prepared by filing bulk specimens in a helium-filled glove box and annealing the resulting powder in a zirconium-gettered helium atmosphere. A 114.6 mm diam Debye-Scherrer camera and a Weyland nonsymmetrical self-focusing camera were used with filtered copper radiation to obtain the powder X-ray diffraction data. Results The data obtained in this study have been combined to construct the constitution diagram of the U-Cr system shown in Fig. 1 where the arrests observed in cooling curves are indicated by dots. The liquidus arrest was quite distinct in thermal data taken on alloys in the range 0 to 20 atomic pct Cr. The eutectic arrest was not observed in studies on the 2.5 and 4.5 pct Cr samples but appeared in the 7.5 pct samples, which suggested some solubility of chromium in y uranium. On quenching from 859 °C, the 2.5 pct sample showed but one phase while the 4.5 pct sample contained a small amount of the eutectic along the grain boundaries; see Figs. 2 and 3. From this the maximum solubility of chromium in r uranium has been set at 4 pct. X-ray studies on these samples showed that the r phase was not retained at room temperature by quenching, but in each case a pattern was observed .which has been identified with the ß phase of uranium. Thermal data show the y-ß transformation of uranium lowered to 737°C as a consequence of this solubility. On quenching from the ß range (660°C), precipitation of chromium in the primary uranium is observed in the 2.5 and 4.5 pct Cr samples (see Figs. 4 and 5),

Jan 1, 1956

By A. Kaufmann, B. Cullity, G. Bitsianes

T0 determine the bulk of the phase diagram, techniques for melting, thermal analysis, heat treatment, metallography, and X-ray diffraction that have already been described were used.' It proved difficult to get definitive results for most of the alloy system because of the large number of compounds, their high melting points, and their extreme brittle-ness which made metallography difficult. The final phase diagram, essentially as shown in Fig. 1, was issued with considerable trepidation.' The phase relationships at the two ends of the system were defined with satisfactory accuracy. In the central region between U,Si, and USi, there was great uncertainty concerning the melting points and even the composition of the compounds. At the uranium end of the system, it is quite clear that a eutectic exists at about 985 °C and approximately 8 or 9 atomic pct Si. The a-to-8 transformation is not appreciably altered by silicon, but the 8-to-y transformation is raised to about 795°C. There is appreciable solid solubility in y-uranium, as demonstrated by the micrographs of Fig. 2. The occurrence of the E phase was not suspected at first, since the alloys as melted and furnace cooled gave no X-ray diffraction or thermal arrest indication of it. However, suitable metallography revealed a rim around each of the compound particles as shown in Fig. 3. Heat treatment above 600°C revealed that the rim was another phase which formed by peritectoid reaction between uranium and U,Si,. This reaction proceeds slowly and it is difficult to get it to go to completion. Consequently, there is some uncertainty concerning the exact composition of P and the exact peritectoid temperature. On the basis of various metallurgical studies, it was decided that the most probable composition for t was about 23 atomic pct Si. Metallographic examination shows the closest approach to a single phase at this composition. Likewise, density and hardness measurements of heat-treated alloys shown in Figs. 9 and 10 show a change in slope at 23 pct Si. This is revealed most clearly by plotting the difference in hardness or density between chill-cast and heat- treated alloys as shown in Fig. 4. Zachariasen2 was able to index the diffraction lines of P and to determine the arrangement of the atoms in the unit cell. He concluded that the composition must be US. However, it is felt that the metallurgical studies are reliable and that Zachariasen's X-ray work does not furnish positive proof that the P phase exists exactly at the stoichiometric composition U3Si. There are many instances in alloy systems where a phase does not occur at the composition dictated by the molecular structure. It is possible that the chemical analyses used in the present work were subject to some systematic error. This seems unlikely, however, since recent work in several laboratories seems to show excess compound at 25 atomic pct Si.

Jan 1, 1958

By M. C. Udy, F. W. Boulger

AN incomplete phase diagram for the U-Ti systern was determined earlier 1 and more recently, a tentative diagram was presented for the uranium-rich end of the system.' In the present re-examination of the whole system of U-Ti alloys, high purity materials were used. Melting stock for the alloys was high purity uranium, containing about 0.09 pct C as the only appreciable impurity, and high purity iodide-process titanium purchased from New Jersey Zinc Co. Both metals were cold rolled to about 1/6 in. thickness, sheared to about I/' in. squares, and cleaned by pickling. The alloys were arc melted under a helium atmosphere in a water-cooled copper crucible. A thoriated-tungsten electrode was used. The furnace chamber was evacuated, then flushed with helium, prior to each melting. It was finally filled with stagnant helium at one atmosphere pressure. Each alloy was remelted three times after the original melting, to insure homogeneity. The alloy button was turned bottom side up before each re-melting operation. Some 22 alloys were examined. Their compositions were spaced at appropriate intervals between 100 pct Ti and 100 pct U. Analyses were made on chips taken after fabrication. The major contaminant was carbon, which varied from 0.03 to 0.08 pct. It appeared in the microstructure as titanium carbide. Alloy compositions were calculated to a carbon-free basis for consideration on the diagram. Tungsten and copper, possible contaminants from the melting operation, were generally less than 100 parts per million each. Fabrication All alloys were forged and rolled to bars approximately V8 in. square. They were clad either in SAE 1020 steel or in a 5 pct Cr-3 pct Al-Ti-base alloy, depending on the fabrication temperature. A temperature of 1800°F (980°C) was used for alloys near the compound composition. This necessitated using the titanium-base alloy, since iron reacts with titanium at this temperature, producing a low melting alloy. Other alloys were fabricated at 1450°F (790°C), using steel jackets. No iron-titanium reaction occurred at this temperature. The jackets were welded in place in an argon atmosphere. Those alloys sheathed in steel were declad and then reclad between rolling and forging operations. On the other hand, those clad with the titanium alloy were cut to a roughly rectangular shape prior to clading and were then carried through both the forging and rolling operations without opening. Those alloys near the compound composition were found to be cracked when the clading was removed. The cracked materials had been plastically deformed, however, and at least some of the cracking had OCcurred during cooling. Heat Treatment The rolled bars, after being declad and shaped to remove surface contamination, were all given an homogenizing treatment of 160 hr at 2000°F. (Samples were taken for analysis following the declading and shaping operations.) All were heat treated at the same time in one furnace, but each was sealed in a purified argon atmosphere in an individual Vycor glass tube. Argon pressure was such that it was approximately atmospheric at temperature. One end of each tube contained titanium chips and this end was heated to 1200°F (650°C) for 10 min prior to the heat treatment. This purged the atmosphere of residual reactive gases. The balance of the tube was warmed during the purge to liberate adsorbed moisture and gases, which also reacted with the hot chips. The bars were furnace cooled from the homogenization treatment. Specimens of each alloy were water quenched after 2 hr heating at 1000°, 1200°, 1400°, 1600°, 1800°, and 2000°F (540°, 650°, 760°, 870°, 980°, and 1095°C). In addition, some were treated at intermediate temperatures of 1300°, 1500°, and 1700°F (705", 815", and 925°C) and at 2150°F (1175°C). Specimens, about '/s in. cubes, were cut from the bars, sealed in individual Vycor tubes, and heat treated as described. All specimens heat treated at the same temperature were processed together. Samples were quenched by breaking the Vycor tube rapidly under water. Metallographic Examination Specimens were mounted in bakelite and ground wet on 180 grit paper held on a 1750 rpm disk. They were then ground wet by hand, using 240, 400, and 600 grit papers. The rough grinding was continued long enough to get well below the surface. Specimens were mounted separately because of the variation in the rate of etching between alloys. The specimens were polished with rouge on a 4 in., 1725 rpm wheel covered with Miracloth. Alloys on the titanium side of the compound composition were etched with a solution of 2 pct hydrofluoric acid in water saturated with oxalic acid. A few crystals of ferric nitrate were added as a bright -ener. Specimens were immersed 5 sec, polished to remove the etch, then re-etched. With the higher titanium alloys, it was often necessary to start the etch on the polishing wheel, because of the formation of a passive film. In some instances, a plain 2 pct hydrofluoric etch was satisfactory. For the alloys on the uranium side of the compound, a distinction between the compound and the uranium phase developed after standing a short time in air. This could be hastened by the application of heat, such as obtained by placing the specimen on a radiator. A deep etch was necessary to develop details in the uranium-rich phase, such as the Widmanstaetten pattern sometimes obtained by quenching y uranium. A 2 pct hydrofluoric acid solution was used for this deep etching.

Jan 1, 1955

By H. H. Klepfer, K. J. Gill, P. Chiotti

SOME observations relative to the U-Zn system have been made by other investigators. Chipman1 and Carter2 have reported the preparation of several U-Zn alloys and point out that these alloys are generally difficult to prepare. Chipman1 reported evidence for a high melting compound at about 90 atomic pct Zn and the possible existence of a eutec-tic between the compound and uranium. Raynor," in a theoretical discussion of the alloying properties of uranium, included zinc among the elements predicted to have little or no solubility in a, p, or y uranium. In the present investigation, thermal analyses, X-ray, metallographic, and vapor-pressure data were obtained to determine the phase boundaries. The relatively high zinc pressure over most of the alloys at temperatures of 900 °C and above proved troublesome and special techniques had to be employed in preparing suitable alloys. Materials and Preparation of Alloys The metals employed in this investigation were Ames Laboratory biscuit uranium containing less than 500 ppm total impurities and Bunker Hill slab zinc or Baker Analyzed reagent granulated zinc, both with a purity of 99.99+ pct. Due to the high vapor pressure of zinc and the high reactivity of both uranium and zinc with oxygen at only moderately high temperatures, alloys were prepared in closed containers which had either been evacuated or evacuated and filled with helium. High purity magnesia, magnesia containing 10 pct calcium fluoride, and tantalum proved to be suitable crucible materials. Tw-o different procedures, described below, were used to prepare alloys, the latter being the most satisfactory. The metals, uranium turnings and granulated zinc, were cleaned with dilute nitric acid, rinsed, dried, and placed immediately in a helium-fill'ed dry box. The two metals were placed' in a 10 mil Ta crucible. The charge was enclosed in the tantalum crucible by welding on a preformed tantalum cap. This assembly was enclosed inside a stainless steel (AISI 309) bomb. The bomb was made by welding a piece of stainless steel plate on each end of a stainless steel pipe. All these operations were carried out in a helium atmosphere. These assemblies were heated in a muffle furnace at temperatures between 1100" and 1200°C for 10 to 15 min or held as long as 15 to 20 hr in the 950" to 1000°C temperature range before quenching. Spectrographic and chemical analyses showed no tantalum pickup by the alloys, indicating no reaction between the alloys and the crucibles. However, some of these crucibles failed, probably due to imperfections in the welds of the stainless steel or tantalum crucibles. The second and most satisfactory method was to prepare the alloys by powder metallurgy techniques. The procedure was to press degreased and acid-etched uranium turnings with granulated zinc into 20 g compacts under 20,000-psi pressure. The compacts were placed in MgO crucibles, and sealed in evacuated Vycor or fused silica tubes. The alloys were then heated as long as two weeks at about 550°C in a muffle furnace. The pressed compacts were observed to expand by several volume percent during heating and it was necessary to make allowances for this expansion in order to avoid breaking the crucible and Vycor tube. This method was found very satisfactory for preparing alloys which were suitable for thermal analysis or vapor pressure studies. Experimental Methods and Results The phase diagram for the U-Zn system at 1 atrn pressure, shown in Fig. 1, is based primarily on vapor pressure measurements and on thermal analysis taken at temperatures below 950°C. Fig. 2 shows the U-Zn diagram at 5 atrn pressure, constructed on the basis of thermal analysis of alloys in sealed containers up to 1150°C and on the basis of metallographic, X-ray, and analytical data. The alloys sealed under vacuum were actually under their own vapor pressure and those sealed in an atmosphere of helium were under an additional pressure due to the helium. At temperatures up to 1100°C the zinc pressure is 5 atrn or less for these alloys; consequently the maximum pressure over the alloys sealed under a helium atmosphere was 10 atrn or less at temperatures up to 1100°C. Changes in pressure of this order of magnitude do not appreciably alter the position of most solid-solid or liquid-solid phase boundaries. In constructing the phase diagram for a pressure of 5 atm, the effect of pressure on all phase boundaries except those for liquid-vapor or solid-vapor regions was considered negligible.

Jan 1, 1958

By L. K. Jetter, C. J. McHargue

The use of axis charts for representing the texture of cold-rolled thorium sheet was compared with conventional pole figures. Four texture components were deduced from the axis charts and shown to be consistent with the pole figures. It was shown that consideration of the pole figures alone could lead to an apparent component which is the average of two actually present. It was also shown that the spread about the "ideal" textures and the amount of material associated with each could be readily obtained Fom the axis chart method. BECAUSE of the crystallographic nature of the deformation mechanisms in metals, plastic flow results in the alignment of certain crystallographic directions with the primary flow directions. In the study of the "preferred orientation" thus developed, one wishes to discern the distribution of the crystallographic directions with respect to some reference direction, usually a direction related to the fabrication procedure, such as the rolling direction. The pole-figure method, introduced by ever,' involves the plotting of the distribution of the poles of some family of planes on a stereographic projection which places the reference direction(s) at the north pole and/or in the center. By noting the angles between the reference directions and concentrations of poles on several such pole figures, one can deduce the crystallographic directions which are aligned with the reference directions. If the preferred orientation is relatively simple and well-developed, one to three pole figures may be sufficient for obtaining the desired information; if there are several texture components, more pole figures may be needed. A significant improvement in the method of representing the textures in wires and rods resulted from the introduction of the "inverse pole figure" or axis-distribution chart by Harris.2 Such a chart gives directly the distribution of the reference axis relative to the standard stereographic projection. There has been widespread acceptance of this method for the representation of fiber textures, and methods have been proposed for converting observed diffraction intensities into axis densities.2-6 The suggestion by Mitchell and Rowland3 and Jetter, McHargue, and williams5 that sheet textures could be represented as distributions of rolling, normal, and transverse directions have been criticized on the basis that they offer no significant advantage over a combination of two ordinary pole figures and that it has not been proved that sheet textures can be treated as quasi fiber textures.6 If the purpose of a texture determination is to study some property which is related to the orientation of a particular crystallographic plane or direction, then conventional pole figures certainly are sufficient. For example, the distribution of the slip planes in face-centered cubic metals (the (111) pole figure) serves to indicate the variation in mechanical properties of sheet. On the other hand if the study is directed at an understanding of the factors influencing texture formation, it is desirable to know not only the preferred orientation but the kind and degree of spread from it and the relative amounts of material in each component, if there are more than one. Axis charts give this information as well as the orientation distributions. Consider the procedure for deducing a texture component from a pole figure for the (111) planes

Jan 1, 1961

By H. C. Fiedler

A high degree of cuhe-on-edge grain orientation and good magnetic properties were obtained in Si-Fe strip processed from laboratory heats containing vanadium nitride inclusions. The higher the nitrogen content the greater was the restraint on normal grain growth and the greater was the ability to undergo secondary recrystal-lization. The degree of restraint in a heat containing magnanese sulfide inclusions was slightly greater than that in the lowest nitrogen vanadium nitride heat. The high rate of diffusion of nitrogen in the temperature range in which the texture is developed makes it necessary to retard the loss of the inclusions until the texture has been developed. MAY and Turnbulll have shown that to develop the cube-on-edge secondary recrystallization texture in Si-Fe it is necessary that inclusions, such as manganese sulfide, be present to prevent normal grain growth. They showed that the driving force for the growth of the secondary grains is inversely proportional to the diameter of the matrix grains, and that the function of the inclusions is to maintain a small matrix grain size. Fast3has developed cube-on- edge texture in Si-Fe by nitriding the hot rolled band. The silicon nitride thereby formed promotes secondary recrystallization during the anneal of the final gage strip. To obtain the best magnetic properties, the inclusions must be removed from the strip after the grain orientation has been developed. Nitride inclusions have an inherent advantage over sulfide inclusions in that the greater diffusivity of nitrogen as comared to sulfur permits the use of shorter times and or lower temperatures for the purification heat treatment. The purpose of the work to be described was to examine the usefulness of vanadium nitride inclusions in developing the cube-on-edge texture and to determine the magnetic properties of oriented strip made from heats containing these inclusions. A comparison is made with a heat having had iden-

Jan 1, 1962

By T. P. Floridis

It was deemed desirable to obtain an understanding of the vacuum desulfurization process. McKechnie1 has reported that the sulfur content of nickel- and cobalt-base alloys is reduced in vacuo. Keverian and Taylor2 have observed that vacuum melting causes desulfurization of iron-carbon-silicon alloys. Garnyk and samarin3 have desulfurized transformer steel by vacuum treatment. Experiments were conducted in order to investigate the desulfurization of liquid iron under reduced pressures, and to establish the effect of some alloying elements, and various crucible materials. A vacuum-induction furnace of 50-lb capacity was used. The charge was Armco iron. Silicon was added as 90 pct ferrosilicon, carbon as electrode graphite, and aluminum as high-purity aluminum metal. Magnesia and silica crucibles were used, their inside diameters were 5 1/2 and 5 3/8 in., respectively. The temperature was measured with an optical pyrometer and it was held at 1600 1 15°C. Samples were withdrawn with quartz tubes of 6 mm ID, without disturbing the furnace atmosphere. Pressures were measured with an alphatron gage. The experimental results are shown in Fig. 1. They indicate that: 1) Sulfur can be removed from liquid iron alloys by vacuum treatment; 2) The desulfurization is enhanced by the presence of silicon, carbon, and aluminum, all these elements increase the activity of sulfur dissolved in liquid iron;4 3) The pronounced effect of silicon indicates that probably a volatile silicon sulfide is formed, Alcock and Richardson5 have observed a similar phenomenon; 4) The same rate of desulfurization was obtained when the metal was contained in either magnesia or silica crucibles. Dr. J. Keverian has contributed many discussions and recommendations. M. S. Sutliff has helped in conducting the heats. The permission of General Electric Co. for this publication is gratefully acknowledged.

Jan 1, 1960

By I. R. Kramer, H. Shen, S. E. Podlaseck

The tensile and creep properties of aluminum in vacuum have been investigated. It was found that the general effect of a vacuum environment was to reduce the rate of work hardening. Results obtained from creep tests showed that lowering the test pressure from 2 x 10-4 to 10-7 torr produced a decrease in the activation energy as much as 500 cal per mole. An attempt is made to explain these effects in terms of how the surface-stress field of a specimen may be affected by a vacuum environment during plastic deformation. In a previous paper,' the results were reported of some of our early work on the effect of vacuum environment on the stress-strain curve of aluminum single crystals. This effort has been extended to the studies of the tensile and creep properties in vacuum as a function of various metallurgical parameters. Some of the recent surface-effect studies2'3 have revealed that a region of high dislocation density, the so-called "debris" layer, is formed near the surface of the specimen as a result of plastic deformation. This layer acts as a surface barrier to oppose dislocation movement. The strength of this barrier depends on the surface condition such as oxidation or surface removal. Experimental results obtained from surface-removal work2, 3 indicate that such a surface barrier, which may be characterized by a stress field, affects the applied stress and the rate of work hardening. In the case of aluminum, this stress field was found to be directly proportional to the strain and the strain rate; however, the slope of this curve is independent of orientation of the crystal and may be altered by using different surface treatment. Since the escape of dislocations may be retarded by the presence of an oxide on the surface of a specimen, it is expected that a change in the rate of oxidation in vacuum will affect the process of dislocation accumulation in the debris layer. Consequently, it also affects the plastic behavior of a metal. This paper reports the changes in some of the creep and tensile properties of aluminum in vacuum and discussions of the observed effects in terms of the concept of the "debris" layer. EXPERIMENTAL PROCEDURE Tensile Tests. A sketch of the tensile apparatus is given in Fig. 1. This apparatus was contained in a bell jar (14 in. diam by 24 in. high) connected to a 10-in. vacuum system. The load was applied to the specimen by means of a gear train which was driven by a magnetic clutch located outside the vacuum chamber. The specimen, held by a pair of square shoulder grips, was attached to the load cell at the top and the drive gear at the bottom. A displacement transducer was placed between the two grips as shown in the figure to measure the elongation. Signals of elongation and load were amplified and recorded on an X-Y recorder. The system was capable of detecting loads as small as 0.025 lb and displacements of 0.0002 in. Shoulder-grip specimens of aluminum single crystals, having a nominal square cross section of 1/8 by 1/8 in. and a 3-1/2-in, gage length, were grown by the Bridgeman techniques. The original purity was 99.997 pet. The orientations of the crystals investigated are shown in Fig. 4(a). The crystals were annealed at 600°C for 2-1/2 hr, electropolished, and placed in the vacuum test chamber without any delay. All the crystals were deformed at a strain rate of 4 x 10-5 sec-1 at 24o ± 3°C. A study was made to find whether a chemically passive surface film would give the same effect as a vacuum environment. Specimens of aluminum 1100 (0.187 in. diam with a grain size of 0.1 mm)

Jan 1, 1965

By H. T. Sumsion, A. U. Seybolt

The results of an investigation of vanadium-rich V-O solid solutions are presented, indicating the structure and lattice parameters of two solutions, a and ß, and their approximate temperature-composition existence. The a solution is the terminal body-centered cubic one, and contains up to 3.2 atomic pct 0. The ß solution has an ordered body-centered tetragonal structure, is formed at 1270°C, and exists from about 15 to 22 atomic pct 0. From the evidence available, the various phase boundaries have no appreciable temperature dependence. Evidence has been found for a polymorphic transformation in pure vanadium at 1550°C. IN an earlier investigation' dealing with the preparation of pure vanadium by calcium reduction of the oxide, it was found that small amounts of oxygen drastically reduced the ductility of the metal. Because this effect was so marked, it was decided to make a study of the solubility of oxygen in solid vanadium. This report deals principally with this solubility and the nature of the phase relationships in the vanadium-rich region, particularly at temperatures below 1300 °C. However, during the investigation enough data on the V-0 system were obtained to make it appear worthwhile to present a tentative phase diagram up to the composition VO. The only significant prior work found on this system are the contributions of Klemm and Grimm,' and Mathewson et al.3 Klemm and Grimm prepared a wide range of V-O compositions by powder techniques including the compositions VO.l, VO.2, VO.3 and VO4 (9.1, 16.8, 23, and 28.6 atomic pct 0, respectively). The first three compositions were found to consist of a body-centered tetragonal solid solution, while the last also showed lines of VO (NaCl structure). They found that the parameter c, increased and the parameter a, decreased with increasing concentration of oxygen. For their composition VO.27, or about 16.8 atomic pct O, they cite the values a, = 2.948A, c, = 3.53A, and c/a = 1.2. Klemm and Grimm made no attempt to determine the solid solubility limit nor to construct a phase diagram. They did, however, give some data on the homogeneity range of VO, and they proposed a structure for the body-centered tetragonal solid solution; these points will be taken up later. Materials and Preparation of Samples The vanadium used in this investigation was prepared in the laboratory by the method previously mentioned.' A typical analysis is as follows: Fe, 0.007 pct; Si, 0.02; Ca, 0.06; C, 0.224; O2, 0.044; N,, 0.0017; H2, 0.003; and V, 99.34 ±0.3 (assay). The vanadium assay is probably low by about the error given. The impurities total about 0.36 which, if subtracted from 100, gives a purity of about 99.6. At the time material was being prepared for this work no suitable technique was available for melting vanadium without appreciable contamination. The procedure adopted therefore was to cut the calcium-reduced regulii into slices which were then rolled to strip about 0.025 in. thick for oxygen diffusion. Pieces of rolled vanadium of approximately 0.025xY4xl in. and weighing about 0.3 g were suspended in a vertical fused-silica tube which was part of an ordinary gas absorption apparatus. The silica tube was heated by an electric resistance-tube furnace which could be raised around the silica tube or lowered away from it as desired. This apparatus had no novel features which require detailed description. Other than the silica tube and furnace, it consisted of a glass system evacuated by a liquid nitrogen trapped mercury diffusion pump, a mercury-operated gas burette, a McLeod and a Pirani gage, and a mercury manometer. It was also equipped with suitably located stopcocks for isolating various parts of the system; the vacuum ordinarily attained was between 10" and 10-6 mm of mercury. Oxygen generated by decomposing MnO2 was passed through anhydrous magnesium perchlo-rate before introducing it into the gas burette and thence to the absorption chamber.

Jan 1, 1954

By H. A. Saller, F. A. Rough

Studies of the V-U system have been made to determine the constitutional diagram. The diagram is fairly simple, since no intermediate phases are formed. Additions of vanadium lower the uranium melting point and allotropic transformation temperatures. The maximum solid solubility of vanadium in uranium is 12 atomic pct and the maximum solid solubility of uranium in vanadium is about 4 atomic pct. WITH the advent of high purity iodide vanadium,' it became of interest to study the alloying characteristics of vanadium and uranium. Alloys were prepared covering the entire system and were studied using a combination of metallographic examination and thermal and X-ray analyses. As a consequence of this work, the V-U constitutional diagram was evolved. In order to prepare suitable alloys for equilibrium-phase studies, it is desirable to have the purest possible starting materials. For this investigation, calcium-reduced vanadium, de Boer process crystal-bar vanadium, and good Mallinkrodt Chemical Co. uranium were used. The first work was done using the best vanadium then available, that being calcium-reduced vanadium produced at Knolls Atomic Power Laboratory. Subsequently, the important features of the diagram were checked using alloys made with crystal-bar vanadium produced at Battelle. These checks are in agreement with the earlier work. Analyses of the starting materials are contained in Table I. Only minor amounts of any impurities were present in either the vanadium crystal bar or uranium. Nevertheless, care was exercised in selecting the uranium for the alloys. The procedure followed was to section the stock into suitable pieces for melting, and macroetch each piece. Only those pieces which appeared to be clean by macroetch examination were used for alloying. Experimental Procedure Melting: Alloys with 0 to 50 atomic pct V were induction melted in beryllia (BeO) crucibles under vacuum of 10 microns pressure or better. A number of remelts was frequently required to insure uniformity of the individual ingots. Although no visible reaction between the alloy and crucible was found in any of these operations, the alloys were checked by spectroscopic analysis. Only trace amounts of beryllium were detected in the alloys. Alloys containing 50 to 100 atomic pct V were arc melted in a water-cooled copper crucible under an atmosphere of helium. A tungsten electrode was used. Ingots of 10 to 15 g were produced readily by turning and remelting a few times. Such small ingots were necessary because high purity vanadium was not available in appreciable quantities at this time. After melting, the ingots were examined metal-lographically for segregation and were consolidated by remelting before fabrication was attempted. Spectroscopic analyses showed that only small amounts of tungsten were picked up during the arc-melting process. Fabrication of Alloys: Alloys containing up to 20 atomic pct V could be fabricated by rolling in air at 600°C with 5 min between passes. After each pass through the rolling mill the alloys were returned to

Jan 1, 1954

By J. T. Williams

The equilibria in the V-Zr alloy system were investigated by solidus temperature determinations, thermal analysis, dilatometry, electrical resistance measurements, microscopic examination, and X-ray diffraction analysis. There is a eutectic reaction at 1230°C between a compound, V2Zr, and a solid solution containing 10 pct V in ß zirconium. V2Zr decomposes at 1300°C into liquid and a solid solution containing about 10 pct Zr in vanadium. The eutectic composition is probably about 30 pct V. A eutectoid reaction between V2 Zr and a zirconium takes place at 777°C at a very high rate. The eutectoid composition is 5 wt pct V. The limit of solubility of zirconium in vanadium was estimated to be 5 pct at 600°C. No attempt was made to determine the liquidus for the system. THE recent availability of large quantities of high purity zirconium has stimulated the study of zirconium binary systems. The equilibrium diagram for the V-Zr system has received little attention, however. Wallbauml appears to have made the first report concerning the equilibria in these alloys. He reported the existence of a compound, V2Zr, having the MgZn2 ((214) type of structure with a, = 5.277 kX and c° = 8.647 kX. Anderson, Hayes, Roberson, and Kroll2 made a survey of some potentially useful zirconium binary alloys and found that zirconium probably dissolves a small amount of vanadium. They reported the probable existence of a compound between the two elements and suggested that the zirconium-rich solid solution undergoes a eutectic reaction with this compound. Pfeil," in a critical review of the existing information, estimated that the solubility of vanadium in zirconium is less than 4.7 pct and probably less than 1.8 pct. Rostoker and Yamamoto' proposed a partial diagram for the V-Zr system in a survey paper on vanadium binary alloys. Their diagram indicates the compound, V,Zr, a eutectic reaction at 1360°C, a peritectic reaction at 1740°C, and a limit of solubility of zirconium in vanadium of about 3 pct. They obtained no information on the equilibria in the zirconium-rich alloys. In view of the potential utility of the V-Zr alloys and the incomplete knowledge concerning the equilibria in the system, an attempt was made to establish the constitutional diagram. Preparation of the Alloys Raw Materials: The vanadium for making up these alloys came from the Electro Metallurgical Corp. Zirconium came from two sources. In the beginning of the investigation, sponge zirconium from the Bureau of Mines was used in making some of the alloys. Later, iodide metal made at the Westinghouse Atomic Power Development Laboratories became available. This material was used in the preparation of all the dilatometric and resistance specimens and about two-thirds of the solidus temperature specimens. A typical manufacturer's analysis of the vanadium is shown in Table I. No other analysis of the vanadium was made. The metal contained a dispersed second phase and did not have a sharp melting point. Typical results of spectrographic analysis of the Westinghouse zirconium are shown in Table 11. These data indicate a very high purity. The Bureau of Mines sponge metal was probably less pure but had good ductility. Melting: All of the alloys used in the investigation were made by melting pieces of vanadium and zirconium together in a dc electric arc furnace similar to those of Geach and Summers-Smith, craighead, Simmons, and Eastwood," and others. Melting was done in an atmosphere of helium scavenged of residual air by the preliminary melting of a separate charge of zirconium. Each ingot was turned over and melted at least three more times before removal from the furnace to aid in the attainment of homogeneity. Alloys prepared for use in the investigation are listed with the results of solidus determinations in Table III with the exception of the following compositions upon which no solidus determinations were made: 0.29, 0.54, 4.57, and 5.55 pct V. Analysis: The weight of each ingot made from iodide zirconium was within 0.1 g of the total weight of the initial charge, about 90 g. Since each component of each charge was weighed to the nearest 20 mg for amounts less than 10 g and to the nearest 0.1 g otherwise, the gross composition of an ingot could be calculated accurately. Chemical analysis for the vanadium content of several alloys agreed

Jan 1, 1956

By P. Chiotti, G. R. Kilp

IN the study of zinc-zirconium1 and zinc-uranium2 systems, vapor-pressure measurements were made in order to establish the equilibrium phase diagrams for a pressure of 1 atm. In each of these systems, only zinc has an appreciable vapor pressure over the temperature range investigated. The zinc pressure was measured over a sufficiently wide temperature and composition range to permit calculation of the standard free energy, enthalpy, and entropy of formation for the compounds ZrZn, ZrZn2, ZrZn3, ZrZn6, ZrZn14 , and U2Zn17. In these calculations the compounds were treated as line compounds. Although the exact composition range was not determined, metallographic and thermal data indicate that the solubility range for these compounds is small. It is to be noted that the compound initially reported to correspond to the formula UZn9, 2 has been shown by the crystal structure work of Makarov and vinogradov3 to correspond to the formula U2Zn17,. The crystal structure is similar to that for the compounds Th2M17, investigated by Florio,4 et al; here M represents iron, cobalt, or nickel. Therefore, thermodynamic calculations are based on the formpla U2Zn17 EXPERIMENTAL The dew-point method employed to measure the vapor pressure of zinc over uranium-zinc alloys has been described previously.2 The same procedure and apparatus were used in measuring the zinc pressure over zinc-zirconium alloys. In the case of alloys in the two-phase region consisting of a zirconium and the compound ZrZn, the Knudsen effusion method was employed to obtain a few values in the temperature range 600" to 560°C. In most cases, tantalum crucibles were used to contain the zirconium alloys. THERMODYNAMIC PROPERTIES OFU2Zn17 A least-squares treatment of the vapor-pressure data obtained for 35 and 65 wt pct U alloys in the temperature range 650° to 940°C gave lOg10PAtm = 7,713/T + 6.33 Here, as in following equations, T represents degrees kelvin and P represents the zinc vapor pressure in atmospheres. The probable errors for the two constants were calculated to be 60.68 and 0.0563, respectively. This relation represents the equilibrium dissociation pressure for the reaction U2Zn17 (s) - 2U(s) + 17Zn(v) The solubility of zinc in solid uranium is negligible and the only additional information required to calculate the free energy of formation of U2Zn17 is the free energy of vaporization of pure zinc. K. K. Kelley5 gives the relations, AF°= 30,902 + 6.03Tlog T + 0.275 X 10-3T2 - 45.03T

Jan 1, 1961

By Glenn R. Zellars, J. P. Morris

The carrier gas method was used to measure the vapor pressure of copper over liquid copper and Fe-Cu alloys. From these results, the activity of copper in the alloys was determined at 146°, 1550°, and 1600°C. Calculations of the activity of iron were made from the results for copper. The activities of copper and iron show strong positive deviations from Raoult's law. KNOWLEDGE of vapor pressures of liquid metals has both practical and theoretical value. Fuming of metal baths is directly related to vapor pressure, and the data should bc particularly important to the technology of vacuum metallurgy. Theoretical applications include determinations of heats of vaporization, heats of solution, and activity coefficients in alloys. Methods for measuring vapor pressures of metals have been reviewed by Speiser and Spretnak.' High-vacuum procedures, such as the Langmuir and Knudsen methods, have been used most widely in recent years. However, these methods are not applicable to vapor pressures greater than lo-' mm Hg and often cannot be used in the temperature range of greatest interest. The purpose of this investigation was to develop a procedure for measuring vapor pressures of pure metals and alloys that would be applicable to pressures greater than lo-' mm Hg. The transportation or carrier gas method has been used. Copper was chosen for the initial work because its vapor pressure has been established fairly well and the reliability of the technique could be ascertained. Applicability to alloys was studied by determining the vapor pressure of copper over liquid Fe-Cu alloys. Experimental Procedure The tests were carried out in a graphite-spiral resistance furnace. The charge consisted of 50 g of metal contajned in a covered sintered-alumina crucible 1/2 in. deep and 13/8 in. ID. Two sintered-alumina tubes led from outside the furnace to within the crucible, passing through close-fitting holes in the crucible cover and radiation shield. The arrangement of the crucible and the two alumina tubes is shown in Fig. 1. The small-bore tube (1/16 in. ID by 5/16 in. OD) served as the inlet for the carrier gas. The tip of this tube touched the metal surface and the flow of gas imparted a rippling motion to the metal. In this manner the gas was brought into close contact with the metal and became saturated with metal vapor. In most of the tests the flow rate of the carrier gas was 560 ml per min, measured at room temperature. Approximately half of this saturated gas was allowed to escape from the crucible through the second alumina tube (3/16 in. ID by 5/16 in. OD) and was discharged to the atmosphere through a

Jan 1, 1957

By C. C. Herrick

The vapor pressure of indium has been measured by the torque-effusion technique, as a function of temperature between 1102o and 1422oK. For liquid indium, the vapor pressure (in atmospheres) can be represented by the equation -R In P = 5.782 x104/T - 25.29from 1000o to 1422oK. Extrapolation yields a calculated normal boiling point of 2329°K and a heat of vaporization at 298 oK of 58.09 kcal per mole by the third-law method and of 58.39 kcal per mole by the second-law method. Calculations indicate that the free-energy functions of Hultgren are to be preferred to those of Guruich. BECAUSE of the potential uses of liquid metals, it is important that convenient experimental methods be developed which will lead to a sufficient body of thermodynamic data so that reasonable estimates of properties of liquid metals can be made. Indium is particularly well-suited to investigation, since it remains liquid over a long temperature interval, it alloys readily, suitable containers for it are easily procured, and relatively few investigations of its thermodynamic properties in the region above 1000°K have been made. PREVIOUS RESEARCH The first determination of the vapor pressure of indium was made by Anderson,' using the Knudsen method and silica-lines stainless-steel cells. Kohlmeyer and spandau2 determined the normal boiling point, and a mass-spectroscopic investigation has been made by Lyubimov and Lyubitov.3 (Unfortunately, only an equation representing the data of the latter investigators has been available to us.) Hultgren et al.4 and Gurvich5 have published discordant tables of the free-energy functions for indium. EXPERIMENTAL METHOD The torsion-effusion principle has been described amply in the literature.8,7 With this method, one observes the angular deflection induced by effusion of vapor from a multihole effusion cell which is suspended from a filament producing a small rotational restoring force. The pressure within the effusion cell can be computed from the relation ziaifiai where k is the torsion constant of the filament, 6 is the observed angular deflection, and ai, fi, and qi are the area, force factor, and moment arm of the ith effusion orifice. The force factor f is analogous to the Clausing factor,' and corrects for the reduction in effusive force attributable to the finite geometry of the effusion orifice. Schultz and Searcy' have tabulated the force factors for various tube geometries. EXPERIMENTAL APPARATUS The physical arrangement employed in this study is shown in Fig. 1. The furnace elements are six Globar heat rods with 15-in. hot zones. Temperature is maintained to * 1°K by a Beck mechanism driving two 26-amp Powerstats in parallel. A Pt (Pt, Rh 10 pct) thermocouple serves as the sensing element. Four inner tubes act as vibration isolators; the air pressure required for maximum isolation was determined empirically, and is maintained by a solenoid valve. The tungsten fiber, 9 by 0.002 in., is joined to an 0.008-in. tungsten wire or graphite rod via a machined lightweight chuck.

Jan 1, 1964

By C. E. Birchenall, C L. McCabe

IN attempting to extend vapor pressure measurements of the type previously reported by Schadel and Birchenall1 for silver and by Schadel, Derge, and Birchenall' for silver-silicon to other systems, it was observed that the materials melted at indicated temperatures 10" to 15" below their accepted melting points. Further investigation revealed that the thermocouple readings were in error due to appreciable conduction losses along the reference thermocouple wires. If the wire diameter of the reference couple inserted into the Knudsen cell was reduced, the correction for the indicating couple changed in a manner tending to explain the melting behavior. When extrapolated to zero wire diameter from measurements with several reference thermocouples of different wire thickness, the melting point of silver then agreed with the indicated temperature at which silver chips were observed to coalesce into a sphere. Approximately the same calibration was given by observing the melting of small wires of silver or gold in the Knudsen cell connected in series with an ammeter, where the leads into the cell were very fine in order to minimize heat conduction. Unfortunately neither of these methods seemed to yield a sufficiently precise temperature calibration to match the apparent precision of the other aspects of the vapor pressure measurement. It was decided. therefore, to redetermine the vapor pressure of silver in another setup under conditions permitting precise temperature measurement. The vapor pressure of pure silver could then be used as an internal calibration of temperature in the older unit in making runs on alloys. This has been done; the present report is a correction to ref. 1. Experimental Procedure The apparatus, shown in Fig. 1, was very similar to that employed by Harteck,3 except that the orifice sizes were smaller and the residual pressure in the vacuum system was probably much lower. A small, sharp-edged hole, nearly circular in shape, was ground into the rounded end of a quartz tube. The orifice area was then measured by tracing the image at known magnification on graph paper and counting the squares enclosed. The silver specimen was sealed into the tube to make a Knudsen cell. A tantalum jacket surrounding the cell served to increase the uniformity of temperature. This assembly was placed in the bottom of a long quartz tube with an inside diameter of about 1 in., which was connected to the vacuum system through a ground joint sealed with picein wax well removed from the furnace. A thermocouple tube inserted through the top of the vacuum line reached into the tantalum jacket so that the thermocouple junction was immediately adjacent to the Knudsen cell except for the protection tube wall. A resistance furnace could be raised to cover the end of the quartz tube containing the cell in such a way that the cell was in the uniform temperature zone 13 in. from the end of the furnace. An ionization gage was included in the vacuum system in the cold lines of wide diameter, immediately beyond the ground joint. The vacuum system consisted of a mercury one-stage diffusion pump, backed by a Welch duo-seal mechanical pump. The pumps were separated from the reactor chamber by a dry ice trap. The ionization gage always read less than 10-5 mm Hg after initial outgassing and before each run was started. Each newly filled Knudsen cell was evacuated at high temperature overnight before the first weighing was made. The cell was returned to the system, heated for a measured time at constant temperature, cooled, and reweighed. The heating and cooling times were quite short since the hot furnace was raised to receive the reactor at the beginning of the run and removed again at the end. The tube heated or cooled quickly. The total mass loss was attributed entirely to effusion of silver vapor from the quartz cell, since empty quartz cells maintained constant mass through similar heating cycles. The vaporized silver condensed on the cold walls of the quartz tube extending above the furnace. Earlier studies in the induction heated unit had shown that the same vapor pressure was found for silver, whether the silver was in contact with the tantalum metal cell or with porcelain or quartz liners. The Pt-Pt-10 pct Rh thermocouple was calibrated against a secondary standard of the same material and found to agree with the published tables. Always operating in air at temperatures below 100O°C,

Jan 1, 1954

By C. E. Birchenall, H. M. Schadel

Vapor pressure of silver over silver-gold alloys has been measured over a range of temperatures for four compositions. Orifice effusion has been compared with electromotive force measurement as a means for determining thermodynamic activities in these solid metallic alloys. Correlated activity coefficients are given for silver-gold alloys for the temperature range 200' to 1000°C. SOUND development of a technique for the measurement of some property of matter requires a comparison with other recognized procedures which have been used to measure the same property. Such a check is especially useful if the assumptions on which the applications stand are basically different for the two methods. Agreement in the results tends to confirm both sets of assumptions. The electromotive force method for determining thermodynamic activities depends upon the reversibility of the electrode processes and a knowledge of the valence states involved. The orifice effusion technique of determining vapor pressures requires that certain mean free path conditions be met, that the state of association of the vapor be known, and that efficient collection of vapors be maintained. Thus the requirements are quite different in nature. Although procedures often exist by which the fulfillment of these conditions may be tested internally, agreement of thermodynamic activities for a single system obtained by both methods constitutes a more severe test of the assumptions. The binary alloy system silver-gold is excellent for this purpose for several series of emf measure-

Jan 1, 1954

By E. A. Gulbransen, K. F. Andrew

Weight loss measurements were made using a sensitive microbalance operating in a high vacuum system. The Langmuir equation was used to Calculate the vapor pressures of the several metals and alloys. Aluminum lowered the vapor pressure of i9,on in a 5.4 Al-94.6 Fe alloy, and of iron and chromium in a 4.8 Al-21.5 Cr-73.7 Fe alloy. The influence of oxide films formed on the alloys on the effective vapor pressures of the alloys was also studied. These studies were used to aid in the inte9,pretation of the high temperature oxidation of alloys based on iron, chromium, and aluminum. RECENT studies on the oxidation of chromium1 and the heat resistant alloy 5 Al-22 Cr-73 Fe showed transitions in the rate of oxidation at 900o and 1050oC respectively. The transition for chromium occurred at a temperature where the rate of evaporation of chromium from oxide-free surfaces equalled the rate of chromium atoms reacting with oxygen. This paper presents new vapor pressure studies on iron, chromium, and the specially prepared alloys 5.4 Al-94.6 Fe, 21.9 Cr-78.1 Fe, and 4.8 Al-21.5 Cr-73.7 Fe. The ternary alloy is similar in basic composition to the patented heat-resistant commercial alloy known as Kanthal. The binary alloys were studied to show the individual effects of aluminum on the vapor pressure of iron and of iron on the vapor pressure of chromium. The purpose of the proposed studies was to test the influence of metal volatility on the oxidation behavior of heat-resistant alloys based on iron, chromium, and aluminum. Since oxide films formed on the metal may limit metal transfer to the surface, the influence of oxide films on metal volatility was also studied. LITERATURE A) Vapor Pressure. The vapor pressure of iron has been studied by Jones, Langmuir, and Mackay,' Dornte and Nrton. Edwards. Johnston. and Ditmar. and McCabe, Hudson, and Paton. Stull and SinkeT have calculated a heat of sublimation at 298° K of 99.83 kcalper g atom. The vapor pressure of chromium has been reported by Bauer and Brunner,' Speiser, Johnston, and Blakburn, Gulbransen and Andrew,'' Vintaikin,o McCabe, Hudson, and Paton, and Kubaschewski and Heymer.12 Good agreement was obtained except for the early work of Bauer and Brunner.' Stull and Sinke7 calculate a heat of sublimation at 298°K of 95.0 kcal per g atom. Vapor pressure measurements on aluminum have been made by Brewer and Searcy,13 Bauer and Brunner,aand Farkas.I4 Stull and Sinke,7 giving Brewer and Seary's' data the most weight, derive a heat of sublimation of 77.5 kcal per g atom at 298 oK. B) Method. Two methods15 are used for measuring the vapor pressure of metals: 1) The Langhuir free evaporation method, and 2) the Knudsen effusion method, In the Langmuir method, used in the present work, the vapor pressure of the metal P is given by the equation: Here M is the molecular weight of the vapor species. T is the absolute temperature. (dw/dt) is the rate of sublimation in g per sq cm per sec and a is the condensation coefficient. a is a measure of the efficiency of condensation of molecules striking the metal surface from the vapor. If condensation results from each collision a is unity. This coefficient has been found to be unity,a, le for metals. These results establish the general validity of the Langmuir free evaporation method as applied to metals. EXPERIMENTAL A) Vacuum Microbalance Method. A modification of the Langmuir free evaporation method was used. Strip specimens were suspended from a sensitive microbalance operating inside of a high vacuum system. For alloys, microbalance methods17 must be used on large area samples at relatively low temperatures to minimize composition changes. The

Jan 1, 1962