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By D. H. Baldwin, R. E. Reed-Hill

DURING the course of experiments involving oxygen equilibrations with a high-purity Pd-5 at. pct Rh alloy, the appearance of a subscale was noted. Most of the heat treatments in a pure oxygen atmosphere were performed at 900°C. The presence of the subscale was noted when the oxygen partial pressure was in excess of approximately 95 mm, up to 1 atm, the highest pressure used. According to available thermodynamic information' on the free energy of formation of Rh2O3 (the only reliably reported solid oxide of rhodium), this oxide should not be stable under these conditions, nor, of course, any oxide of palladium. Extraction of the subscale from the matrix was affected by treatment with aqua regia, and the resulting X-ray diffraction powder pattern is shown in Table I. This pattern is not that of Rh2O3, nor that of any known rhodium or palladium oxide. It can be indexed on a hexagonal lattice with a. = 5.22Å and co = 6.0Å. In addition to this unknown material, there was a weak Rh2O3 pattern. An X-ray fluorescence analysis of the powder gave 3214 palladium counts per sec and 5381 rhodium counts per sec, thus indicating the likelihood of a double oxide of palladium and rhodium. In view of the fact that the above result was obtained on a single sample, it is perhaps best not to insist upon the above ratio of about 0.6, but possibly to tentatively assume 0.5 or 1/1. The small amount of Rh2O3 in the X-ray pattern might have occurred on cooling, and it seems possible that the double oxide, below some temperature level, is not as stable as Rh2O3. To examine this point, some of the extracted double oxide was heated on a sheet of pure palladium in air at 900°C, and allowed to furnace-cool over a period of several hours. The X-ray diffraction pattern of the powder showed approximately the same intensity of Rh2O3 and the "hexagonal" oxide as before, but, in addition, a rather faint pattern of another crystalline species made its appearance. This finding again suggests partial transformation of the major (double) oxide to other oxides on cooling from 900°C. Application of the phase rule shows that only one equilibrium oxide (subscale) could have been present at temperature. Vacuum-fusion, oxygen analysis of the extracted subscale showed 7.7 ± 0.7 pct O. Combining this information with that of the X-ray fluorescence analysis suggests the formula PdRhO. This compound should have an oxygen content of 7.07 pct, which is within the range of error of the oxygen analysis. However, until a more careful stoichiometric study is made, it is suggested that the formula be written

Jan 1, 1965

By Robin O. Williams

The increase in internal energy as the result of deformation has been measured for lead, aluminum, silver, nickel, iron, and zirconium by using rapid, adiabatic compression. The stored energy increase is roughly Proportional to the strain; the propor-tionality constant increases rapidly with increasing melting point. The fraction of the mechanical energy which is stored increases more slowly, since the strength of the metals also increases with melting point. The values of the stored energy are considered accurate to about 10 pet. The present values appear about 50 pet larger than the more reliable published results where comparisons are possible. It is possible that this difference is due to the high strain rate used in this investigation. Immediately after deformation all these metals release energy at a rate roughly proportional to (time)-'. This release is considered to be associated with dislocation motion but in aluminum (and copper) some additional process seems to be present. This release can represent 20 pet or more of the stored energy. WHEN pure metals are plastically deformed, most of the mechanical energy is converted into heat. The energy remaining within the metal is significant in that it is the energy of the disorder produced and thus detailed knowledge of this energy is a powerful tool in understanding the nature of deformational disorder. While much effort has been expended on this problem, the amount of information available is limited. The situation as of 1958 has been carefully reviewed by Titchener and ever.' The present results have been obtained by a new experimental approach to this problem. The method necessitates high-strain rates which make comparisons with published results less certain, but a high-strain rate is an advantage in that the energy release immediately after deformation can be followed. EXPERIMENTAL PROCEDURE In the experimental method used, the internal (stored) energy is given as the difference between the mechanical work used in deforming the sample and the heat which is released (the first law of thermodynamics). The work is supplied by two identical hammers swinging freely from a fixed height, the available energy being the product of the mass, the gravitational constant, and the distance through which the center of gravity moves. The initial temperature rise of the sample represents the heat produced by the deformation. The sample temperature is determined by a small thermocouple embedded in and supporting the sample. This process can be repeated over and over to produce increased strains. Most samples are run through about five cycles to a total strain of around 0.7. Further details are covered elsewhere. The determination of the heat is dependent upon the sample weight, its specific heat, the rise in temperature, and any gains or losses to the surroundings. One is dependent on published results for specific heals (the values were taken from a collection).3 The values must be accurate to about 0.1 pet in order that they not affect the results. The best values thus do not contribute to the uncertainty, but this is probably not always the case. The accuracy of the temperature rise is limited primarily by knowledge of the thermocouple characteristics over short temperature intervals, but careful calibration eliminates this as an important factor. One can calculate readily the heat flow between the sample and hammers if the time of contact and the sample temperature are known and if there were no oil at the interface. Assuming the maximum temperature difference, upper values for this interchange have been calculated (the time of contact is determined from the decrease in sample length) and it may or may not be significant. However, no correction is colnsidered necessary because of the presence of a thin oil film which has a thermal conductivity much less than the metals. Except for the possible uncertainty in specific heat, the heat is considered to be adequately known. The mechitnical losses which are recognized as important are: 1) friction between the sample and the hammer:;, 2) the kinetic energy in the hammer suspension system which may not be entirely usable, 3) the rebound energy of the hammers, and 4) the vibration of the hammer heads. All these have been covered in detail elsewhere2 and only the more significant points are made here. The friction turns out to be small except for soft materials (lead)

Jan 1, 1962

By James H. Bechtold

The yield strength of annealed tungsten was found to have a strain rate exponent 12 times as great as that of low carbon steel. The effects of temperature and strain rate could be correlated through the Zener-Hollomon parameter with a heat of activation associated with yielding of 32,000 cal per g-atom. This heat of activation is independent of strain although both the temperature and strain rate dependence of the yield strength vary with strain. MOST metals have mechanical properties that are relatively insensitive to strain rate, especially at temperatures below that at which creep is important. In iron and molybdenum," however, relatively large strain rate effects are observed in the temperature range where the rapid increase in yield strength occurs which leads to the ductile to brittle transition. In this temperature range Zener and Hollomon' have suggested that the effects of strain rate and temperature T on the yield strength could be correlated through a single parameter, P, in the following manner where Q is a heat of activation characteristic of the metal and associated with some unspecified relaxation process. They also observed that if the temperature is restricted to a relatively narrow range— below room temperature for steel—that the logarithm of the yield strength could be expressed as a linear function of the logarithm of P. From this relation the logarithm of stress should be a linear function of the logarithm of the strain rate at constant temperature: where r will be referred to herein as the strain rate exponent. At a constant strain rate the logarithm of stress should be a linear function of l/T, ay(e = C) ~ eT [4] where q = Qr [5] and is referred to herein as the temperature exponent. These equations have been found to relate the effects of temperature and strain rate on the yield strength of molybdenum in the temperature range wherein the ductile to brittle transition occurs, —75° to +150°C.2 Bechtold and Shewmon3 have shown recently that the ductile to brittle transition in annealed tungsten occurs in the temperature range +150° to +350°C. In this temperature range the yield strength in-

Jan 1, 1957

By John W. Weeton, Max Quatinetz, Robert J. Schafer

Nickel powders of 1-, 0.4-, and 0.2- average particle size were combined with 0.05- p MgO powder. Oxide was added in quantities of 4, 12, and 20 vol pct. The mixtures were vacuum hot pressed and extruded at 2100°F. In spite of a general tendency to growth of oxide particles during processing, dispersions of fineness comparable to those of Sintered Aluminum Powder, SAP, were achieved. A small amount of oxide considerably improved 1800°F stress-rupture strength, but further oxide addition resulted in a decrease of strength. RECENTLY there has been much speculation concerning the requisites for strong and stable dispersion-strengthened alloys. The nature of the dispersed phase, the stability of the alloys, and variables influencing the mechanical properties of alloys have been given consideration. Such topics have been reviewed by Bunshah and Goetzel. Of the many parameters related to the mechanical properties of dispersion alloys, mean free path between dispersed particles (a measure of the fineness of the dispersion), particle size, and vol pct of the dispersed phase have been shown to be very important. Generally, strength increases with decreasing mean free path, and there is an optimum vol pct of dispersed phase. Just how fine a dispersion ought to be is somewhat open to debate, but Irmann, Lenel, Backensto, and Rose, Gregory and Grant,4 and others have shown in the original SAP-type alloys, aluminum plus aluminum oxide alloys, that better strength is generally obtained with finer dispersions and the best strengths were obtained when the mean free path was less than . Further, Cremens, and rant' showed that in nickel plus aluminum oxide alloys, stress-rupture strength at 1500°F was definitely increasing with decreasing mean free path. In their investigation, however, the finest dispersions that were produced were about 2.0 mean free path. To determine the potential of the nickel plus oxide type of product, it would be neces- sary to obtain dispersions analogous to those of SAP alloys; namely, ones where the mean free path between oxides is less than . Unfortunately, however, it is very difficult to produce extremely fine dispersions of oxides in metals other than those which naturally form refractory oxide skins. There are several approaches to this problem including such procedures as mechanical mixing, internal oxidation, reduction of mixed oxides, and metal deposition from solution. Work at the Lewis Research Center has shown that metal powders 0.1p or finer could be produced by ball milling, when selected grinding aids were utilized.' Because these unusually fine powders were available, it was an initial objective of this study to determine if by mixing of these fine Ni powders with various amounts of MgO (0.05) a fine dispersion of MgO in Ni comparable to those in Al-Al203 might be achieved. A second objective was then to attempt to relate the parameters of structure to 1800°F stress-rupture strength of the alloys. Nickel powders of 1-, 0.4-, and 0.2-p average particle size were combined with 0.05- average particle size MgO powder. This oxide was added in quantities of 4, 12, and 20 vol pct. The mixtures were hot pressed and extruded at 2100°F.

Jan 1, 1962

By G. K. Manning, E. R. Stein, E. E. Underwood

Columbium, carbon. and nickel additions were made to iron-base alloys with 20 pct CY. The effects on microstructure, precipitation-hardening characteristics, and High-temperature properties were investigated by means of metallographic, hot-hardness, and magnetic measurements. An austenitic alloy consistilzg of Fe-20Cr-14Ni-0.8C-2.5Cb had higher hot-lzaciness at 1700°F than the other alloys investigated. Its favorable standing is attributed to an enhanced interaction of Cb and Cr4C particles dispersed within a stable austenitic matrix. FOR the past 20 years, cobalt- and nickel-base alloys have been used for the more stringent high-temperature applications because of their excellent creep resistance. Although their performance is quite satisfactory, other base materials must be considered for large tonnage, low-cost applications such as for the automotive gas turbine. Iron-base alloys are also used extensively, but at somewhat lower temperatures, primarily because of the instability of particles or because precipitation occurs at relatively low temperatures. Since unusual stability has been reported for Fe-Cb and Cb-C compounds.1-3 it was hoped that effective amounts of these compounds could be taken into solution at temperatures just below the melting range, then precipitated upon reheating. Actually, strengthening was found to occur by both precipitation hardening in some alloys and by phase transformations in others. Among the alloys investigated, that with the best high-temperature properties owes its strengthening to a combination of carbide particle interaction and a relatively stable austenitic matrix. EXPERIMENTAL PROCEDURES Small preliminary melts of about 100 g were made in a vertical Glotube furnace, using alundum crucibles and an argon atmosphere. Alloys weighing 35 lb were melted in a basic-lined induction furnace, also under argon. The compositions of the alloys are given in Table I. The nickel-free alloys were usually ferritic, while those containing nickel were generally austenitic. Depending upon the particular heat treatments, or the combination of alloying elements, two-phase alloys of ferrite and austenite could be obtained. In the following text, however, the nickel-free or nickel-bearing alloys are referred to simply as either ferritic or austenitic alloys respectively. The small ingots were swaged at 2000°F, in steps, to rods of approximately 3/16-in. diam. The large ingots were forged at 2100°F to 1-in. sq bars. No difficulties were encountered in swaging or forging any of the alloys to as much as 90 pct reduction in area. For heat treating at temperatures above 2000°F, specimens were sealed in quartz tubes under a partial pressure of helium. No special precautions were found to be necessary at lower temperatures. Intermetallic compounds in the preliminary alloys were identified by means of X-ray diffraction measurements. Correlation of these results with micro-structural characteristics (such as color, shape, size, and location) of the compounds aided in their identification. Magnetic susceptibilities were also obtained at room temperature with a Magne Gage instrument after the various heat treatments. To

Jan 1, 1962

By W. Rostoker, R. P. Elliott, B. W. Levincer

Phase equilibria in the Ti-Cr-Mo system have been investigated for alloys containing 100 to 40 pct Ti in the temperature range 550° to 1300°C. Five experimental isothermal sections and a surface of incipient melting are presented in addition to a summary projection of the space. Seven vertical sections have been constructed from the isothermal sections. A STUDY of the system Ti-Cr-Mo was undertaken as part of a program sponsored by Watertown Arsenal toward the development of titanium-base phase equilibrium diagrams of potential technical importance to alloy development. The scope of the investigation included the composition range contained by the titanium corner and the line of 64 pct Ti (all compositions are by weight) and the temperature range 550" to 1100°C. A surface of incipient melting was constructed from measurements on a large number of alloys. The system Ti-Cr has been established.&apos;-? The phase diagram is characterized by a wide miscibility field of the solid solution at elevated temperatures and a eutectoid reaction of the type: (15 pct Cr) + TiCr, (685C). The a solid solution has a very limited capacity for solution of chromium (< 0.5 pct Cr). The intermediate phase TiCr2 has at best a very narrow composition range. It has been shown that two allotropic modifications of this phase exist. However, this behavior does not appear to be perceptibly reflected in the form of phase boundaries of the type + TiCr,. In the system Ti-Mo,9-10 the phase is stabilized to successively lower temperatures with increasing molybdenum content. There are no known intermediate phases. The a solid solution is limited to less than 0.5 pct Mo. Because of the very extensive range of the phase, the ternary system was ideally suited to the determination of the phase boundaries by the parametric method of Andersen and Jette." A parametric surface was constructed to provide the necessary basis for the application of this method. Preparation and Treatment of Alloys More than 60 alloy compositions were prepared for the study of this system. Alloy buttons of 10 grams were arc-melted in a nonconsumable electrode furnace using inert atmosphere protection and a water-cooled copper crucible. The construction and operation of this type of melting unit have been adequately described elsewhere.&apos;&apos; To insure complete solution of alloy additions, the buttons were remelted alternately on top or bottom as many as five times. This was accomplished by a suitable mechanism for "flipping" over the ingots. Iodide titanium (99.97 pct Ti) was used as the alloy base. An especially high quality chromium (99.9 pct Cr) was obtained from the National Research Corp. Molybdenum sheet (0.005 in.) of 99.9 pct purity was used.

Jan 1, 1954

By J. R. Hornaday, A. E. Morris, N. A. Parlee, D. C. Carmichael

Rates of absorption and effusion of hydrogen in solid iron were measurede by a Sieverts type of apparatus. With clean a iron these rates are diffusion controlled down to 420°C and are represented by the equation D = 6.70 x e-&apos;100&apos; crn2 sec-I Certain surface impurities can affect these rates at all temper-atuves. Thin oxide films lower, the apparent diffusivity of annealed iron markedly. Cold-worked iron (75 pct) annealed at 430°C has practically the same solubility and diffusivity as anealed iron. THE mechanism of absorption and evolution of hydrogen in solid iron alloys is not fully understood despite the large number of investigations on the subject. While there appear to be many instances of simple solid-state diffusion control of rate, there are many cases where investigators have either concluded that surface reactions are rate controlling or that some "nondiffusible" hydrogen resides in discontinuities. The purposes of this research were 1) to develop a sound method for studying this absorption and evolution behavior, 2) to examine it in pure annealed iron, and 3) to determine the effects of important variables, such as cold work, surface condition, and composition. This paper relates the progress toward these objectives. Much of the present knowledge of the movement of hydrogen through iron alloys comes from permeability studies which involve the measurement of the rate of passage (under steady state conditions) through membranes of the order of 1 mm in thickness. This method is open to question on various grounds. If surface reactions are involved, two different surfaces operating under two different conditions must be considered. The presence of defects in the membrane even of microscopic dimensions may lead to erroneous results. A few people have studied evolution of hydrogen into vacuum, under carefully controlled (unsteady state) conditions, and some of the best diffusivity data come from these sources. Most of these methods measure rates of evolution only and thus do not make a rigorous test for reversibility of behavior. It appeared that an "unsteady state" method that would permit the measurement of both rates of evolution and absorption under controlled temperature, pressure, and other conditions would offer several advantages. Accordingly the modified Sieverts apparatus described below was designed and built. APPARATUS This apparatus is shown in Fig. 1. The sample bulb is detailed in Fig. 2. Hydrogen of 99.5 pct purity and argon of 99.99 pct purity were used. The hydrogen was deoxidized and dried by means of a Deoxo unit followed by a "Dri-erite" tower. The argon was only dried. The gas burette was graduated in tenths of a cubic centimeter, and, with the aid of a travelling magnifier, readings could be estimated to the nearest 0.01 cc. The level of mercury in the burette was adjusted by means of a leveling bulb operated by a fine screw adjustment. An electrical system, consisting of two slightly separated platinum electrodes inserted in one leg of the manometer and a small light bulb was used to obtain rapid accurate adjustment of the pressure. The capillary tubing of the Vycor sample bulb was 3 mm bore as was all tubing back to the gas burette. The samples were discs made from vacuum-melted Ferrovac-E iron of greater than

Jan 1, 1961

By Irving Cadoff, Jack Medoff

THE linear thermal expansions of oriented single crystals of zinc were measured in the range from 20" to 416°C using a Leitz HTV optical lever differential dilatometer. The single crystals, supplied by Dr. J. J. Gilman, General Electric Research Laboratory, Schenectady, N.Y., were prepared from 99.999 pct-purity zinc, and were in the form of rods approximately 2 in. long by 0.15 in. in diameter having the following orientations: x = 1.5, 12, 20, 30, 37.5, 49, 60, 73, and 82 deg where x is the angle between the specimen axis and the (0001) plane of the crystal. From the family of expansion vs temperature curves, the instantaneous expansion coefficients were calculated at 100°C intervals, and Fig. 1 shows a plot of these derived coefficients for each orientation. Coefficients were obtained from the raw data using the chord-slope method. The curves in Fig. 1 indicate a positive temperature coefficient for low x, a negative temperature coefficient for high X, and an expansion coefficient invariant with temperature, amounting to about 43 x 106, for a crystal orientation of approximately 52 deg. From the available data, it is possible to arrive at the values for the axial thermal-expansion coef- ficients a, and a, by extrapolating the curve of the expansion coefficient vs orientation to zero and 90 deg orientations. More precise values for a, and a, are obtained if the instantaneous coefficients are plotted against the square of the cosine of the orientation angle, as shown in Fig. 2. The extrapolated terminal points of the best straight line through the data points of such a plot define the axial coefficients at that temperature. The theoretical justification for the linearity of this curve comes from inspection of a rearranged form of the general anisotropic equation for expansion: ax=a, + (a, -a,)cosZx [1I in which a, is the expansion coefficient in the "c" or [0001] direction and a, the expansion coefficient in the "a" or [ll50] direction. From Fig. 2, the axial coefficients at 100°C, for example, were found to be a, = 13.9 x 10&apos;6 and a, = 62.4 x 10"6 per "C. In the same way, the coefficients were determined for 20", 200°, 300°, and 400°C. This data is summarized in Fig. 3. It is of interest to observe the common intersection of the curves in Fig. 3, which seems to bear out the earlier observation, from Fig. 1, of a

Jan 1, 1964

By M. G. Benz, J. F. Elliott

The austenite solidus of the iron-carbon system has been determined using a series of diffusion couples, each of which consisted of a specimen of austenite held in contact with a melt saturated with austenite. After the equilibrium distribution of carbon had been established by diffusion at a specified temperature, i.e. the austenite specimen had become austenite of the solidus conzposition, the diffusion couple was cooled, sectioned and anazlyzed for carbon. The solidus was found to be a straight line: A revised iron-carbon temperature-composition diagram is presented, COMPOSITION has a marked effect on the temperature at which austenite begins to melt and the austenite solidus of the iron-carbon system describing this effect has been the subject of many investigations.1-10 The significant results of these investigations are presented in Fig. 1. The experimental methods and purity of the materials used in them are summarized in Table I. In plotting the data in Fig. 1 no attempt has been made to convert these results to the International Temperature Scale of 1948," except for the data of Adcock,5 as this conversion would do little to reduce the uncertainty that exists as to the position and shape of the solidus. A point of major concern in evaluating these data is that the alloys studied, except those used by Adcock, were not binary alloys of iron and carbon, Table I. Also, it would appear that several of the methods did not permit equilibrium to be established through the system being studied. All that can be said for the results of these investigations is that they indicate only approximately the location of the solidus with the uncertainty as to its location at 1 pct C being approximately 100°C. The current investigation was undertaken to provide reliable data by which the austenite solidus could be established. It was hoped that information on the liquidus also could be developed at the same time, but experimental limitations prevented this as austenite segregated from the liquid on cooling. After a careful study of possible experimental methods and extensive laboratory tests, it was decided that the use of a series of austenite-liquid diffusion couples would provide the most reliable results. This paper describes the method and its results and also includes a complete iron-carbon temperature-composition diagram based on what are considered to be the best available data. EXPERTMENTAL METHOD The diffusion couple used for this investigation consisted of a small cylindrical pellet of austenite held in contact, at a specified temperature, with a melt saturated with austenite. The composition of the cylindrical austenite pellet was chosen to be approximately 0.1 wt pct C less than the estimated

Jan 1, 1962

By Nicholas J. Grant, Bill C. Giessen

The phase diagram Ta-Ni has been treated repeatedly; investigations up to 1958 are summed up in Ref. 1. Since then, an equilibrium diagram has been presented by Kornilov and Pylaeva.2 They found the following intermediate phases: Ta2N; (CuA12 type),3 "TaNi" (p phase, Mo8Co7 type), TaNiz, and TaNi3 (Ticus type).5 Of these, the crystal structure of TaNin remained unknown, and no further literature reference was found. According to Kornilov and pylaeva,&apos; TaNil melts peri-tectically at 1420°C. A 5-g alloy of Ta + 66.7 at. pct Ni was prepared by inert-gas arc melting using 99.95 pct Ta rod (NRC-MG) and 99.9 pct Ni shot (Fisher). The weight loss after three melts was <2 mg; a microscopically homogeneous ingot resulted. As shown in Ref. 6, no oxygen or impurity pickup needed to be considered under these conditions. Specimens were annealed for 15 hr at 1350°C, 15 hr at 1205oC, and 15 hr at 875°C with subsequent powdering and stress relief for 10 min at 1350°C, 15 min at 1205oC, and 1 hr at 875oC, respectively. Micrographs and X-ray analysis (114.6 mm Norelco camera, filtered CuK, radiation) showed homogeneous one-phase alloys in the as-cast state and at all three temperatures, all having the same powder pattern except for broad lines in the as-cast state. From this it is concluded that either TaNi, melts congruently or its melting behavior is a transition between the eutectic and peritectic cases. Since it was found that very slight surface contamination on stress relieving can considerably alter the structure of similar phases (e.g.,TaPt3),7 patterns before and after stress relief were also compared; they were identical except for line broadening in the non-stressrelieved condition, and two faint impurity lines in the stress-relieved samples. The powder pattern is presented in Table I. All lines were accounted for by indexing with a tetragonal lattice, 0Yh- I4/mmm , with a, = 3.154 ± 0.001A, co = 7.905 * 0.002i and c/a = 2.506. The atomic positions are: 1000; 1/2, 1/2, 1/2] + 2Ta(a): 000; + 4Ni(e): *(002); z = 0.333 * 0.002 This proves TaNi2 to be a representative of the MoSiz (Cllb) type. The variable parameter was determined by grouping the visually estimated intensi- ties according to their 2 indices, plotting their F values, and varying z for best fit, which was obtained at z = 0.333= 1/3. Since absorption was not considered, only the relative intensities with similar 28 can be expected to agree. The interatomic distances are presented in Table 11; the standard error is *o.o2A, resulting chiefly from the uncertainty in z. These distances indicate a coordination number 10 for both kinds of atoms and show that the pseudohexagonal layers along (110) are rather un-distorted. The mean atomic volume of 13.1A3 is close to the Vegard&apos;s law value of 13.3A3, showing

Jan 1, 1964

By B. S. Tani, R. V. Schablaske, M. G. Chasanov

In the Ti-Cd system two intermediate phases, TiCd and Ti Cd, have been found and characterized by X-ray diffraction powder techniques. An earlier reference&apos; had reported the existence of the compound Tisd. No structural details, however, were given nor was there any indication of the possible existence of other binary compounds. Studies of the solubility of titanium in liquid cadmium2 led to the discovery of both TiCd and Tisd. The solid phase in equilibrium with saturated solutions of titanium in liquid cadmium at temperatures below 620°C is TiCd, whereas TizCd is the equilibrium solid phase above 620°C. Both thermal analysis and X-ray studies show that a peritectic reaction between the y phase (~isd) and cadmium to form the 6 phase (Ti Cd) occurs at 620°C. X-ray diffraction powder photographs of material isolated from the ingots used in the solubility measurements as well as material prepared by a direct heat of the appropriate stoichiometric amounts of titanium and cadmium can be indexed on the basis of a tetragonal unit cell with a = 2.904 ± 0.002A, co = 8.954 * 0.004A, c/a = 3.083. No variation of lattice parameters indicative of a range of composition is evident. With four atoms in the unit cell the calculated density of the intermetallic is 7.05 g per cm3; the observed value is 7.07 g per cm3. The TiCd lattice is primitive as indicated by the fact that there are no systematic absences among hkl reflections. The absence of all hko reflections for which h + k = 2n + 1 is also evident. This systematic absence indicates the following two space groups, C\h — P4/n and D&~4/nmm; both groups, however, have the same two- and four-fold positions. Stoichiometry, symmetry and space group indicate two likely structure types. They are the PbO-litharge3 (B10) type and the Ti CU (B11) type. Spatial considerations favor the B11 structure with atoms placed in the following positions: 2Ti (c) 0, 1/2, Z; l/2, 0, Z 2Cd(c) 0, 1/2, Z; 1/2, 0, Z The TiCd structure may be summarily described as a layered structure in which double cadmium layers alternate with double titanium layers along the c-axis. The best agreement between calculated and observed intensities is obtained for the values ZTi = 0.608 and ZCd = 0.130. Calculated and observed powder intensities as well as other pertinent powder data are recorded in Table I. Interatomic distances for the proposed structure are recorded in Table II and are to be compared with the distances of closest approach for titanium and cad-

Jan 1, 1962

By J. B. Darby

In an investigation of various equiatomic alloys of transition elements, Dwightl did not find evidence for the CsC1-type ordered structure in VMn. More recently waterstrat2 has reported the occurrence of this structure in the alloy his sample had been homogenized at 1150°C and subsequently reannealed at 725°C after crushing to powder. The purpose of this note is to confirm the existence of the CsC1-type ordered structure in VMn and to report a lattice parameter. An equiatomic alloy was prepared by arc-melting under an argon atmosphere vanadium and electrolytic manganese having a purity of at least 99.9 pet. The composition was obtained by charging an excess of manganese and monitoring the weight loss until the stoichiometric composition was achieved; the assumption that the loss during melting was due to the vaporization of manganese appears to be a valid one.&apos; The calculated composition of the sample prepared was 49.85 at. pct V and 50.15 at. pct Mn. Specimens taken from the fractured as-cast buttons were wrapped in molybdenum foil and encapsulated in quartz capsules under vacuum; there was no evidence of a reaction between the alloy and foil. The X-ray patterns were obtained using filtered copper or chromium radiation with a Straumanis-type Debye-Scherrer camera having a diameter of 114.6 mm. The CrKa radiation is best suited for detecting the superlattice lines since it overcomes the slight difference in scattering powers of the components by making use of the anomalies in atomic scattering factor which occur when the frequency of the radiation is close to the characteristic absorption frequency of one of the elements.394 A specimen in the as-cast condition, as well as one which had been annealed for 1 week at 900°C and water-quenched, gave a diffraction pattern for the disordered bcc solid solution. However, when samples were homogenized at 900°C for 3 weeks followed by a subsequent anneal at 700°C for 6 to 7 weeks and then water-quenched or furnace-cooled, the (loo), (1111, and (210) Ka superlattice lines were observed in films taken with chromium radiation. An ao = 2.904 * 0.001 was determined by a Nelson-Riley extrapolation from a film taken with CuKa radiation. The author wishes to thank Mr. J. W. Downey and Mr. L. J. Norton for carrying out the experimental portion of this investigation. This work was performed at the Argonne National Laboratory under the auspices of the U.S. Atomic Energy Commission.

Jan 1, 1963

By R. T. DeHoff

Saltykov&apos;s meihod for the determination of the size distribution of spherical particles from measurements made on random plane sections is extended to particle shapes which are not equiaxed. The analysis permits the deterrrzination of the size distribution of aggregates of ellipsoidol particles provided all particles have tlze same shape. V ARIOUS solutions to the problem of the determination of the size distribution of spherical particles from measurements made on random plane sections have appeared in the literature since the original solution was first published by Scheill in 1934. Solutions have been offered utilizing measurements of the diameter1y2 or area3&apos;4 of circular sections that result when a test plane is passed through an aggregate of spherical particles, and the length of intercept5y6 that results when a test line is passed through the aggregate. To the best of the knowledge of the author, no method has been proposed which permits the determination of the size distribution of particles that are not equiaxed. The purpose of this paper is to develop a method which permits the determination of the size distribution for a specific class of particle shapes which cover awide range of axial ratios; namely, ellipsoids of revolution. In the interests of simplicity, it is necessary to limit consideration to aggregates of particles which are all similar in shape; that is, the ellipses which, by rotation about one of their axes, generate the particles, all have the same axial ratio. The size distribution, which is to be determined by the analysis, is arbitrary. It is also assumed that the section, or sections, upon which the measurements are made represent the average of all possible sections through the structure. This ideal condition can be approached as nearly as may be desired by observing a sufficient area of sections taken at random positions in a randomly oriented aggregate, and in random positions and orientations in an aggregate that has one or more axes of orientation. SIZE DISTRIBUTION FOR PROLATE ELLIPSOIDS Consider an aggregate of cigar shaped, geometrically similar particles (prolate ellipsoids), dispersed in an opaque matrix. Let the size of each particle be specified by the length of its minor axis, By Fig. 1, and the shape, q, be defined by the ratio of its minor to its major axis, B/A. Let the index, j, refer to dimensions of ellipsoids in a particular size class, and the index, i, to dimensions of an ellipse resulting from the intersection of a test plane and a particle. The size of such a section is specified by the length of its minor axis, b. Let the particles be arbitrarily divided into k size classes, and let the increment of size between each class be A, where A = &,$, B,,,being the minor axis of the ellipsoids in the largest size class. Since all sections are represented on the test plane, the minor axis of the largest of the sections observed, results from the intersection of the test plane with an ellipsoid in the largest size class at its center, and is therefore equal to Bmax Thus the increment, A = The final result will consist of the k numbers which represent the number of particles in each size class per unit volume of aggregate:

Jan 1, 1962

By J. Szekely

The paper discusses simultaneous heat conduction and diffusion. These problems may arise in slag-metal kinetics and in connection with rapid heating or cooling of specimens. A mathematical formulation is given and it is shown that the transients due to either the variability of the diffusion coefficient with temperature, or thermal diffusion, are unlikely to be important for molten or solid metals and slags. The negligibility of these effects is due to the large values of the ratio thermal diffu-sivity/molecular diffusivity. In many problems relating to transport-controlled metallurgical kinetics the diffusion process is accompanied by the transfer of heat; large-scale metal refining or metal reduction may be quoted as examples. It has been shown that due to turbulent mixing in these systems the relevant transport phe- nomena are best described in terms of "surface renewal models" which require the solution of the unsteady-state diffusion equation.1-1 Problems involving simultaneous diffusion and heat conduction are also encountered in the laboratory in connection with rapid heating or cooling of specimens. The effect of variable temperature on diffusion may be twofold: a) in the absence of "cross effects" the problem, due to variable temperature, is that of diffusion with a variable diffusion coefficient; b) if the Soret Effect is nonzero, thermal diffusion must also be taken into account in the formulation. Although information is available on thermal-diffusion transients in a steady-temperature gradient,4&apos;5 little work has been done on the effect of unsteady temperatures on either the diffusion coefficient or on thermal diffusion.&apos; In the following a formulation is given and solutions presented with a view of examining the areas where these effects may be important.

Jan 1, 1965

By T. E. Leontis, R. S. Busk

WORK was initiated several years ago at The Dow Chemical Company to determine the applicability of powder metallurgy processes to magnesium and its alloys. Although it was found possible to apply the normal techniques of compacting and sintering, it was discovered early that the technique of extrusion offered special promise. This technique and the alloys produced were therefore studied extensively, and the present paper is a partial report of those studies. It is expected that a second paper will discuss more extensively some of the alloy systems simply mentioned here. Although the literature on powder metallurgy is very extensive, there are only two articles dealing with magnesium13,14 and there is no reference to the direct extrusion of the powder of any metal without the preliminary step of compacting. Furthermore, no mention has been made of the possibility of designing alloys by the technique of powder extrusion, particularly alloys that cannot be made in any other way. Sauerwaldl3 discusses the extrusion of magnesium and aluminum powder to produce alloys containing nonequilibrium phases, but again he emphasizes the importance of pre-compacting before extrusion. There seems to be general agreement that some form of hot-pressing is essential to obtain in a compact the full theoretical density of the metal used.&apos;-" The nearest approach to the technique of direct powder extrusion is that of extrusion of pre-sintered compacts,&apos;-" for which the best review through 1944 is that by Kieffer and Hotop.11 Ensign* discussed a specialized technique of "extruding" already compacted and sintered gears for the purpose of sizing. Sauerwa1d13 discussed the compacting and extrusion of magnesium and magnesium-aluminum alloy band-saw chips. The properties of the extruded compacts were found to be comparable with those of extrusions prepared from solid billets. He also extruded a compacted mixture of magnesium and aluminum powder and found the properties to be somewhat different from those obtained on an extrusion made from a cast billet of the same composition. The second reference dealing specifically with the powder metallurgy of magnesium is Stern&apos;s patent," in which is described the extrusion of scrap chips and turnings following a hot-pressing and sintering step. Considerable emphasis in most of these articles is placed on the necessity of atmospheric protection to minimize the danger of oxidation. Experimental Procedure The first experiments were made with the small extrusion press shown in fig. 1. With this press, billets of about 1/2 in. diam are used to produce wire

Jan 1, 1951

By G. Krauss, W. Pitsch

The fine structure of the bcc martensite formed in an Fe-33 wt pct ATi single crystal of arrstenite is sho~on by transmission electron microscoPy to consist of combinations of transformation twins, stacking faults, deformation twins, and regular arrays of parallel screw dislocations. These structures constitute evidence for the multiple lattice-invariant deformations which operating during the formation of martensite could produce the real habit-plane scatter measured by a two-surface analysis of the plates formed in the single crystal of this investigation and reported in the literature for other Fe-Ni rnartensites. CRYSTALLOGRAPHIC theories1,2 of martensitic transformation show that the habit plane of martensite in a parent lattice is dependent in part upon an inhomogeneous distortion or lattice-invariant deformation which takes place on a fine scale within a martensite plate during its formation. Several recent theoretical papers3,4 have addressed themselves to an analysis of a wide variety of conceivable lat-tice-invarient deformations and the habit planes which they produce, while experimental investigation have been concerned with either the measurement of habit planes or the description and identification of the martensitic fine structure which reflects the nature of the lattice-invariant deformation operating during transformation. In Fe-Ni alloys with subzero Ms temperatures, the group of alloys with which this paper concerns itself, habit planes have often been found to scatter an amount greater than might be expected from possible experimental errors,5-7 and fine twinning has been identified as a major constituent of the fine structure of martensite.8-11 It has been suggested3,4 that more than one type of invariant shear occurs during martensitic transformation. This possibility has been experimentally supported12,13 by the observation of both dislocation configurations and twinning in a single martensite plate. The purpose of this paper is to report additional evidence for multiple lattice-invariant deformations in martensite and so to account for the real scatter in the habit planes of the martensite plates formed in Fe-Ni alloys. EXPERIMENTAL PROCEDURE The Fe-Ni single crystal was produced by pulling a high-purity iron and nickel charge through a single-crystal vacuum furnace in an alumina crucible. The crystal was double-melted to promote homogeneity and to increase its size by further additions on the second pass. In its final form the crystal was 4 cm in diam and 5 cm long. The nickel and carbon contents were analyzed at 32.9 and 0.006 wt pct, respectively. The austenite of this alloy first transformed to martensite by bursts at about -120°C, and, to preserve as much of the austenite as possible, all transformation was performed just below -120°C. Some observations were made on transformed samples which had been heated for 2 min at 340°C. It is assumed that the features of the martensite of these samples, Figs. 1 and 4, are the same as those of the as-quenched martensite. Orientation of the crystal by X-ray diffraction established 10.735 0.609 0.3161? as the axis of the crystal, an orientation that was checked within 2 deg by neutron diffraction. Further checks by electron diffraction of samples cut normal to the axis confirmed this orientation within the larger limits of error inherent in electron diffraction of thin foils. The X-ray orientation was the one used for the two-surface analysis of the martensite habit planes. A two-surface analysis was performed on the quadrant of the single crystal which had been oriented by both X-ray and neutron-diffraction techniques. Photomicrographs at X50 were made on two surfaces along an edge 2 cm long. Fiducial marks and the fact that many of the plates were almost completely surrounded by retained austenite made good matching of individual plates on two surfaces possible. The habit-plane trace on a surface was taken as the best line parallel to the long axis of a plate. A measure of the accuracy afforded by this criterion was provided by a family of very large plates which appeared at intervals along the entire 2 cm length of the edge. The plates all had habit-plane traces within 2 deg of one another. Many of the plates did not show midribs and, therefore, the use of midribs7 to represent habit-plane traces was not feasible in this investigation. The over-all experimental accuracy is estimated to be better than ±2 deg. Samples for transmission examination in a Siemens Elmiskop I at 100 kv were prepared by cutting 2-mm-thick discs from the single crystal, removing about 0.5 mm by chemical polishing,14 trans-

Jan 1, 1965

By J. C. Fisher

Martensite transformations in steels and other alloys are characterized in part by the absence of composition changes during the growth of a new phase. Transformation occurs rapidly, and there is insufficient time for long range diffusion or partition of alloying elements to take place; martensite reactions in alloys thus are similar to phase transformations in single component systems. A fundamental understanding of martensite transformations in steels is impossible without knowledge of the free energy change upon transforming austenite (face centered cubic iron containing alloying elements) to ferrite (body centered cubic iron containing alloying elements) of the same chemical composition. The present paper assembles the best information available concerning the influence of temperature and composition on this free energy change. Most of the material has been taken from the work of Johansson,&apos; Mehl and Wells,2 Zener3,4 and Smith;5 and indirectly, through these authors, from the work of Austin.6 In agreement with the generally accepted viewpoint, martensite is assumed to be an ordered solution of carbon in ferrite of the same composition as the parent austenite; only at high temperatures and low carbon concentrations is the carbon in ferrite distributed at random. The properties of the disordered solution are estimated by extrapolating the known properties of iron-carbon solid solutions into the range of supersaturation, and the free energy change associated with ordering is estimated from the theory developed by Zener. By incorporating Smith&apos;s recent thermodynamic measurements and Zener&apos;s theory of ordering, the present analysis modifies previous estimates of the free energy change associ- ated with martensite transformations. Consider a two component system consisting of a solvent A and a solute B. Let Na and Nb represent mol fractions of A and B respectively, let aa = raNa and ab = YbNb represent activities, and let superscripts 1 and 2 refer to phases 1 and 2. The partial molal free energies of A and B in phases 1 and 2 can be summarized as follows: free energy standard state Fa&apos; - Fao1 = RT In an1 pure A&apos; Fa2 - Fa2 = RT In aa2 pure A2 Fb1 - Fbo = RT In ab1 pure B Fs2 - FB2 = RT In aB2 pure B. The free energy of a gram atom of phase 1 is* AF1 = Na&apos;Fa1 + Nb&apos;Fb1 and that of phase 2 is AF2 = NA2FA2 + NB2Fb2. A martensite transformation from phase 1 to phase 2 requires Na1 = Na2 = Na and NB1 = NB2 = Nb, and the free energy change per gram atom accompanying transformation is AFi-2 = NA(Fa2 - Fa1) + Nb(Fb2 - Fb1) = NA[RT In (aA2/aA1) + AFa1?2] + Nb RT In (aB2/aB1) = Na[RT In (ra2/ra1) + AFa1?2] + Nb RT In (TaVTB1). [1] where yn2, yn1, yB2, yB1 are activity coefficients, and where Ma&apos;+* is the free energy change upon transforming a gram atom of pure A from phase 1 to phase 2 at the temperature in question. For martensite transformations in plain carbon steels, A = iron (Fe), B = carbon (C), 1 = austenite (7). 2 = ferrite (a), and Eq 1 is AFr ?a= NFe[RT In (rf.a/rfer) + AFF.r?a] + TVc RT In (eca/rc1)- L2] Nothing is known concerning the values of yFea and yca for carbon concen- , trations in excess of 0.025 pct. However, the approximation rf.7 = 7f." = 1 cannot be appreciably in error for small carbon concentrations, and Eq 2 reduces to Afr-a = NfeFfer?a1 + NC RT In (rca/rcr)- [3] Johanssonl and Zenera have calculated MFfr?a from the specific heat measurements compiled by Austin.6* Their calculated values agree closely, and are summarized in Table 1. The activity coefficients relative to graphite for carbon dissolved in iron vary with temperature according to the relationships d In rcr/d(1/T) = AHcr-R dlnyca/d(l/T) = AHCa/R where AHc? and AHca are the heats of solution of graphite in y and a iron. Assuming the values of AH to be independent of carbon concentration and temperature, In rCr = AHcr/RT + I1 In rca = AHCa/RT + 12

Jan 1, 1950

By R. W. Heckel, R. L. DeGregorio

The DeHoff method of determining the size distribution of ellipsoid-shaped, second-phase particles has been applied to the spheroidization of cementite in a eutectoid steel. The surface area of Precipitates determined from the various size fractions in the distribution was correlated with surface-area measurements based on the number of intersected interfaces on a random lest line. The precipitate particles were found to be oblate ellipsoids with an axial ratio of about 0.90. The size distributions were found to be log-normal. A method is proposed whereby the shrinkage of small particles and growth of large ones can be determined from the experimental data. The experimental data are compared to previously proposed mathematical models describing diffusion-controlled kinetics and various types of interface-controlled kinetics. The experimental growth and shrinkage rates are considerably slower than those predicted by diffusion-controlled kinetics. The best fit is obtained for a model describing interface-controlled kinetics limited by the rate of formation of cementite at the growing interfaces where the interfacial reaction rate is proportional to the solute thermodynamic activity gradient across the surface. THE subject of spheroidization of second-phase particles has been considered previously by other investigators. Livingston1 has studied the precipitation of the cobalt-rich phase from copper alloys containing 0.7 to 3.2 wt pet Co. His results indicate that the average particle size increases as the cube root of the heat treatment time in accord with diffusion-controlled kinetics.&apos;-&apos; Komatsu and rant&apos; in their studies of the growth of SiO2 in a dispersion-strengthened copper-silica alloy found that the initial growth of SiO2 proceeded by diffusion-controlled growth and later stages were limited by interface-controlled growth. The values of the activation energy obtained for the interface- controlled process led them to the conclusion that the process was limited by the dissociation of SiO2 at the shrinking interface. The transition from diffusion-controlled growth to interface-controlled growth was characterized by a particle size which varied with the heat-treatment temperature. It is also interesting to note that the particle-size distribution as found by Komatsu and Grant exhibits log-normal behavior when replotted on log-probability coordinates. Dromsky, Lenel, and Ansell9 have observed the growth of Al2O3 particles having a mean free path of about 1.5 to 13 µ. Their photomicrographs indicate that the Al2O3 particle sizes were larger than those observed by Komatsu and Grant. Dromsky, Lenel, and Ansell concluded that the Al2O3 coarsened by an interface-controlled growth mechanism limited by the solution of A12O3 at the shrinking interfaces. Bannyh, Modin, and Modin10 have studied the spheroidization of a eutectoid steel (0.83 wt pet C). Their spheroidization (tempering) treatments were carried out in the range from 210o to 700°C for times between 1.5 sec and 20 hr. Measurements of mean particle size as a function of time indicated a cube root of time dependence of size at 700°C, in agreement with previous analyses of diffusion-controlled kinetics.2°7 At lower temperatures, the time dependence was less than the one-third power. It is important to note that, although mathematical models of growth* considering particle-size distribution have been available, measurements of only mean particle size have been carried out. DeHoff11 has presented a quantitative metallography technique which is applicable to the determination of the size distribution of ellipsoids of constant shape. This method is applicable to both oblate and prolate ellipsoids of all ratios of minor to major axis and is based upon an extension of Saltykov&apos;s analysis12 of the distribution of spheres of varying size. DeHoff&apos;s analysis is based upon measurements of the minor axis of the particles on a random plane of polish in a unit area. This method provides a measure of the number of ellipsoids in various size ranges per unit area. As pointed out by DeHoff, such measurements may be used to obtain surface area and total precipitate volume data. Thus, the accuracy of the distribution analysis may be checked by comparison of the surface and volume

Jan 1, 1965

By D. Mizer, J. B. Clark

In a recent investigation of the entire magnesium-yttrium phase diagram, Gibson and Carlsonl report the maximum solid solubility of yttrium in magnesium as 8.0 wt pct Y at 1050°F (565°C). However, independent studies of the precipitation in magnesium-yttrium alloys, 2 conducted prior to the publication of the above phase diagram, indicated a greater solid solubility of yttrium in magnesium. Accordingly, the magnesium-rich region of the diagram was re-investigated to determine more accurately the liqui-dus, solidus, and solvus of the magnesium solid solution in this system. Metallography and thermal analysis were chosen as the best corroboratory methods for accurately determining this region of the diagram. A magnesium-yttrium master alloy was prepared from yttrium of 99.9 pct purity and singly sublimed magnesium of 99.9 pct purity. Subsequently, the twelve magnesium-yttrium alloys shown in Fig. 1 were prepared using standard alloying procedures for magnesium. Metallographic samples of the alloys were heat treated under an SO, atmosphere in an electric furnace capable of control within ±3oF of the desired temperature. Alloy samples for the metallographic determination of the solvus, and for the precipitation studies were solution treated at 1000° F for 16 hr. Subsequent aging times for the solvus determination varied from a minimum of 6 hr for the higher temperatures up to 816 hr for lower temperatures. Periods of 30 min at temperature were used to generate liquation for the metallographic placement of the solidus and eutectic horizontal. The liquidus and the eutectic horizontal were determined by thermal analysis. Simple time-temper- ature heating and cooling curves were obtained from 10-g alloy samples heated and cooled in a well-insulated electric furnace. An Iron-Constantan thermocouple was placed in direct contact with the sample. The thermocouple did not contaminate the alloy sample since metallographic examination showed that the alloy melt did not wet or attack the thermocouple. The heating and cooling rates were approximately 5 to 8OF per min. The standard metallographic procedures for magnesium were used in polishing and etching the magnesium-yttrium alloys studied. Glycol etchant* was etchant was used to reveal precipitate in the solvus determination. The redetermined magnesium-rich region of the magnesium-yttrium diagram is shown in Fig. 1. The

Jan 1, 1962

By S. T. Wlodek

The surface and subscale oxidation reactions were followed by means of continuous weight-gain and metallographic techniques over the range 1600" to 2200°F (871° to 1204 °C) for up to 400 hr. Full identification of all scale and subscale reaction products was obtained by electron and X-ray diffraction. At or below 1800°F (982°C) a linear rate of reaction (QL = 46.0 kcal per mole) governed the oxidation process, extending for up to 100 hr at 1600°F (871 "C). During linear oxidation the surface scale consisted of an amorphous SiO2 film overgrown with Cr 2O 3 and NiCr204. This initial linear process was followed, and above 1800°F completely replaced, by two successive parabolic rate laws (Qp = 60 and 57 kcal per mole). This parabolic reaction involved the formation of a complex scale consisting of Cr2 O3 and smaller amounts of NiCr2O4. Parabolic oxidation appeared to coincide with the disruplion of the silica film present during linear oxidation and was followed by subscale (internal) oxidation of crystobalite and NiCr2O4. The balance between the subscale and surface oxidation reactions controls the oxidation of this commercial alloy. The amorphous silica film appears to result in the linear rate and diffusion through Cr2O3 is the more likely rate-limiting step during parabolic oxidation. THE oxidation of a multicomponent composition is a complex phenomenon not presently amenable to a rigorous classical interpretation. Nevertheless, even a qualitative understanding of the scaling and subscale reactions that occur in a commercial composition can illuminate the reactions that limit its high-temperature stability in an oxidizing environment. This study of the oxidation of Hastelloy Alloy X presents the first of a series of studies with the above approach in mind. Hastelloy X exhibits one of the best combinations of strength and oxidation resistance available in a wrought, solution-strengthened, nickel-base alloy. Although during long time exposure some precipitation of M6C and M23C8 carbides as well as a complex Laves phase occurs, the amounts are probably small enough to have no appreciable effect on the chemistry of the matrix. Radavich has identified the oxidation products on Hastelloy X oxidized for 5 min to 10 hr at 1115°F as NiO and the NiCr2O4 spinel. Oxidation for 5 to 15 min at 1500°F produced a scale of spinel, NiO, and a rhombohedra1 phase, probably Cr2Os. Sannier et 2. have reported continuous weight-gain data for Hastelloy X at 1650" and 2010°F and internal-oxidation measurements after 150 hr at 2010°F. In addition, much of the data on binary Ni-Cr alloys recently reviewed by Kubaschewski and okins&apos; and Ignatov and Shamgunova4 as well as studies of binary Ni-Mo alloys5 are also pertinent to the oxidation of this composition. EXPERIMENTAL Continuous weight-gain measurements and metallographic measurements of subscale reactions were the main experimental techniques used in this study. X-ray and electron diffraction backed up by a limited amount of electron-microprobe analysis served to characterize the nature of the scale- and subscale-reaction products. Two heats of commercial sheet of the composition given in Table I and identified as A and B were used in the bulk of this study. Internal-oxidation measurements were made on a third heat of material in the form of a 0.5-in.-diam bar. In order to assure homogeneity, all heats were reannealed 4 hr at 2175°F prior to sample preparation. weight-Gain Measurement. All specimens (1.5 by 0.4 by 0.03 in.) were abraded through 600 paper, electropolished, and lightly etched in an alcohol-10 pct HCl solution. An electrolyte of 150 cu cm H,O, 500 cu cm HsPO4 (85 pct conc), and 3 g CrO3 at a current density of 0.9 amp per sq cm or a solution of 10 pct HaW4 in alcohol used at 4 v and 0.3 amp per sq cm was used for electropolishing. The resultant surface exhibited a finish of 3 ± 1 p rms. Continuous weight-gain tests were made at 1600°, 1700°, 1800°, 1900°, 2000", and 2200°F on auer&apos; type balances capable of recording a total weight change of 110 mg with an accuracy of k0.1 mg. All tests were made in air dried to a dew point of -70°F and metered into the 2-in.-diam reaction

Jan 1, 1964