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By Paul A. Farrar, Harold Margolin

The system Ti-Cr has been re investigated From 40 to 100 pct Cr. The high-temperature modification of TiCr, was found to exist in the temperature range from 1365 where it terminates in the eutectoid reaction TiCr,(TI)* — HE chromium-rich region of the Ti-Cr system was investigated by Vogel and Wenderott,' Mc Quillan,' Van Thyne et dl.,' and Cuff Duwez and Taylor determined the crystal structure of the low-temeperature modification of TiCr, to be fcc a = 6.943A and Levinger' discovered the existence of a second (high-temperature) modification of TiCr, and determited its crystal structure to be hexagonal a = 4.932A, c/a = 1.617i. No information was obtained as to the temperature range over which each modification of TiCr? existed; therefore the present investigation was initiated. I) Experimental Techniques Alloys employed for the delination of this system were prepared from iodide titanium and high-purity chromium. The purity of these materials is as follows: Iodide titanium: purity 99.99 pct; major impurities—carbon 0.0001 pct; oxygen 0 to 0.002 pct; nitrogen 0 to 0.002 pct; silicon 0 to 0.005 pct; manganese 0.001 pct; copper 0.001 pct (Foote Mineral Co., Lot 609-29). Chromium #1: metallic purity 99.997 min.; interstitial content—oxygen 0.27 to 0.46 pct; hydrogen 0.028 to 0.038 pct; nitrogen 0.0002 pct (Ave-quipo Inc.). Chromium #2: purity 99.9 pct; major impurities—carbon 0.01 pct; oxygen 0.013 pct; nitrogen 0.02 pct; silicon 0.01 pct; sulfur 0.03 pct; iron 0.01 pct (Fulmer Research Institute Ltd.). A total of sixteen alloys of thirteen different compositions was prepared in the range from 40 to 95' *All percentages are given in weight percent. pct Cr. The compositions of the alloys prepared were as follows: Ti-40, 45, 50, 60, 62, 64, 66, 68, 70, 80, 85, 90, and 95 pct Cr. Two buttons were prepared of the 66, 90, and 95 pct Cr alloys. The 40 pct Cr and one of the 66 pct Cr alloys were prepared from the chromium # 2. All the other alloys were prepared from chromium #1. The alloys were arc-melted as 25-g buttons in a copper-hearth glass-walled furnace7 under an inert atmosphere of high-purity helium. The alloys were sintered at 250 amp and melted twice at 350 amp for 30 sec with the buttons being turned over between melts. Weight losses during melting were determined and in the range 45 to 68 pct Cr they were found, in all cases, to be less than 1.25 pct of the total alloy weight. A number of the high-chromium alloys shattered during cooling from either the first or second melt, which made the weight loss considerably larger than usual. The 70 and 80 pct alloys lost approximately 5.7 pct of their total weight, the 85 pct alloys 1.5 pct, the 90 pct alloy 2 pct, and the 95 pct alloy 3 pct. However, it is felt that the actual composition of the alloys in this region is probably within approximately 1 pct of the intended composition, as a large portion of the material lost was that due to shattering after melting. Nominal compositions were therefore used in plotting data and referring to the alloys. The alloys were cut into specimens, wrapped in molybdenum sheet, and encapsuled in argon-filled quartz or Vycor tubes for heat treatment. The heat-treatment times and temperatures are given in Table I. The specimens were quenched in iced brine (I.B.Q.) following heat treatment. The standard techniques of grinding, polishing., and etching were used in preparing the specimens for metallo-graphic examination. For alloys below 68 pct Cr

Jan 1, 1963

By Wilhelm in der Schmitten

A necessary condition for a uniaxial stress distribution in compression testing is that the specimen end surfaces make full contact with the apparatus compression plates. In addition all compression surfaces must be machined perpendicularly to the loading axis. In general this condition is not sufficiently fulfilled in a simple apparatus. Therefore on application of the load the compression plate may first touch the specimen at one point only. This may cause crumbling of the specimen if the material under investigation is brittle (e.g., ionic crystals). It also precludes a precise study of the beginning of deformation as the effective cross sectional area is unknown. In addition a complex stress state is set up which renders quantitative work impossible, e.g., on dislocation density and arrangement as a function of stress. To overcome these difficulties a "self-adjusting" upper compression plate is required. A simple and reliable method of achieving this is shown in Fig. 1: A steel hemisphere is placed on the specimen. As soon as the compression plate (which no longer needs precise machining) makes contact with this hemisphere the load is uniformly distributed over the whole specimen surface, and a homogeneous uniaxial state of stress results. An advantage of the method is that the loading is always central, even when the specimen surface is not parallel to the compression plate (Fig. 1, dashed line). Therefore to a first approximation a uniaxial state of stress is maintained if the misorientation of the specimen end surface is small. Two typical stress strain curves of identically prepared sodium chloride crystals are shown in Fig. 2; curve 1 was taken without and curve II with the hemisphere under otherwise identical conditions. The improvement in the accuracy of defining the beginning of elastic and plastic deformation is clearly indicated. The writer wishes to thank the Deutsche Forschungsgemeinschaft for financial support, and Professor P. Haasen for providing facilities for this work.

Jan 1, 1962

By P. M. Kelly

TWO analytical methods for calculating the positions of matrix and twin reflections on electron-diffraction patterns of a twinned cubic crystal have recently been published.1,2 Meieran and Richman1 considered the case of twinning on (1 12) in a bcc structure, while a more general treatment, capable of analyzing twinning on any plane in a cubic crystal, was developed by Johari and Thomas.2 The latter authors applied their general rotation matrix to both (111) twinning in face-centered crystals and {112} twinning in body-centered crystals and they derived separate expressions for these two cases. This separate treatment masks the fact that in a holosymmetric cubic crystal a (111) twin is crys-tallographically equivalent to a {112} twin.3 This is analogous to the crystallographic equivalence of rotation and reflection twins in centrosymmetric crys tals. A particularly simple example of this equivalence occurs when the two twin planes are perpendicular to each other—for example (111) and (112). Then, if twinning on (111) moves a point P to a new position P1, and twinning on (112) moves P to P2, P1 and P2 will be related to each other by a 180-deg rotation about [110]— the direction normal to both (111) and (112). Since there are diad axis axis along (110) direction in a holosymmetric cubic crys tal the two positions P1 and P2 must be crystallo-graphically equivalent. In other words if a particular pole (hk l) is in a position P1 after twinning on (Ill), then after twinning on (112) a pole of the type {hkl}—but not the particular pole (hk1)—will also be in the position P1. As a result, when only the general (as opposed to the particular) indices of poles after twinning are required, twinning on (111) gives the same results as twinning on (112). To convert the results for a (111) twin to those applicable to a (112) twin, when the particular (as opposed to the general) indices of poles after twinning are required, all that is necessary is to multiply the indices of the twin poles by the matrix for 180-deg rotation about [lie], i.e., The analytical methods suffer the disadvantage that a new set of calculations must be performed for each different diffraction pattern or, as is done in the Meieran and Richman treatment,1 a lengthy list of the location of the twin spots in reciprocal space must be prepared. All the tedious computation involved in these analytical approaches can be avoided by simply plotting, on tracing paper, a stereographic projection of a crystal twinned on say (112). By superimposing this twin stereogram on a standard (001) projection (preferably with the points and indices in a different color from those on the twin stereogram) the relative positions of both matrix and twin reflections can be determined by inspection. Alternatively, a composite projection can be

Jan 1, 1965

By P. Schwarzkopf, J. H. Brophy

A modification of the Mendenhall wedge blackbody1 has been used to determine solidus temperatures and to anneal alloys in the tantalum-rhenium binary system. The technique has proven to be simple and accurate with modest power requirements. For temperatures up to 2800° C a piece of 10-mil tantalum sheet was folded double and formed as shown in Fig. 1. The middle of the resulting ribbon filament was a cylindrical-shaped crucible with less than a 1/8-in. opening. After inserting a specimen in the cavity, the filament was clamped between molybdenum electrodes in a high vacuum chamber. A representative filament 2 in. long and % in. deep was heated to 2800°C with a current of 700 amp at 4 v. The specimen temperature was determined by sighting into the cavity with a Leeds and Northrup optical pyrometer. The twisted filament shape was necessary to avoid serious geometric changes during thermal expansion. For temperatures between 2800" and 3000°C a 4-mil tungsten filament was employed. To facilitate fabrication it was warm-folded, and a twisted segment of 10-mil tantalum was used to support one end to absorb thermal expansion. This element assembly required about 700 amp at 10 v to reach 3000°C. Two geometric modifications of the filament were helpful for solidus temperature determinations. A shallow chamber, with depth to opening ratio of 3 to 1, permitted direct observation of the specimen through the optical pyrometer. The fact that the specimen was visible permitted the melting point to be determined when sharp features on the specimen surface became rounded; it also indicated that blackbody conditions did not prevail. By standardizing against the melting point of molybdenum, the observed melting temperature was corrected by the following form of Wien's Law: in which T is the true temperature, To is the observed temperature, and K is a constant characteristic of the filament geometry and determined experimentally for the molybdenum standard. The resulting corrected melting point for each alloy served as a reference temperature for the second more precise filament geometry. With a 6 to 1 depth to opening ratio, the specimen was indistinguishable from the filament interior, indicating good blackbody conditions. Using this technique, a series of specimens was isothermally treated at temperatures above and below the corrected direct reading. Incipient melting was detected by visual or micrographic examination of the specimen after cooling. This procedure was checked with specimens of pure chromium, molybdenum, and tantalum. Table I represents the values obtained, together with the solidus temperatures in the tantalum-rhenium binary system determined by the same technique. The alloy specimens were approximately 1/8 in. in diam. They were pieces of 10-g buttons melted four times in a nonconsumable tungsten electrode arc furnace on a water cooled copper hearth in 2/3 atmos of titanium gettered helium. Several sources of error are possible in such temperature measurements. Within the accuracy of the pyrometer no gradients were detectable near the specimen, and element distortion was virtually eliminated by twisting the tantalum. Compensation was made for the 1/8-in. sight glass by inserting several identical thicknesses and extrapolating to zero thickness. As a result an additional 10°C correction was added to all observations in the temperature range of 2500" to 3000°C. The principal source of error was then in the operation of the

Jan 1, 1961

By B. Post, D. Moskowitz, F. W. Glaser

Lattice parameters of various hafnium compounds, two borides, a carbide and a nitride were determined. Some physical properties were measured and comparisons with the properties of the corre-sponing zirconium compounds were drawn. A LTHOUGH borides of transition metals have A been extensively investigated in recent years, there are few references in the literature to borides of hafnium. This is not surprising, in view of the extreme scarcity and great cost of the purified metal. Becker' reported the existence of a boride, carbide, and nitride of hafnium; he listed the formulas as HfB, HfC, and HfN. However, in his work no mention was made of chemical analyses of the reaction products. In this paper results of a study of various hafnium compounds, including borides, carbides, and nitrides, are reported. It was expected, of course, that these compounds would closely resemble their zirconium analogues, which have previously been extensively investigated. Three borides of zirconium are known: ZrB,² ZrB,² and ZrB212.4 In addition, one carbide, ZrC,5 and one nitride, ZrN,5 have been reported. It should be noted that ZrB and ZrB12 are stable phases only in relatively high-temperature regions, unlike ZrB2, which is stable from room temperature to its melting point." It was felt that hafnium would yield corresponding borides, carbides, and nitrides with similar stability ranges. Experimental Procedures and Results A quantity of hafnium metal that was better than 97 pct pure, with zirconium as its major impurity, was furnished by the Argonne National Laboratories. The boron powder used was approximately 96.5 pct pure. In general, the procedure involved mixing the metal and boron powders in the desired proportions, and subsequently hot pressing them. The resulting specimens were analyzed chemically and phases present were identified by X-ray diffraction methods. A Philips Norelco Geiger counter spectrometer was used for identification of components and for measurement of lattice constants. Filtered copper radiation was used for all X-ray work. Electrical resistivity and density measurements were carried out at room temperature. The stable diboride, HfB2, will be described first. Stoichiometric quantities of boron and metal were heated to approximately 2000°C during hot pressing. The reaction product was shown by X-ray diffraction methods to be isomorphous with ZrB2 (space group Dlgh; P6/mm). Values of lattice constants of HfB2 compared with those of ZrB2 are shown in Table I. The calculated density of HfB2 is 11.2 g per cc, compared with a measured (pycnometric) value of 10.5 g per cc. HfB2 specimens prepared under a wide variety of conditions (including variations in initial Hf-B ratios) showed only insignificant variations in lattice constants, indicating a very narrow range of homogeneity. Like ZrB2,7 HfB2 is an excellent electrical conductor. The resistivity at room temperature of hot-pressed samples, approximately 85 pct dense, is 12 microhm-cm. This compares favorably with the electrical resistivity of 10 mi-

Jan 1, 1954

By B. R. Banerjee

MATERIAL used throughout this investigation was high-purity aluminum (99.998 pct). The 1/2-in. cubes were cut out of a cold-rolled slab and annealed at 550°C for 1 hr before deformation. The single crystal specimen was 1 in. in diam x 0.33 in. thick, cut from a cylindrical crystal produced by cooling at about 55°C per hr in a furnace which had a convenient gradient of about 10°C in every 4.5 in. vertical crucible length. Specific data on the ten specimens studied are summarized in table I. Three types of deformation were employed: (I) compression in a hand-operated hydraulic press; (2) impact with a 7-lb sledge hammer; and (3) rolling, using a wedge-shaped specimen' to obtain a gradient of deformations varying from 0 to 50 pct along the length of the wedge. Specimens were either deformed at room temperature, or first cooled in a freezing mixture of dry ice and ether at a temperature of —72°C and then deformed between steel blocks cooled with dry ice. The specimens deformed at the low temperature were ground and polished through 000 emery paper and wheels impregnated with alundum and alumina No. 1. During this entire procedure the specimens were dipped frequently in the freezing mixture of dry ice and ether in order to maintain constantly a frosty surface as insurance against overheating. All specimens were given a final electrolytic polish, using two different electrolytes* to insure that the surface effects observed were not char- acteristic of the polishing technique. With both electrolytes, the annealed specimens gave even surfaces with good polish. Several standard etchants were tried without success. The best results were obtained with an immersion etchant, developed in this connection, containing: hydrofluoric acid, 4 parts; nitric acid, 1 part; glycerine, 3 parts. Discussion of Results Typical strain markings have been observed within deformation bands, as shown in figs. 1, 2, and 4. By stereographic plotting of metallographic measurements, the strain markings within a deformation band in a single crystal of aluminum deformed by compression were found to be traces of (111) planes within a maximum scatter of 6". The edges of the deformation band were found to be traces of the (100) planes also within the maximum scatter of 6" (in agreement with previous workers). At 30 pct deformation, this scatter was considered quite normal because of the fragmentation and rotation of the crystal. This solution, based on the traces on one plane of observation is not unique, but as the planes were observed to be those normally and logically to be expected, further work in this regard was not felt to be necessary. The study of strain markings in a wedge-shaped specimen, rolled to obtain a gradient of deformation from 0 to 50 pct, revealed prominent strain markings at 25 pct reduction in thickness (fig.5). Faint and weak strain markings were visually observable up to a minimum reduction in thickness of about 12 pct. Since strain markings have been observed in other face-centered cubic metals after much smaller deformations;1, 2 and since the rate of recovery in aluminum is high, it is believed that, at lower deformations, self-recovery removes the localized strains at the slip interfaces, thus removing the source of the etching effects that give rise to strain markings. At high magnifications, strain markings sometimes seem to appear as very thin, light-banded re-gions bounded by a pair of dark edges. This so-

Jan 1, 1951

By E. J. Rapperport, E. D. Levine

The boundaries of the (a +ß) field in the Al-Li system were determined between 150°and 550°C utilizing quantitative metallography and lattice-parameter measurements. The solubility of lithium in aluminum decreases from 12.0at. pct Li at 550°C to 5.5 at. pct Li at 150°C. P Al-Li is saturated with aluminum at 45.8 at. pct Li and has this boundary value constant over the temperature range 150°to 550°C. THE solid solubility of lithium in aluminum has been determined by several investigators, 1-6 but, as shown in Fig. 1, there is little agreement among the various determinations. The earliest investiga-tions'-' are suspect because of the use of impure materials. Although high-purity materials were employed in more recent work,4'5 the experimental techniques may have led to contamination of the specimens. Probably the best work has been that of Costas and Marshall,6 who obtained close agreement between results obtained by two independent phase-boundary techniques: electrical resistivity and mi-crohardness. No detailed studies of the solubility of aluminum in the bcc ß phase, Al-Li, have been reported. Cursory investigations1,2,6 have indicated only that the (a+ß) -p boundary lies between 40 and 50 at. pct Li and is relatively independent of temperature. The present work was undertaken in order to provide an independent check on Costas and Marshall's determination of the solubility of lithium in aluminum, to extend knowledge of this solubility limit to temperatures below 225°C, and to make an accurate determination of the solubility of aluminum in Al-Li. EXPEFUMENTAL Alloy Preparation. In view of the difficulties encountered in previous investigations of the A1-Li system, close attention was paid to the use of methods of alloy preparation and treatment that would minimize contamination. Aluminum sheet (99.99 + pct Al) was vacuum-induction melted in a beryllia crucible to remove hydrogen. Lithium (99.9 pct Li) was charged with pre-melted aluminum into a beryllia crucible, in a helium-filled drybox. The crucible was sealed in a Vycor tube and transferred from the drybox to an induction furnace. Melting of alloys was performed by induction heating in a helium atmosphere. Solidification was accomplished by means of a suction apparatus, shown in Fig. 2, in which the alloy was forced by changes of pressure into a 3/16-in. inside diam closed-end beryllia tube. This technique produced rapid solidification of a small portion of the melt, resulting in alloys with a high degree of homogeneity. Typical lithium distributions are presented in Table I. Transverse sections 1/8 in. long were cut from the alloy rods, and each section was split in half longitudinally. One half of each section was analyzed for lithium, and the opposing halves were employed for phase-boundary determinations. Lithium contents were determined by flame photometry with an accuracy of 1 pct of the amount of lithium present. Thermal Treatments. Homogenization and equilibration heat treatments were performed in electrical-resistance furnaces with temperatures controlled to ± 2OC. Calibrated chromel-alumel thermocouples were employed to measure temperature. Homogenization was performed in helium-filled l?yrex tubes for 1 hr at 565°C. The encapsulated specimens were then transferred directly to furnaces maintained at lower temperatures for equilibration. Equilibration times were 2 hr at 550°C, 8 hr at 450°C, 27 hr at 350°c, 90 hr at 250°c, and 285 hr at 150"~. These times were chosen on the basis of conditions employed by previous investigators. Alloys were quenched from the equilibration temperatures by breaking the capsules into a silicone oil bath. By performing all possible operations either in sealed capsules or in a helium-filled drybox, the specimens were given minimum exposure to the atmosphere. Quantitative Metallography. Metallography of Al-Li alloys is difficult because of the atmospheric reactivity of the ß phase. It was found possible, however, to prepare surfaces of good metallographic quality by preventing contact with moisture during preparation. Grinding through 4/0 paper was performed in the drybox. The specimens were then transferred under kerosene to the polishing wheel. Three polishing stages were employed: 25-p alundum with kerosene lubricant on billiard cloth, 1-µ diamond paste on Microcloth, and 1/4-p diamond paste on Microcloth. Between stages the samples were cleaned by rinsing in trichloroethylene and buffing

Jan 1, 1963

By J. W. Pugh, J. D. Nisbet

THIS study of the ternary has been made as one phase of a metallurgical investigation which began nearly four years ago in the General Electric Company's Research Laboratory in Schenectady, N. Y. The objective of this program is the discovery of new metallurgical information which will lead to the development of better high-temperature materials. Combinations of the four pure base elements—iron, chromium, nickel, and cobalt—are being studied at the present time. It is essential in an investigation such as this to know as much as possible about the constitutional diagrams involved. The study of the iron-chromium-nickel system has been made to this end. Preliminary Explanation of Experimental Procedure: In all, fifty-five alloys at steps of 10 at. pct were made for the investigation of this ternary. Several ternary alloys in the chromium rich corner had to be omitted because their extreme brittleness made testing impractical. All alloys were vacuum melted with hydrogen reduction and centrifugally cast. The apparatus and technique of this process has been described in detail by Nisbet.1,2 The purity of the alloys prepared in this way is considered to be quite good. Impurities are listed as follows: Pct Carbon.............................0.02 Oxygen................................0.02 Nitrogen.........................0.005 Magnesium..................0.03-0.05 Sulphur...........................trace Hydrogen....................trace Phosphorous..................trace Silicon........................ trace After casting, all samples were given a homo-genization treatment which consisted of holding them for 15 hr at 1150°C (2100°F) and water quenching. Testing was begun with the samples in this condition. The tests employed in the study of this system were (1) dilatometer, (2) hardness versus aging temperature, (3) tensile strength and elongation versus temperature, (4) microstructure analysis, and (5) electrical resistance versus temperature. Dilatometer tests were made at the rather rapid heating rate of 1093°C (2000°F) in one hour on a Bristol-Rockwell type instrument. Hardness data were taken on specimens cooled from 204", 427", 649°, 760°, 871°, 982°, and 1093°C under conditions which are thought sufficient to bring the specimens to equilibrium in all but the cases of the very sluggish transformations. Electrical resistance data for several specimens were taken for both heating and cooling conditions in a vacuum furnace especially designed for this work by D. W. Bainbridge, formerly of this laboratory. Provisions were made for heating and cooling standard specimens at a constant rate, while autographic records of resistance and temperature were made simultaneously. Micrographs of all the alloys were made in the quenched condition. Some question may exist as to how such physical values as hardness, tensile strength, and elongation were interpreted to indicate a change in phase. Original data were recorded on physical property versus temperature graphs, each of which was made from the data of a single alloy (fig. la). From these, another series of graphs were plotted with the physical property as a function of composition for constant temperatures (fig. 1b). Sharp deviations in the slope of these composition versus hardness curves often indicate a change in phase. The example of fig. 1 will serve to describe this experimental technique. The plotting of a single point from graph "a" to graph "b," and finally to the phase diagram "c," is illustrated. The location of the point on graph "b" is indicated as point (1). If all the alloys of the A-B binary system were solid solutions at temperature X, the hardness in this curve would be expected to rise at a fairly even rate as A is diluted with B to a maximum at some intermediate value between A and B, as

Jan 1, 1951

By J. W. Pugh, J. D. Nisbet

F. B. Foley—The use of data published by Wever and Jellinghaus in 1931 to fix boundaries of the sigma phase in the Fe-Cr system, in the face of the author's own references to the suggestions of Bradley and Goldschmidt, Aborn and Bain, and Hougardy that the phase is much more extensive, and the very much more accurate work, which weighs heavily in favor of these suggestions, of Cook and Jones, published in 1943 and evidently disregarded by the authors, makes their derived ternary diagram, especially fig. 15 for 400°C, quite inaccurate on the Fe-Cr side and affects the extent of phase boundaries in the most controversial and important part of this ternary system. In commercial Fe-Ni-Cr alloys the occurrence of sigma has been observed time and time again at lower chromium contents than that of the 24 pct Cr, 16 pct Ni, 60 pct Fe alloy which is the lowest permitted by fig. 15 of practically carbonless metal. Newel1 reports sigma in 27 pct Cr-Fe even with some carbon present to be one of the greatest detriments to its extensive application, whereas Pugh and Nisbet set a low limit of 36 pct Cr for sigma in the binary Fe-Cr system. J. J. Heger—The diagrams presented by the authors do not agree with observations made on commercial Fe-Cr and Fe-Cr-Ni alloys, nor do they agree with two recent investigations made on the Fe-Cr and the Fe-Cr-Fe systems. I refer first to the investigation made by Cook and Jones1' on the sigma region of the Fe-Cr system. On the basis of their results which were published in 1943, Cook and Jones established new boundary limits for the sigma and the alpha plus sigma regions. These new limits have been accepted and are incorporated in the Fe-Cr diagram that appears in the 1948 edition of the Metals Handbook. This diagram is shown here as fig. 34. As will be noted, the boundary limits of the alpha plus sigma region in this diagram extend to much lower chromium contents than do those in the diagram presented by the authors in fig. 2. These new limits indicate that sigma phase should form in an Fe-Cr alloy containing 26 pct Cr, and experience with commercial alloys confirms this finding. Indeed, recent studies on a commercial Fe-Cr alloy containing 17 pct Cr have shown sigma phase to be stable in this alloy at 550°C. I recognize that the authors' work did not include studies on the sigma region; however, I believe this is a serious omission because sigma phase may profoundly affect the physical properties of these alloys and should be evaluated in any investigation which attempts to relate physical properties to the equilibrium diagram. The second investigation to which I refer is that made on the Fe-Cr-Ni system by Rees, Burns, and Cook,'' who used pure alloys and employed heating 17 W. P. Rees, B. D. Burns, and A. 3. Cook: Journal Iron and Steel Institute. (July 1949) 162, Part 3. p. 325. times that extend to 200 days. These investigators published their results in July 1949 and from these results constructed isothermal sections at 800" and 650°C. The isothermal section at 650°C is shown here as fig. 35. This diagram indicates sigma phase should form in a pure 18 pct Cr-8 pct Ni alloy. Although sigma phase has not been observed after 10,000 hr at 1200°F in commercial 18-8 alloys containing carbon and nitrogen, it has been observed under the same conditions in 18-8 alloys modified with titanium and columbium, both of which serve to reduce the effect of carbon and nitrogen. This section and the one at 800°C suggest that the constant iron sections which are presented by the authors, should be considerably altered, if they are to represent an accurate picture of the system. A few of the suggested alterations are as follows: In fig. 9, the section at 50 pct Fe, the gamma plus sigma region should be widened, not narrowed at 800" and 650°C. In fig. 10, the section at 60 pct Fe, the gamma plus sigma region should be widened at 800° and 650°C. In fig. 11, the section at 70 pct Fe, an alpha plus sigma, an alpha plus gamma plus sigma, and a gamma plus sigma region should be added. In fig. 12, the section at 80 pct Fe, the alpha plus gamma region should be widened at 800°C. In fig. 13, the section at 90 pct Fe, the alpha plus gamma region should be widened at 650° and 800°C. Undoubtedly, the chief reason for the discrepancies between the authors' data and those of Rees, Burns, and Cook is that the authors did not employ long enough heating times to allow for transformation. In this connection, I wish to warn that transformations in these alloys, particularly transformations at temperatures below 800°C, are extremely sluggish and may require a year or more to approach completion. Therefore, unless extremely long heating times are employed or steps are taken to accelerate these transformations by such means as mechanical working, the results will not yield an accurate picture of the alloy system. Certainly, an accurate picture is needed for the development of better high-temperature materials. E. J. Dulis-The purpose of this work is not clear. If an improvement of the ternary equilibrium diagram of the Fe-Cr-Ni system was wanted, it seems logical that testing techniques superior to those previously used would be a prime requirement. To approach equilibrium in this system, long holding times are needed; a fact established long ago but apparently ignored by the present authors, who used the continuous heating and cooling tests in equilibrium studies. A publication on the same system by Bradley and Goldschmidt6 was criticized by Monypenny in a

Jan 1, 1951

By D. J. Mack, R. E. Reiswig

Six examples of the peritectoid transformation were selected from the literature and studied by the method of isothermal transformation. The kinetics and mechanisms of five of the examples are presented as TTT diagrams and photomicrographs. The exist-enc- of a peritectoid in the sixth case is doubtful. ALTHOUGH the peritectoid transformation per se has been known for many years, no precise published data exist concerning the kinetics or mechanisms involved in transformations of this type, except for the brief treatment by Rhines, et al. 1,10 Bearing in mind the fact that investigations of recent years are uncovering more and more peritectoids and suspected peritectoids, a thorough study of the well-established peritectoids appeared to be in order. It was for this reason that a study of the kinetics and morphological mechanisms of six binary peritectoids was undertaken. The six peritectoids selected from the literature for study were those reported at 7.02 wt pct Al-Ag, 26.0 wt pct Sb-Cu, 30.5 wt pct Sb-Cu, 32.3 wt pct Sn-Cu, 8.35 wt pct Si-Cu, and 21.2 wt pct Al-Cu. These selections were based on availability and purity of components, ease of preparation and heat-treatment, and estimated reliability of the available equilibrium diagrams in the regions of interest. EXPERIMENTAL PROCEDURE The alloys used in this investigation were induction melted in electrode-grade graphite and chill-cast in cast-iron split molds. In all cases, the alloys were so brittle that they could easily be broken into samples weighing 1 or 2 g. Chemical analyses showed that the alloys used were close to the respective peritectoid compositions reported in the literature and that the impurity levels were low in all cases. Metallographic examination showed uniform distributions of phases in all samples, indicating uniformity of composition in the samples studied. Isothermal transformation studies were carried out in fused-salt media, using the familiar inter-rupted-quench method. Uniformity of temperature in the salt baths was maintained by continuous stirring with a stainless-steel agitator. On the basis of actual observations of the temperature fluctuations, the estimated temperature control was + 10C for the Ag-Al and Cu-Sb alloys and ±30C for the Cu-Sn, Cu-Si, and Cu-Al alloys. The accuracy of all temperature measurements was estimated to be ±1°C. It was found necessary to mount metallographic specimens of the Ag-Al alloy in cold-curing methyl methacrylate, since the temperatures encountered in mounting in bakelite or lucite caused an appreciable degree of transformation to the ß phase. For the other alloys, wood-flour-filled bakelite mounts were used to avoid extraneous X-ray diffraction lines during the later examination of the metallographic specimens on a Norelco Geiger-counter d if f r ac tomete r. In the X-ray diffraction procedure, agreement between the published diffraction patterns and those obtained in this study was good. This was particularly important for phase identification, since the literature contained little in the way of micrograph description in some cases. Etching of the silver-aluminum alloy for metallographic examination was done by swabbing with either of the following reagents: 1) 10 g CrO3, 1 g (NH4), SO2, 0.5 g NH4NO3, 100 ml H2O, or 2) 10 ml NH4OH, 1 ml 20 pct KOH, 4 ml 3 pct H2O2, 5 ml H2O. The other alloys were etched with the usual bichromate etchant: 2 g K2Cr2O7, 1.5 g NaC1, 8 ml conc. H2SO4, 100 ml H20 (swabbed vigorously). EXPERIMENTAL RESULTS A) The Ag-Al peritectoid at 7.02 wt pct Al— The phase equilibrium involved in this peritectoid is shown in Fig. I.2 The phase boundaries in the vicinity of the peritectoid were most comprehensively established by Hume-Rothery, et al,3 who placed the equilibrium temperature at 448 °C and the equilibrium compositions of the a ß' and y phases at 6.11, 7.02, and 7.24 wt pct Al, respective The alloy used in this study analyzed4 6.95 wt pct A making it slightly hypoperitectoid according to the accepted equilibrium diagram. The rate of the transformation a+ y — ß' varies rapidly with degree of undercooling below the equilibrium temperature, passing through a maximum in the vicinity of 350°C. Thus the TTT dia-

Jan 1, 1960

By E. F. Adkins, R. W. Douglas

The variation of the spectral emissivity of osmium and ruthenium with temperature can be expressed by the following relations: DURING a study of the sintering characteristics of osmium and ruthenium, discrepancies in the previously accepted melting points of the metals were noted. The melting temperatures had not been measured at the time of the work, but were estimated in the literature as 2700°C for osmium and 2450°C for ruthenium.' A recent investigation by Baird2 has shown the melting temperatures of osmium and ruthenium are 3000" and 2250° C, respectively. In the present work, osmium bars sintered at temperatures as high as 2900° C showed no evidence of melting, whereas ruthenium showed definite evidence of melting at temperatures as low as 2250°C. In view of the apparent discrepancies, the melting temperatures of both metals were determined. The spectral emissivities of both metals were determined as an aid to studies of the sintering characteristics of these metals. MATERIALS AND PROCEDURES The materials used in these studies were metal powders supplied by Englehard Industries, Inc., through the courtesy of The International Nickel Co. Analytical data from three laboratories, although not in complete agreement as to the purity of the material or the major impurity elements, indicated that the maximum impurity levels were 0.5 and 1.0 wt pct in osmium and ruthenium, respectively. The metal powders were pressed to 1/4 by 1/4 by 6-in. bars at 20Tsi using carbon tetrachloride as a lubricant. The pressed bars were presintered 1 hr at 1200°C after heating 1 hr at 400°C. The presintered densities of the several bars of the metals ranged from 65 to 80 pct of ideal density. Specimens were self-resistance heated in a vacuum sintering bell. Temperatures were measured with an L&N, disappearing filament optical pyrometer. True temperatures were read at the base of an 0.30-in.-diam, 0.150-in. deep hole drilled into the center of a presintered bar. Apparent temperatures were read from the surface of the bar adjacent to the hole. MELTING POINT DETERMINATION The melting temperature of the metals was determined as the temperature at which a slight additional increase in current through the specimen resulted in a reduced temperature at the base of the hole. The apparent reduction in temperature is a result of the liquid metal filling the hole, thus altering the black-body conditions. The melting points of osmium and ruthenium, based on three deter-

Jan 1, 1962

By Nicholas J. Grant, A. W. Mullendore, Paul Predecki

The Duwez technique of splat cooling in which a molten droplet of metal is accelerated and made to impact on a cold, highly conducting substrate was investigated.- An apparatus for producing "splat" was constructed employing an explosive powder charge to accelerate the molten metal. Transport of the molten metal just prior to impact with the substrate was studied by means of high-speed photography. The molten particles are small spherical droplets from about 1 to 50 µ diameter. The average cooling rates for aluminum, silver, and a gold alloy splatted on nickel substrates at room temperature were determined experimentally and were found to vary from 18 to 5 x 10' C per sec. The heat-transfer coeficients for pure aluminum and pure silver splats cooled on nickel substrates at room temperature were found to be 2.7 to 6.8 and 13.6 to 54.2 cal per sec sq cm 'C (2 to 5 x 104 and 1 to 4 x 18 Btu per hr sq ft OF), respectively. Solidification rates in pure aluminum and silver splats were calculated. "SPLAT cooling" is a term describing a technique for extremely rapid freezing and cooling of molten metals and alloys to room temperature or below. The technique was developed by Duwez et al.' in 1960, and after refinement2 consisted in transferring a few tenths of a gram of molten metal to near sonic velocity to strike a suitably placed cold copper surface. Upon impact, the metal spread into a thin nonuniform film called a splat, about 10-4 cm thick. The splat particles produced in this manner were thin enough in some areas to be suitable for transmission electron microscopy, without further treatment, and, together with X-ray data, yielded a variety of structures which would be classified as follows: 1) supersaturated solid solutions (increase in solubility limit), 2) metastable crystalline stoichiometric and non-stoichiometric intermediate phases ( in Au-20.5 at. pct Si and in Au-14 at. pct Sb, respectively), 3) amorphous alloys, 4) retained high-temperature phases, 5) alloys having equilibrium phases present, but with unusual, markedly altered structures. The crystalline phases present in splats were usually extremely fine-grained and had low dislocation densities. Present interest in splat cooling centers around a study of the unusual structures produced by the technique, their contribution to alloy theory, and the possibility of the development of new or unusual properties. In addition, the technique is being examined for the production of bulk quantities of new high-strength alloys for low- and high-temperature use. Although a number of alloy systems have already been investigated by the technique, relatively little is known about the mechanism of splat formation and the physical conditions encountered by the metal during solidification and cooling. The purpose of this work was: 1) to determine the velocity, shape, and size of the molten metal droplets just prior to impact with the cold substrate; 2) to try to estimate or measure solidification rates, heat-transfer coefficients, and cooling rates in splats. EXPERIMENTAL TECHNIQUE Splatting Apparatus and Procedure. An apparatus similar to the type originated by Duwez et a1.,2 but employing an explosive charge to accelerate the molten metal, was used. A "Ramset" fastening tool (a gun normally used for driving studs into concrete, steel, and so forth) was mounted vertically above a resistance-heated graphite crucible as shown schematically in Fig. 1. A splat product was produced by melting a few tenths of a gram of metal at the position shown in Fig. 1, and then exploding a powder charge in the breach. The shock wave thus generated traveled down the barrel into the furnace, ejecting the molten metal through the 0.06-in.-diameter hole at the bottom of the crucible. The ejected metal was impacted on a high-conductivity metallic substrate where it formed a splat. High-Speed silhouette Photography. In order to investigate the size and shape of the ejected, molten metal just prior to impact with the cold substrate, high-speed silhouette photographs were taken, and are shown in Figs. 2 and 3. These photographs were obtained with the aid of an Edger-ton, Germeshausen, and Grier microflash unit and a submicrosecond flash drive equipped with a variable time delay. Two types of triggers were

Jan 1, 1965

By Elmars Ence, Harold Margolin

The partial isothermal section of the Ti-Cu-Zr system at 750°C has been studied. The crystal structure of Ti2Cu has been detevnlined as tetragonal and when expressed as face-centered tetragonal, a = 4.164, c = 3.611 A, and c/a = 0.867. TITANIUM systems in which ß-titanium decomposes eutectoidally into a and a compound rather readily are known as active eutectoid systems. Recently in active eutectoid alloys based on the Ti-Cu system, and containing additionally A1 and Zr, it has been found that zirconium may produce both moderate weakening and marked strengthening at elevated temperatures.' In order to obtain some insight into the behavior of zirconium, an attempt was made to determine whether compounds of the type Tix Zr and TixZryCu exist. Only results of the Ti-Zr-Cu studies are reported here since data on Ti-Zr alloys will be presented in a separate publication. EXPERIMENTAL PROCEDURE Alloys were nonconsumably arc melted in argon atmosphere as 15-g buttons, prepared from materials of the following purities: Bureau of Mines electrolytic Ti — Bhn 63 OFHC copper Reactor Grade Zirconium Bhn .C 150 Prior to alloy preparation the electrolytic titanium was premelted to avoid excessive spattering of the charges during melting. The weight losses during melting did not exceed 1 pct of the total weight of the charge. The alloys prepared are listed in Table I together with analyzed compositions. in cases where they were determined. Specimens for heat treatment were wrapped in molybdenum sheet, annealed in argon-filled quartz capsules, and quenched by breaking the capsules under iced brine. About 1 to 2 sec are required to remove the capsule from the furnace, drop it into the iced brine and break it. Because of the insulation effect of the quartz, this delay in quenching should not produce a large drop in temperature. In many instances, no difference or little difference has been detected between the same specimen quartz capsule quenched or directly quenched in vacuum by an impinging stream of water.' Ti-Cu-Zr alloys were homogenized for 48 hr at 825°C prior to annealing at 750°C. Annealing for 5 days at 750°C produced the same results as 11 days at temperature and it was therefore assumed that 5 days was sufficient to establish equilibrium. Titanium-copper alloys were homogenized for 40 hr at 900°C prior to heat treatment for 6 days at 750°C. Specimens for metallographic examination were electropolished and etched with the following solution:

Jan 1, 1962

By J. Gordon Parr, G. R. Purdy

The phase boundaries of "TiNi" have been accurately determined by high-tewperature X-ray diffraction and metallo-graphic techniques. The phase (centered at 51 pct Ni) has a restricted range of homogeneity at 500°C, and a much greater range at 800°C. "TiNi" in all alloys containing less than 51 pct Ni decomposes rapidly and reversibly at 36°C to form a previously unreported phase. ThE region of the titanium nickel system corresponding to the phase "TiNi" has been the subject of several investigations.1,2,4-7 Laves and wallbaum1 reported "TiNi" as a body-centered cubic (a2) Phase, and Duwez and Taylor2 determined its para- meter to be 2.986. Duwez and Taylor also reported that "TiNi" decomposed into "Ti2Ni" and "TiNi3" on prolonged heating at 650" and 800°C, which led McQuillan3 to suggest a sluggish eutectoid reaction between "TiNis" and "Ti2Ni". Margolin, Ence, and Nielsen,' working with higher purity alloys, found no evidence of this decomposition. They showed the "TiNi" phase region extending from 49 to 54 pct*

Jan 1, 1962

By E. Eichen, G. S. Ansell. L. R. Sefton

A Technique for the Preparation of Thin Films of Two-Phase Alloys Suitable for Use in Transmission Electron Microscopy In order to obtain foil sufficiently thin to permit transmission electron microscopy of two-phase alloys, particularly alloys in which the second phase is present as a very finely dispersed second phase, a technique has been developed to permit the thinning of these materials without either the large degree of pitting present in the electrochemical or chemical thinning process or the slicing of the second phase particles accompanying the mi-crotoming process. The metal or alloy to be thinned is first ground to initial sheet thickness of approximately 0.002 in. by hand grinding on 4/0 emery paper. This starting sheet of material is then reduced in cross section by an air-abrasion technique using an S. S. White dental abrasive unit. This unit consists of a supply of compressed CO, which provides a stream of high velocity gas impinging on the metal surface. Incorporated in the machine is a hopper-type of device in which fine abrasive particles of alumina are introduced into this high velocity gas stream. The abrasive particles are irregular in shape, approximately 30 across. If any particles are introduced into the sample during thinning, they should be readily observable. Impingement of this gas abrasive mixture on the metal surface causes rapid removal of the metal in a rather uniform manner. The rates of metal removal may be adjusted by both regulating the gas pressure and amount of abrasive material introduced into the gas stream. By moving the gas stream abrasive mixture slowly about the specimen or sheet surface, a large area of the sheet may be uniformly reduced to a thickness which is then suitable for use in transmission electron microscopy. Attendant with this metal removal is a result in cold working of the metal surface due to the impingement of the abrasive particles. The resultant foils, therefore, if made from materials which are desirable to view in the annealed condition, have to be subsequently heat treated. If the alloy from which these foils have been made has been initially heavily worked or deformed during manufacture, then no subsequent heat treatment is, necessary as in the case of the SAP-type alloys. Following this initial abrasive thinning technique, the material is then given a final electropolish which is just sufficient to remove any oxide which has formed on the metal surface during the abrasion technique. Care has to be exercised in this technique to make sure that none of the attendant problems associated with the chemical or electrochemical removal process are encountered. An electrolyte solution which has been found suitable for aluminum-aluminum oxide SAP materials is one containing 78 ml of perchloric acid, 120 ml of distilled water, 700 ml of ethanol, and 100 ml of butylcellosolve. The current density used is 1 amp per sq cm and a polishing time of 3 sec on each side of the foil is used. Extreme care must be used in choosing the current density and time to ensure that just the oxide film on the specimen is removed. Times or current density which are greater than this will lead to perforation of the film due to the removal of the A1,0, particles or a complete loss of the film due to complete solution of the matrix. Once the oxide film is removed the specimens must be kept in alcohol until placed in the microscope to prevent any further oxidation. Fig. 1 shows -an electron micrograph at X28,000 of a thin foil of an aluminum-aluminum oxide SAP-type alloy prepared using this technique. The alloy consists of a matrix of commercial purity aluminum containing a very finely dispersed second phase of aluminum oxide flakes. In the electron micrograph, the dark patches are the aluminum oxide particles. The lighter portions of the micrograph are the aluminum matrix. Within the aluminum matrix, the dislocation structure of the alloy is clearly visible. This structure consists of intersecting twin boundaries and single dislocations. No abrasive particles were observed in the samples. In addition, thin films of commercial purity aluminum were prepared using this method and observed in the electron microscope. Here also, there was no evidence of abrasive particles introduced in the samples. The technique presented provides a satisfactory method for the production of thin films of aluminum-aluminum oxide SAP-type alloys for use in transmission electron microscopy where none of the usual thinning techniques appears to be satisfactory.

Jan 1, 1962

By M. Herman, G. J. London

A series of dilute Fe-Co alloy specimens were prepared from zone-melted iron. The iron used was given ten floating-zone passes in purified hydrogen. Alloying was accomplished by plating a gradient of metallic cobalt onto a 1/4-in.-diam bar of material from a pure plating bath. The deposit was radioac-tively tagged by a small addition of co80 to the plating bath. Alloy homogenization was accomplished by passing an additional floating-zone pass through the plated section of the bar. The final concentration gradient achieved was from residual to about 1000 ppm (0.1 pct) added Co. After homogenization, the bar was swaged to 0.057 in. wire and cut to tensile specimen blanks 1-1/16 in. long. Each specimen was counted in a well-type scintillation counter. The concentration of cobalt was computed by combining the normalized specimen counting rate with the known weight (measured to the nearest 0.01 mg) of cobalt added during plating. Fig. 1 shows the cobalt gradient obtained, exclusive of residual cobalt which was estimated by neutron-activation analysis to be approximately 50 ppm. Results of recrystallization for 1 hr at 750°, 700°, 650°, and 600°C with specimens containing 40, 170, 470, and 870 ppm added Co are shown in Figs. 2, 3, 4, and 5. Recrystallization was accomplished in a dynamic vacuum of less than 1 x 10 torr. The specimens were mounted and polished parallel to the swaging direction. The recrystallization re- sults show a radical coarsening of structure with increasing cobalt concentration. The coarsest structure obtained was at the 470-ppm level (at all recrystallization temperatures) with a slight grain refinement at the 870-ppm level shown. A more normal solid-solution effect is seen in Fig. 5 (60OoC, 1 hr) where the only specimen completely recrystallized is the one containing 47 ppm Co. Since the possibility existed that some atmosphere-contamination phenomenon might be responsible for the observed results the recrystallization series was repeated using sealed quartz capsules. These were filled with purified argon which was gettered during heat treatment on metallic titanium included in the capsules. The results obtained were essentially the same as shown. Another identical alloy was made using a 5-mm-diam bar of Westinghouse Puron. Specimen preparation and recrystallization treatment were also as described. The results obtained (see Table I) were somewhat erratic. Some coarsening in structure does appear at the highest cobalt concentrations. There is no apparent relationship between the coarse, columnar, directionally solidified structure of the as-melted bars and any of the recrystallized grain structures observed, all of which are much finer than the former structure. A recrystallization process can be most simply described as one of homogeneous nucleation of recrystallized grains in a cold-worked matrix followed by more or less uniform growth of the strain free grain into the cold-worked material.' An impurity element may influence one or both of these processes. Aust and utter' have shown that impurities strongly influence grain boundary mobility in lead. Moreover, definite recrystallization textures are produced by small concentrations of tin,3 while additions of silver or gold4 produce large boundary in-

Jan 1, 1964

By D. H. Polonis, M. B. Kasen

AN unusual thermal etching phenomenon has been observed on the surface of a Cu-Si alloy. The observations were made during studies of vacancy condensation pit formation which were reported in previous papers.1,2 Figs. 1 and 2 are photomicrographs of the surface of a Cu-4.8 pct Si alloy after heating in vacuo (3 x X mm Hg) to 525°C and subsequently holding at that temperature for 16 hr in an impure (air-contaminated) argon atmosphere. Under these conditions preferential removal of silicon from the matrix by oxidation results in a surface which is entirely a phase overlying an a plus y structure. The light areas of Fig. 1 are regions in which very little thermal etching has occurred, while the majority of the surface has experienced severe etching. Similar regions in Fig. 2 are seen as plateaus above contours produced by net transfer of material from the specimen surface. Twin boundaries are also visible in Fig. 2. The difference in matrix mi-crostructure between Figs. l and 2 reflects the par- ticular orientation of low-index planes with respect to the surface of observation. Thermal etching contours of this type have been studied in detail by Hondros and Moore Both of the photomicrographs show that one or more dark spots are located approximately at the center of every protected area; where two spots are in close proximity, the protected areas have merged. The microstructures indicate that the anomalous resistance to thermal etching is related to the presence of these spots. Williams and Hayfield have shown that local regions of unusually high oxidation resistance can be caused by a surface monolayer of atoms with a work function higher than that of the base material. The result is an inhibition of the rate of electron removal from the base material and a decrease in the oxidation rate. The source of the protected layer may be either surface diffusion from contaminants on the specimen surface or pipe diffusion of solute atoms followed by migration across the surface. In the latter case the driving force is a difference in chemical potential resulting in a Gibbs adsorption phenomenon. In the present case, chemical polishing of the specimen surface precluded contaminants of a strictly surface nature. Consequently, diffusion of impurity atoms from within the lattice appears to be the most tenable explanation of the observed phenomena. The authors propose that the structures of Figs. 1 and 2 reflect the diffusion of impurity atoms along extended dislocations within the individual a phase grains prior to surface oxidation. The central spots within the protected regions are seen as sites of dislocation emergence which have become locally decorated with oxide in the manner discussed in Ref. 1. Williams and Hayfield observed protected regions adjacent to the grain boundaries of iron-contaminated copper which they attributed to the segregation of iron to the grain boundaries coupled with high dif-fusivity along the complex dislocation array at such boundaries. No evidence of such grain boundary protection was observed in the Cu-Si alloy; this

Jan 1, 1964

By P. D. Southgate

VIBRATION, both ultrasonic and sonic, can affect the course of a variety of metallurgical processes. Reports of work on this subject have appeared at intervals over the last twenty years and a thorough review has been made by Hiedemann,' who gives a full list of references. It has been established that both refinement of structure and outgassing of the melt can be produced in most metals with sufficiently intense vibration. Many of the experiments in this field have been carried out using alloys of direct commercial interest in production- type casting equipment. Here a different approach is adopted, and simple binary alloys which exhibit different types of structure are used to facilitate examination of the effects produced under well-controlled conditions. Results from this kind of experiment should enable the basis of larger scale experiments with similar types of materials to be better substantiated. The alloys selected for this work are: 1) aluminum of 99.2 pct purity, the principal impurity being iron, near 0.4 pct; 2) A1-1.8 pct Fe, which is near eutectic composition; 3) A1-4 pct Fe—the second phase, of composition FeAl3, forms distinct acicular crystals; and 4) A1-12 pct Si—although this is of eutectic composition, the silicon normally appears in the form of large plates. Experimental Method There are three practicable methods by which acoustic vibration can be transferred to liquid metals: a) vibration of the whole of the container; b) electromagnetic induction; and c) insertion of a vibrating probe into the liquid. Method c is chosen here, and a frequency of 8 kc per sec is used. The end of the probe is electrically heated so that the metal near it will be the last to solidify. This ensures that vibration is transmitted directly to the liquid metal without passing through any solidifying and pasty intermediate region, which would attenuate acoustic waves severely. In addition, the heating prevents solid metal adhering to the probe, with consequent detuning. The transducer assembly is shown in Fig. 1. A stack of Permendur laminations 30 cm long acts as a half-wave vibrator and a winding along its length carries polarizing and energizing currents. It is joined by a half-wave coupling bar to a mild steel exponential horn, which tapers from 3.8 to 1.1 cm diam. A stainless steel stub 3 cm in length is welded on the free end. The assembly is supported by a plate located at the node of the coupling bar. Acoustic radiation losses are minimized by allowing the free end of the transducer to protrude above the surface of the cooling water. The exponential horn serves the function of matching the acoustic output impedance of the transducer to the radiation resistance of the probe into the melt. Electrical drive is provided by a 1 kw generator. Measurements of probe vibration amplitude during melt treatment are obtained from a gramophone pick-up with the needle attached to the wide end of the horn. This device is calibrated with a microscope before use. Liquid aluminum is very actively erosive and solution of materials with which the molten metal comes into contactcauses difficulties. Erosion of the probes in particular can be severe since it is greatly accelerated by the vibration. Tests of various acous-

Jan 1, 1958

By Louis E. Toth, Alan W. Searcy

A modification of Le Claire's microscopic model for self-diffusion is developed in a form suitable for prediction of activation energies for diffusion in disordered substitutional solutions as well as in pure metals Bonding is considered as a localized interaction, and the energy of bonding between atoms of different types is taken as the arithmetic mean of the energies in the pure elements. The activation energies for vacancy formation and migration in substitutional alloys are shown to depend on the empirical constants developed for self-diffusion when the equations are adjusted for the mole fractions of the two elements. The differences in activation energies between alloys and pure metals are calculated and compared with the experimental differences. MaNY different correlations have been proposed for use in predicting activation energies for self-diffusion or diffusion in dilute alloys. Usually theories for self-diffusion cannot be applied to alloys nor can theories for dilute alloys be satisfactorily adapted to self-diffusion. We present here a correlation scheme which can approximate activation energies for diffusion, vacancy formation, and atom migration in substitutional-alloy systems of any concentration. The activation energy for self-diffusion QD has been correlated to the melting temperature Tm,1, 2 the sublimation energy L,,l the heat of fusion Lf,3 crystal structure, and valence.4 Working with mechanical parameters Le claire5 proposed that the activation energy for vacancy migration Qm could be related to appropriate shear moduli. He also proposed that the activation energy for vacancy formation 62, is proportional to Ls. THEORY We consider that Le Claire's expression for 4, is satisfactory. But an alternate expression for Qm in terms of the bulk modulus is just as satisfactory from a theoretical point of view for the prediction of the energy of vacancy migration in pure metals and can more easily be used for alloys. In order for Qm to be a minimum, not only the moving atom but also adjacent and nearby atoms must undergo distortions. These distortions are neither isotropic as is implied by a correlation with the bulk modulus nor unidirectional as implied by a correlation with a shear modulus, but are complex and multidirectional. A complete analysis in terms of the bulk modulus and all the shear moduli is not yet available. In a first approximation in which only one parameter is used, there is no obvious theoretical reason to prefer the shear modulus over the bulk modulus. The bulk modulus is a much more convenient parameter for use in predicting Qm. Shear moduli are known for less than half of the metallic elements while bulk moduli are usually available. Furthermore, because of the directional dependence of shear moduli, their estimation for alloys appears inherently more difficult than is the estimation of bulk moduli. We assume that the bulk modulus Bs reflects an average value of the shear moduli such that the strain energy at the point of maximum distortion of the lattice during vacancy migration can be expressed as

Jan 1, 1964

By B. F. Decker, D. Harker

The recrystallization reaction in OFHC and spectroscopically pure copper has been followed by X ray diffraction determinations of the amount of material with the cold-worked and recrystallized textures in specimens which had been given various heat treatments. The heats of activation for recrystallization are found to be: 29.9 kcal per rnol for OFHC copper and 22.4 kcal per rnol for spectroscopic copper. A reliable and speedy method for measuring the amount of recrystallized material in a piece of rolled metal, together with a new scheme for low temperature heat treatment, have made possible a determination of the activation energy for recrystallization in rolled copper. This method for studying recrystallization rates is different from others reported in the literature.' Oxygen free, high conductivity copper—OFHC copper—of purity 99.98 pct gave an activation energy for recrystallization of 29.9 kcal per mol with recrystallization data taken in the temperature range 208-245°C. Copper of purity 99.999 pct from the American Smelting & Refining Co. gave an activation energy for recrystallization of 22.4 kcal per mol with recrystallization data taken in the temperature range 43-135°C. The great change in activation energy for recrystallization as the small amount of impurity in the metal is decreased suggests that the motion of grain boundaries is conditioned by the impurities which must be concentrated in them. Thus the OFHC copper recrystallizes with an activation energy of the order of magnitude to be expected for the diffusion of the kind of impurities likely to be present. The high purity copper, on the other hand, recrystallizes with a much lower activation energy, in accord with the notion that the grain boundaries need not wait for the impurities to diffuse with them. Copper when strongly cold-rolled has a well-defined crystallite texture which can be described by saying that the [111] direction of the face- centered cubic crystals is aligned approximately in the cross-rolling direction and that faces of the form {110} lie approximately in the rolling plane. When such rolled copper is completely annealed the crystallite texture is quite different: the [loo] direction is in the cross-rolling direction and faces of the form (100) are in the rolling plane. At intermediate stages of annealing, the copper contains both textures. If a Geiger counter X ray diffraction spectrometer is set up so as to measure the intensity of the (200) Bragg reflection from the rolled surface of a piece of copper strip, this intensity increases from a very small value in the cold-worked state to a maximum as the specimen is annealed. The same would be true for (200) reflections from planes normal to the rolling or cross-rolling directions. In view of this fact, the amount of material in the copper which has recrystallized can be measured by comparing the intensities of one of these reflections just mentioned with the intensity from a similar specimen after complete annealing. In much the same way the intensity of (111) reflections from planes normal to the cross-rolling direction could be used to measure the amount of unrecrystallized material in the specimen. In the work to be described here the amounts of recrystallized and unrecrystallized material in partly annealed rolled copper specimens were determined in this way. Experimental Procedure Samples of oxygen free high conductivity (OFHC) copper, cold-rolled 99.7 pct to 0.002 in. thickness, were subjected to heat treatments for various chosen times at four different temperatures, and high purity copper samples, cold-rolled 98.0 pct to 0.0015 in. thickness, at six different temperatures. Each sample, after treatment, was placed in an orientation sample holder for use with a Geiger counter X ray spectrometer and a reading was taken of the intensity of the (200) reflection from planes perpendicular to

Jan 1, 1951