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By D. C. Buffum, L. D. Jaffee

THAT the path of brittle fracture in temper brit-tle steel differs from the path in unembrittled steel has been known for some years1 More recently it has been recognized that alloy steels, at least, display two forms of temper brittleness1 which may be called upper-nose and lower-nose. No experimental evidence has been published, however, to show that the two forms can be distinguished by simple inspection of brittle fractures. Fig. 1 is a photograph of the fractured surface of a V-notch Charpy specimen of SAE 3140 steel in which temper brittleness developed during heat treatment in the lower-nose temperature range.' The steel was austenitized 1 hr at 900°, tempered 1 hr at 675", and embrittled 48 hr at 500°C, with water-quenching after each treatment. The temperature of transition from ductile to brittle fracture in the Charpy test was + 2°C (criterion of transition: 100 pet fibrous). The hardness was Re 25; the austenite grain size ASTM 8. This specimen was broken at 150°C in the standard Charpy test and absorbed 6.6 ft-lb; its fracture was rated 0 pet fibrous. Fig. 1 shows that the fracture appeared rough to the naked eye. At the side of the specimen away from the notch (the side which was in compression during the test), the surface as a whole deviated sharply from a plane. At higher magnification, X60, the surface gave

Jan 1, 1959

By G. W. P. Rengstorff, H. B. Goodwin

A method is described for purifying iron in batches of 150 Ib or more. Oxygen, carbon, nitrogen, and sulphur are removed from flakes of electrolytic iron by treatment in wet and then dry hydrogen. A special consumable-electrode arc furnace is used to remove hydrogen and to melt the flakes into ingots. A LTHOUGH iron and steel have been made and used for hundreds of years, there are still wide gaps in the knowledge of these metals. One reason for this is that the true properties of the base metal, iron, are not well known, simply because really pure iron has never been readily available for research purposes. Commercial iron, and even the "pure" irons which are readily available, are really complex alloys of iron with oxygen, nitrogen, hydrogen, carbon, sulphur, silicon, and other elements. With such materials, it has not been possible to prepare steels and other iron-base alloys in which all elements are present in exactly controlled amounts. It is an almost impossible task to determine the separate effects of each element or the intereffects of several elements unless alloys can be prepared containing only the desired elements. The ability to make controlled alloys of low impurity content is of far more than academic interest. It will assist in the study of such practical problems as the brittle failure of ship hulls or high pressure gas pipelines, determination of the factors which affect the deep-drawing properties of steels, the improvement of magnetic and electrical properties, the phenomenon of aging, and many others. In 1949, the American Iron and Steel Institute asked Battelle Memorial Institute to prepare high purity iron on a larger scale than had been attempted up to that time. The work was directed toward the preparation of iron as low as possible in carbon, oxygen, nitrogen, hydrogen, and sulphur because these are the contaminating elements of major interest in many research investigations on iron-base alloys. The initial goal was to prepare iron with not more than a few thousandths of one percent of each of these impurities. The purpose was to set up a "bank" of high purity iron on which researchers could draw when a supply was needed and which could be replenished from time to time as the stock became depleted. Of several possible purification methods, heating in hydrogen was selected as the most promising for the removal of both carbon and oxygen from solid iron. Metallic impurities are not removed by hydrogen. Electrolytic iron was chosen as a starting material because it is the most readily available form of iron low in metallic impurities. Decarburization and desulphurization were obtained by annealing in wet hydrogen at relatively low temperatures. The optimum temperature, as found by previous investigators,' is about 1400°F. This treatment also removes nitrogen. Deoxidation is favored by dry hydrogen and higher temperature. There is evidence from previous work that an anneal at 2200°F in dry hydrogen will deoxidize the iron, reduce the nitrogen content, and reduce the sulphur content. These facts suggested the use of a two-stage process: treatment in wet hydrogen followed by treatment in dry hydrogen. To carry out the purification process in a reasonable time, it was necessary that the iron be in thin pieces. Flakes of electrolytic iron were used. Because most research requires that the iron be in massive form, it was nccessary to melt the purified iron to consolidate it into large pieces. For this work, it was decided to prepare ingots weighing at least 15 lb by using a cold hearth arc-melting process in which the iron was melted in an inert atmosphere of argon or helium. This method of melting was selected because the water-cooled copper crucible is not wet by the iron and does not contaminate it as a refractory crucible would. Because the crucible is also the ingot mold, contamination from mold materials was also eliminated. Contamination from electrodes was avoided by making them from purified electrolytic iron. The melting procedure finally developed was an adaptation of the Kellogg electric ingot process.' Two batches of purified iron weighing about 100 Ib each in ingot form were prepared with analyses as shown in Table I. The remainder of this paper describes the procedures in more detail. In general, the process con-

Jan 1, 1956

By Nicholas J. Grant, Noboru Komatsu

Metallographic and X-ray studies were made of oxide dispersion strengthened Cu-12 vol pet SiO2 and Cu-3.5 vol pet Al2O3 alloys following time exposures at temperatures approaching the melting. point of the matrix metal. Changes in hardness, oxide particle size and crystal structure were studied to determine the time-temperature stability of these alloys. An interpretation of the mechanism of oxide agglomeration is proposed. One of the most striking features of dispersion strengthened metal-metal oxide alloys is their stability at elevated temperatures, both with respect to recrystallization and maintenance of mechanical properties after exposure at high temperatures. A small amount of work has been done on the re-crystallization phenomenon in dispersion hardened metal-metal oxide alloys'l and several theoretical treatments have been advanced regarding the stability of these alloys by Lenel and Ansell,8 and by Grant and his co-workers.9, 10 This paper reports on metallographic and X-ray studies following various heat treatments of Cu-AI2D3 and Cu-SiO2 alloys at elevated temperatures and the effect of subsequent cold work on the recrystallizalion temperature. An interpretation of the recrystallization behavior of this type of alloy, based on the decomposition and agglomeration of oxides, is proposed. EXPERIMENTAL PROCEDURE Cast ingot bars of copper-1.59 pet Si and Cu-0.77 pet Al, 7/8 in in diam, were annealed in argon for 80 hr at 1000CC for homogenization. The homogenized rods were machined into fine chips and ball milled. Milled particles in the range minus 20 to plus 60 mesh were separated for internal oxidation. The internal oxidation was accomplished by the same method used by Preston and Grant,9 except that the oxidalion temperature for the Cu-Si alloy was 750°C, and for the Cu-A1 alloy was 800°C. These powders were then hydrogen-reduced at 450°C, compacted in a copper container, evacuated at room temperature to 0.04 µ, and sealed for extrusion. They were extruded at 1400°F (760°C) with a ram speed of 55 in. per min for an extrusion ratio of 15.6 to 1 to give a rod 0.38 in. diam by about 3 ft long.

Jan 1, 1962

By R. B. Russell

The expansion coefficients of hexagonal (a) zirconium have been calculated from the lattice expansionparameters of both coefficientsofhexagonallow and high hafnium alloys in the range 0° to 600°C. It is found parametersofboththat the coefficients are straight-line functions of the temperature and that the effect of hafnium is to depress the mean coefficients by about 1 pct at room thattemperature and 9 pct at 600°C. The cubic (ß) form was stabilized so that its coefficient could be determined from 400' to 900°C. COEFFICIENTS of thermal expansion for hexagonal-close-packed (a) zirconium were determined up to 600°C (1112°F), without any interference from the effects of preferred orientation, by measurement of the lattice parameters obtained from a symmetrical back-reflection focusing X-ray camera. This paper combines the results for a 1.2 atomic pct hafnium alloy from an earlier report' with those for hafnium-free zirconium. The problem confronting the investigator is essentially that of getting sharp, high-angle diffraction lines which represent the material at a known constanl. temperature and chemical composition. In order to do this, it was necessary to use the proper X-ray emission wave length (Co K was chosen) and to have a fine-grained metal whose atomic spacings would be free from effects of cold work (the specimen was annealed) and to prevent a changing composition owing to gas absorption during the heating of the metal to temperature (the sample was heated in a vacuum). A summary of the effect of temperature on properties is listed in Table I which gives data for ordinary zirconium, as well as new values for the low hafnium metal. The coefficients of the linear thermal expansion that have been reported are given in Table 11. The discrepancies shown may be explained perhaps by the effect of preferred orientation, and McGeary and Lustman2 have done this successfully for sheet metal. The differences in texture cannot explain all the differences in expansion coefficient shown in Table II. liesults by the author agree most closely with those of Kroll.3 Chemical analysis of the high hafnium metal pre- viously reported1 showed that it was 98.22 atomic pct (97.40 wt pct) pure. The low hafnium metal was dehydrogenated Westinghouse crystal bar, designated SMZ-137A. Its analysis was as given in Table 111 (chemical analyses 5377, 5431, and 5432).

Jan 1, 1955

By K. R. Garr, D. Kramer, A. G. Pard

EXPOSURE of steels to a fast-neutron flux results in helium generation by (n,a) reactions. The elements Fe, Ni, Cr, and N are major participants in the (n,a) reaction and helium concentrations of up to 5 x 10"5 atom fraction could accrue in fuel clads after a year of service in proposed fast reactors.' Such concentrations of helium can severely reduce the ductility of alloys at temperatures of reactor service. The phenomenon of helium embrittlement has been studied in austenitic alloys which were irradiated with particles as a rapid means of producing significant helium concentrations.2-5 The pres-ent work is a continuation of those studies and is concerned with helium effects in a ferritic stainless steel. I) EXPERIMENT Small sheet tensile samples with a gage length of 0.50 in. and a cross-section of 0.009 by 0.040 in. in the gage length were punched from cold-rolled type 405 stainless steel. The composition of the steel used in this work is given in Table I. Prior to helium injection, the samples were annealed 24 hr at 760°C which caused complete recrys-tallization with a resulting grain size of 40 u. This treatment rendered the alloy completely ferritic and carbide particles, 0.5 to 2 uin size, were present at a density of 1 x 1011 cm-3, Fig. 1. These carbide particles were probably a mixture of the M7C3 and M4C types.6 Helium was injected into the samples by irradiation with a particles from a cyclotron, which produced uniform concentrations through the center two-thirds of the sample thickness. Details of this procedure were previously described.' The average helium concentration, as measured on a gas-source mass spec- 11) RESULTS A) Tensile Tests. The results of the tensile tests are given in Table I1 where the values shown are averages of two samples. There was no significant effect of helium on either the yield strength or tensile strength at 550" or 650°C but an increase in both is apparent at 750°C. Samples with helium showed a 20 pct loss in total elongation at 650°C which increased to a 55 pct loss at 750°C. The uniform elongation is not affected by helium at 550" or 650°C but shows an increase at 750°C.

Jan 1, 1970

By N. S. Stoloff, R. G. Davis

It is concluded from an X-ray stztdy that the formation of the hcp Mg3Cd superlattice is a nuclea-tion and growth reaction. A two-phase, ordered-plus-disordered, region is observed between 153" nnd 141"C. A maximum in the flow stress is observed both at temperature and in quenched specimens within the two-phase region. The peak arises fro112 the extra energy required to create antiphase boundary upon cutting the ordered regions. Hardening during isothermal annealing in Mg3Cd has been shown to be similar to the domain hardening in Cu3Au-type alloys in that up to the maximum the structure is ordered domains growing into a short range ordered matrix. The hardening is due to a combination of short range order strengthening and the strengthening produced when antiphase boundary has to be created. In a recent study1 of the effect of long-range order upon the mechanical properties and deformation modes of hcp Mg3Cd a comparison was made be- tween quenched (disordered) and slow-cooled (fully ordered) specimens. The most striking observations were that the disordered alloy yields at a much higher stress than the ordered alloy, and that there is a change in slip mode from primarily (_0001) ba_sal slip in the disordered state to (0001), {1010}, {1122}, and {l011} slip in the ordered state. The main purpose of the present work is to extend these observations to determine whether MgsCd exhibits a peak in flow stress at an intermediate degree of order, or at a critical domain size, as do other orderedstems. However, disagreement exists about the Mg3Cd order transformation with respect to: 1) the location of the critical temperature for order, 165°C or 153"~,' and 2) how the degree of order varies with temperature. Welber et a1.5 concluded from specific heat measurements that the formation of Mg3Cd is a second-order reaction with the degree of order a continuous function of temperature, while Moore and Raynor concluded that it is neither a first-order reaction, with the degree of order a discontinuous function of temperature, nor a second-order reaction. Consequently it was considered necessary to undertake an X-ray investigation into the change in degree of order and structure in the

Jan 1, 1964

By R. N. Stevens

GRAIN boundary sliding plays a part in the high-temperature deformation of polycrystalline solids, and many investigations1-11 have been carried out to determine the deformation contributed by this mechanism. The principal method employed has been to measure one of the components of the grain boundary displacement, see Fig. 1, at a large number of boundaries, and from the mean value so obtained to calculate the elongation contributed by sliding. A formula connecting these mean values and the elongation contributed is evidently required, and many have been derived. They are usually of the form where Egh is the tensile elongation contributed by sliding, is the mean value of the component, x, N is the mean number of grain boundaries intersecting a line of unit length, and K is a dimensionless parameter known as the averaging factor. Table I summarizes the equations used by various investigators. Examination of these shows that their derivations usually involve doubtful assumptions, and are often in serious error. The purpose of this note is to point out that rigorous solutions to the problem of calculating Egb from x have been available in the literature for some years.' Since this is so little known, and also the original derivations do not make some important points clear, it is proposed to rederive the basic equations. These will be extended further, and the conditions in which the extended equations may be applied will be indicated. The mean values of the components of the sliding vector (I, d, and I:, Fig. 1) will depend on the method by which boundaries are sampled in order to make measurements.3,11 Boundaries are usually sampled by marker lines parallel to the stress axis, or lines perpendicular to the stress axis, the measurement being made at the point where the marker line intercepts the boundary. We will denote the mean values determined by these various sampling methods as follows. lT, dT and vT are mean values obtained by making measurements at pointswhere a transverse line intercepts the boundary while lL, dL, and vl, are the mean values from a longitudinal line. The number of boundaries per unit length also may depend on the direction of measurement, and so we define NT and Nl to be the number of boundaries per unit length, after creep, in the transverse and longitudinal directions, respectively.

Jan 1, 1967

By R. E. Grace

WHEN an Fe-C alloy, partially reacted to pearl-ite, is quenched rapidly enough to suddenly stop the growth process, it may be expected that any carbon concentration gradients will freeze in situ in the newly formed martensite. A short tempering treatment should precipitate finely dispersed iron carbide particles in approximately the same distribution in which carbon existed during the growth process. Examination with the electron microscope of areas of tempered martensite adjacent to the pearlite-martensite interface should reveal the existence of preferred areas of precipitation. Fig. 1 is an example of an electron micrograph used in this study. Small sections of a commercial steel (0.9 pct C, 1.6 pct Mn, 0.25 pct Si) were austenitized at 945°C in air for 45 min, furnace-cooled to 697°C, iso-thermally reacted to pearlite for 30 min, and quenched into iced brine. Two samples were tempered 5 and 20 min, respectively, in air at 498°C and were subsequently quenched in water. The samples were mechanically polished and etched in 1 pct picral. Collodion replicas were made by the conventional technique; these were shadowed with chromium at about 650' from the horizontal, and electron micrographs were taken at a linear magnification of 14,200X. Prints were made at a magnification of 44,100X. Lineal analysis of the prints was made as a function of the interlamellar spacing and volume percent of iron carbide. Straight lines were drawn into tempered martensite between adjacent lamellae at a distance S/4 or S/6 apart, where S is defined as the interlamellar spacing measured at the pearlite-martensite interface. The volume percent of iron carbide was measured along each of these traces with a "gage length" of S/10. The results of analyses of two samples are shown in Table I. The area included between the S/2 and S distance from the interface could be considered homogeneous within the sensitivity of the measuring technique. This value was used as a proportionality factor to evaluate an arbitrary carbon content of homogeneous austenite, and also as a parameter by which error in measurements could be evaluated. Results of the lineal analysis are shown in Figs. 2 and 3. Fig. 4 shows the statistical nature of the interlamellar spacing for all lamellae measured in this

Jan 1, 1954

By N. A. Gokcen, S. Fujishiro

THE pressures of Mn(g) in equilibrium with Mn7C3 and graphite have been measured by McCabe and Hudson' and Butler, McCabe, and paxton2 by means of graphite, zirconia, and Ta-Mo Knudsen cells. The details of corrections in their measurements were somewhat different from those of the authors;3'4 hence, further independent results at higher temperatures seemed to be warranted. The stoichiom-etry of Mn7C3 in equilibrium with graphite and its stability up to the vicinity of 1373 °K have been satisfactorily established by Kuo and persson5 (See also Hansen and Anderko).' The method of investigation has been described elsewhere in detail.3'4 The carbide was prepared in the manner described by McCabe and Hudson. In order to avoid possible reactions with refractory oxides, only graphite cells were used for the effusion experiments. The authors' results are the runs M-1, M-2 and M-3 listed in Table I, together with those of McCabe and co-workers. The heat capacity, c;, of Mn7C3 is not known; therefore McCable et al. plotted their data as log P vs 1/T and the results were only expressed first' as, and later2 as logP=- -14220+6.06 without further thermodynamic calculations. However, it is possible to estimate closely the necessary heat capacity as follows: When H; - &98 of CrC, or MnO, is plotted vs x, corresponding to various compounds, with temperature as the parameter, it is seen that the relationship is represented by a line of very small curvature. This relationship is also followed by other compounds which do not exhibit solid state phase transformations. In view of the fact that chromium and manganese have the atomic numbers of 24 and 25 respectively, it is reasonable to assume that H; - H;ae vs x for MnC, follows the same type of lines as those for CrC, for which adequate data are available.7 Similar correlations appear to be valid for oxides of two neighboring metals for which reliable data are available. The procedure is thus to move the lines for CrC, vertically (parallel to the axis of HT - H298) until they coincide with the points for MnC1/3, the only carbide of manganese whose H; - HO298 has been

Jan 1, 1963

By M. B. Bever, A. M. Reti, A. K. Jena

The results of this investigation show that in the system SnTe-PbTe l) the solid solutions have small exothermic heats of formation relative to the binary compounds, 2) solid-solution hardening is almost absent, 3) the tin and lead atoms are .randomly distributed on the sublattice occupied by them and do not undergo an ordering transition during a thermal treatment designed to promote such a transition, and 4) the lattice parameter changes linearly with composition. These findings suggest that the solid solutions of tin telluride and lead telluride approach ideal behavior. THE binary compounds SnTe and PbTe form a continuous series of solid solutions.' Information on several physical properties of these solid-solution alloys has been published.1"3 In the work reported here, the heats of formation, microhardness, and lattice parameters of these alloys have been measured and the possibility of an order-disorder transition has been investigated. Order-disorder transitions occur in some solid solutions of binary compounds. In the system Bi2Te3-BizSes, for example, the electropositive bismuth atoms and the electronegative tellurium and selenium atoms always occupy separate layers. In the disordered alloy Bi2Te2Se, the tellurium and selenium atoms are randomly distributed on the layers occupied by them, while in the ordered alloy they occupy distinct layers.4 The NaC1-type (Bl) structure of SnTe-PbTe alloys consists of two sublattices, one of which is occupied by the electropositive tin and lead atoms and the other by the electronegative tellurium atoms; any ordering would be expected to occur in the sublattice occupied by the tin and lead atoms. PREPARATION OF SPECIMENS Polycrystalline specimens of PbTe and Sn1-xPbxTe (x = 0.25, 0.50, and 0.75) were prepared by melting stoichiometric amounts of lead (99.999 pct), tin (99.98 pct), and tellurium (99.999 pct) in evacuated and sealed Vycor tubes. Because of the deviation of tin telluride from the stoichiometric composition5 this compound was prepared by melting 49.8 at. pct Sn and 50.2 pct Te. The tubes containing the melts were quenched into water; the solidified ingots were removed from the tubes and annealed for 2 weeks in vacuum (PbTe approximately 100°C below its melting point, the three alloys 100°C below their solidus tem-

Jan 1, 1969

By H. J. Snyder

Both the room-temperature and 600°F (315°C) primary solid-solution phase boundaries for the silver comer of the Ag-Cd-In temary system have been determined using X-ray diffraction and metallo-graphic techniques. Solid solubility of indium in silver at room temperature and 600°F (315°C) is shown to be between 20.3 and 21.7 wt pct, and corrects the previously extrapolated value of 18 wt pct appearing in the literature. The solid solubility of cadmium in silver at room temperature, not previously established, is shown to be between 34.9 and 41.3 wt pct. The solid solubility of cadmium in silver at 600°F (315°C) is confirmed as 43.2 wt pct. The shape of the primary phase boundary line for the subject temary system appears straight at room temperature and curves con-vexly towards the solvent (silver) comer at 600°F (315°C). THIS investigation was undertaken to initiate development of the Ag-Cd-In ternary diagram by establishing the primary solid-solution phase boundary in the silver corner of the diagram. At the present time, a diagram for this ternary system is not available, while this system appears to offer potential for a variety of uses. One alloy in the system has already proven to be useful as a control rod material in nuclear reactors.1"3 To encourage additional exploration of the potential of the system an understanding of its ternary diagram would be beneficial. It is towards this end that the present study is directed. LITERATURE SURVEY In undertaking this study, a comprehensive literature survey was first conducted to find information that would be useful for guidance. Useful knowledge did result from investigating the respective solubilities of indium and cadmium in silver as indicated in the binary constitutional diagrams of the systems Ag-In and Ag-Cd and also reviewing possible shapes of ternary primary phase boundaries. Weibke and Eggers~ used metallographic techniques to establish the system In-Ag as illustrated in Fig. Is This diagram shows that silver in the solid state is capable of dissolving approximately 18 wt pct In. This solid solubility of indium in silver holds true from room temperature up to 1276°F (691°C), then decreases continuously up to 1'761°F (960.5"C), the melting point of the silver. The solid-solubility limit of indium in silver established by Weibke and Eggers has been confirmed by Hume-Rothery and Reyno1ds. G Durrant' used a method which consisted of microscopic examination of quenched specimens to establish the Cd-Ag system as displayed in Fig. 2.' The primary phase boundary of Fig. 2 proved to be in good agreement (ie., within 1 wt pct) with other studies*'1' of the Ag-Cd system made using X-rays and metallog-

Jan 1, 1968

By J. Szekely

A rrlethod is presented for the estitnation of oxygen pickup by teettzed molten steel streams. Of the mechanisrt~s considered for oxygen absorption, physical entrainment appeared to be the most significant. The estent of this physical entrainment may be calculated by the application of boundary layer theory. It was found that the values of oxygen pickup thus predicted were in wry good agreement with data reported in the literature, in the case of short, well-defined metal streams. This entrainment model does not provide an adequate description of oxygen pickup by long streams, exoosed to external winds and where breakup of the streatn is likely. WheN molten steel is teemed into a mold or tundish it will absorb oxygen contained in its environment. Recently interest has been expressed in the measurement and possible prediction of the extent of this oxygen absorption: this interest owes its existence to two principal factors: a) the desirability of providing a quantitative assessment of the role played by the absorbed oxygen in impairing the quality of (vacuum degassed. stainless, or continuously cast) steel; b) the need for providing guidelines for stream protection methods. The purpose of this communication is to suggest a simple method for the prediction of oxygen absorption by continuous metal streams and to compare the predicted values with the data available in the literature. It is hoped that this method may provide at least an initial framework for the interpretation of additional data that might become available. FORMULATION Consider a continuous metal stream falling into a molten metal pool as depicted in Fig. 1. This system may absorb oxygen due to the following two mechanisms: 1) physical entrainment of gases (including oxygen) into the molten pool where reactions. thus absorption, will occur; 2) mass transfer (convec:tive diffusion) of oxygen to the metal stream. 1) Physical Entrainment. The physical entrainment may be calculated by using the concept of the diaplace-ttzent tlzickness of the laminar boundary layer, which will be stated briefly in the following.&apos; When a fluid flows past a solid surface, a portion of the fluid close to the surface is retarded due to fric-tional forces. This region in which the velocity of the fluid differs from that in the main stream is called the boundary layer. An illustration of this is given in Fig. 2 where it is seen that the thickness of the boundary layer increases as we move along in the x direction. In the laminar region, when Re, < 10&apos; the thickness of the boundary layer is given by the following relationship: f =4.64 Re;1" [1] where 61 is the boundary layer thickness, x is the length along the leading edge, Us is the velocity of the free stream, (i is the viscosity of the fluid (gas in this instance).

Jan 1, 1970

By N. K. Chen, R. Maddin, K. T. Aust

Lattice rotations occurring on the tension and compression sides during the bending of molybdenum single crystals at room temperature were followed in detail. Observations were also made concerning crystallite rotations and slip traces. &apos;The results indicated the participation of (110) as slip planes and the<111>as slip directions. Crystallite fragmentation during bending was also noted. EVERAL investigations have been conducted on pure bending of face-centered cubic metals.&apos;-" Structural changes as a result of "bend-gliding," and the importance of bending and constraints during conditions of axial stressing, have been noted.&apos;," It was evident that the mechanism by which the bending deformation occurs is more complex than the case of simple shear. Studies of the bending of body-centered cubic metals from the aspect of the crystallographic mechanism appear to be lacking. Molybdenum single crystals therefore were subjected to deformation by bending in order to determine the reaction of the body-centered cubic lattice to bending stresses. It might be expected for the bending of single crystals that lattice rotation as represented by the axis of compression would proceed toward the pole of the active slip plane and that rotation of the tensile stress axis should indicate the slip direction.&apos; It was believed therefore that a study of the tension and compression sides of bent molybdenum single crystals should give data concerning the operative slip system. Specific problems which are pertinent to the overall picture of plastic deformation were also investigated, e.g., whether the constraints offered by pure bending give rise to the deformation band, and also the nature of the asterism occurring on the tension and compression sides during the stages of deformation. Experimental Procedure Eight single crystal specimens were grown from sintered molybdenum rods 143 in. diameter using the method described previously.6 The purity of this material was reported as 99.9 pct as described elsewhere.&apos; The specimens were % to 1/10 in. in diameter x 7 in. long, with single crystals approximately 1 to 2 in. in length occupying the entire center sections of each specimen. The specimens were elec-trolytically polished using an electrolyte of 300 cc methyl alcohol, 60 cc H2SO4, 130 cc HCl, and a current density of 4 amp per sq in. Each specimen was then loaded in a bending apparatus similar to that used by Yen and Hibbard,&apos; as shown in Fig. 1. The load was applied by means of a spring and screw arrangement through four ball-bearing surfaces set 1.5 in. apart. Either the inner two bearings or the outer two (A, B, Fig. 1) were attached to a movable steel plate (D) through which the load was transmitted. Since the apparatus was mounted on a track of the X-ray apparatus, the tension or compression side of the specimen could

Jan 1, 1954

By Harrison Everett Ashley

(Pittsburg Meeting, March, 1910.) Slimes are usually defined as all material passing a certain-sized sieve, which is invariably the finest sieve employed by each metallurgist in his tests; 100-mesh and 200-mesh have been taken as the limits by different writers in recent volumes of the Transactions ; 150-mesh is common practice. The reader of the first two volumes of Richards 1 is impressed by the inflexibility with which all grinding-arrangements are praised or condemned according as they make little or much slime. Until recently, one of the principal aims of the mill-man was to keep down the slimes. He was obliged to have separate apparatus for sand- and slime-treatment. The latter was sometimes puzzling and erratic. Often it was found most .economical to omit treatment of the slimes entirely, or to store them in immense basins or heaps, waiting for possible future improvements in methods of treatment. Recently, improved processes of agitation and filtration have made possible the successful extraction of values from many slimes. The tendency, when the metal is exceedingly finely disseminated, is to cut out all sand-treatment and to slime everything. These processes have been almost entirely mechanical in their conception, and it is the object of the present paper to show that certain chemical considerations may greatly hinder or assist their successful operation. It has long been a source of dissatisfaction to chemists that their analyses fail to be of much help to the clay-worker; that it is a matter of common occurrence for clays of closely identical ultimate chemical constitution to show the widest possible divergence in physical properties. * Published by permission of the Director of the U. S. Geological Survey. Ore Dressing (1903).

Aug 1, 1910

By R. E. Ogilvie, J. I. Goldstein

The a and y solubility limits in the Fe-Ni phase diagram have been redetermined at temperatures above 500°C. Both a diffusion-couple and a quench and anneal technique were used. The solubility limits were measured with an electron-probe micro-analyzer. The nickel concentration at the phase boundary is increased below 700°C and the a solid-solubility range is much larger than had been previously measured. The solubility limits are also extrapolated to 300°C. It is suggested that the solubility of nickel in a FeNi reaches a maximum in the temperature interval 400o to 500°C. THE equilibrium diagram has been of great use in the field of meteoritics, where the phase relations between kamacite (a) and taenite (?) in metallic meteorites can be described by means of the Fe-Ni diagram. The study of metallic meteorites by electron-probe microanalysis has cast some doubt on the accuracy of the presently available Fe-Ni diagrams.1, 2 Recent thermodynamic studies of the Fe-Ni system also suggest that the diagram may be in error.3 For these reasons the high-temperature (800o to 500°C) part of the diagram was redetermined. INTRODUCTION The currently accepted Fe-Ni diagram is that of Owen and Liu.4 Above 910°C, there is a region of complete solid solubility, ? (fee). Below 910°C, the a (bee) phase is stable in pure iron. The effect of increasing amounts of nickel is to stabilize the y phase. The phases that form when Fe-Ni alloys are heated or cooled bear little relation to the equilibrium diagram. If an alloy is cooled from the y state and held at a temperature within the a + y field, no evidence has been found for the occurrence of the y - a transformation.5 If the alloy is cooled to low enough temperatures the y phase breaks down into a supersaturated bee phase called a2. In fact, in alloys over 27 pet Ni the y phase is retained at room temperature.6 The a, phase has the same composition as the original y. The temperature at which a2 forms, the M, temperature, has been determined experimentally.&apos; A state of equilibrium can be approached by cooling below M, to form a2 and then reheating the alloy into the two-phase region of the diagram. The y phase will then begin to precipitate out of the a2 phase and grow. The growth of the phase, however, is quite slow. Using the interdiffusion coefficients -Da of Goldstein et al.,8 it is estimated that it takes about 1 year to grow a 10-µ-wide region of ? at 700°C in a 5 pet Ni alloy. Owen and Liu4 used the technique just described to form the equilibrium phases. The phases present were determined by means of X-ray analysis. The accuracy of their diagram depends on the number of alloys available near the phase boundary at a given temperature. In this study two different techniques were used to determine the a/a + ? and ?/a + ? solubility limits. The inherent accuracy of both techniques was greatly improved over that used by previous investigators because the phase-boundary compositions were measured with an electron-probe microanalyzer. PROCEDURE The two methods used to determine the Fe-Ni diagram are the diffusion couple (D.C.) and the quench-and-anneal (Q.+A.) techniques. In the first method, diffusion couples whose diffusion path goes through a two-phase region of the phase diagram were used. A description of the technique used for making the diffusion couples has been described in a previous paper.&apos; After the diffusion treatment a discontinuity in the resultant concentration vs distance profile was measured. The nickel concentrations in the a and ? phases at the interface of the discontinuity are the solubility limits of the a and ? phases in the phase diagram at the diffusion temperature. If the interface com-

Jan 1, 1965

By T. R. Armstrong

Total crude oil production in Mexico in the year 1941 is estimated at approxi- mately 41,000,000 bbl.; a daily average of 112,700 bbl. and a decrease of about 5 per cent as compared with the daily average of 120,000 bbl. produced in the year 1940, and nearly 11 per cent decrease as com-

Jan 1, 1942

By T. R. Armstrong

Total crude oil production in Mexico in the year 1941 is estimated at approxi- mately 41,000,000 bbl.; a daily average of 112,700 bbl. and a decrease of about 5 per cent as compared with the daily average of 120,000 bbl. produced in the year 1940, and nearly 11 per cent decrease as com-

Jan 1, 1942

By James A. Crookston, William D. Fitzpatrick

Committee C-8 of the American Society for Testing and Materials defines refractories as "material, usually nonmetallic, used to withstand high temperature," and it defines the term refractoriness thus, "in refractories, the capability of maintaining the desired degree of chemical and physical identity at high temperatures and in the environment and conditions of use.” The ability to withstand high temperatures is critical to classifying a material as a refractory. However, as the foregoing definition of "refractoriness" implies, the refractory material may be subjected to and is expected to resist not only high temperatures but a variety of other destructive forces incidental to its use. These include abrasion, impact, thermal shock, chemical attack, and high level loads at high temperatures. The various industrial applications of refractories represent an infinite variety of combinations and degrees of severity of these destructive forces, including high temperatures. Naturally, such a variety of application conditions leads to a need for a wide variety of refractory products to satisfy those conditions. The number of refractory products available is great enough that some form of classification is necessary to a discussion of refractories. Refractories may be classified in several ways, but the most suitable one is based on their general compositions and to lesser degrees on their physical condition and distinctive properties. The primary types of brick include aluminosilicate brick (fire clay and alumina), basic brick (magnesia and chrome ore, alone and in various combinations), silica brick, insulating brick, and special refractories (carbon, SiC, ZrO2, zircon, borides, and nitrides). Brick is, of course, preformed and usually is heat treated at high temperatures before use, but in some instances is unburned and/or chemically bonded, including the relatively new class of tar-bonded magnesia and/or dolomite brick. In addition to preformed refractories, another broad category includes the so-called specialty refractories, which are usually made in the same types as itemized for brick-bonding mortars and high temperature cements, hydraulic setting castables or refractory concrete, plastic firebrick, and ramming mixes. Refractory grains, such as fire clay and high alumina calcines, dead-burned magnesite and periclase, and dead-burned dolomite are also considered refractory products. Bricks are also classified on the basis of physical dimensions, as standard sizes and special shapes. The standard sizes are those which are of definite shape and design and are designated as 228.6 x 114.3 x 63.5 mm (9 x 4 ½ x 2 ½ in.) and 228.6 x 114.3 x 76.2 mm (9 x 4 ½ x 3 in.) straights and series. Bricks of nonstandard sizes are referred to as special shapes. These come about as a result of a particular need for a given furnace or for a particular application. Types and Classes of Refractories Fire Clay Refractories Classification: Fire clay refractories, as the name implies, are composed essentially of fire clay and/or calcined fire clay in varying proportions. The ASTM Standard Classification of Fireclay and High Alumina Refractory Brick, C27-70 (reapproved 1976), gives five classes of fire clay brick designated as super-

Jan 1, 1983

By Huey-Lin Luo, T. R. Anantharaman, William Klement

Two new metastable phases have been obtained in Au-Ge alloys by rapid cooling from the melt. One is of the hep structure with near-ideal axial ratio and may be considered a Hume-Rothery phase. The other has a complex tetragonal structure (a = 11.627Å and c = 22.191Å for the Au60Ge40 alloy) and may be looked upon as a superlattice made up of forty-four fee unit cells. Alloys with the tetragonal phase are charactevized by marked preferred orientation. The metastable constitution of the Au-Ge alloys is discussed with reference to the rate of cooling from the melt. THE Au-Ge equilibrium diagram features a rather deep eutectic near 27 at. pet Ge with a maximum solid solubility of 3.2 at. pet Ge in gold and is similar to the Ag-Ge, Ag-Si, and Au-Si diagrams.&apos; The present study of the effects of rapid cooling from the melt on the structure of Au-Ge alloys constitutes an extension of similar investigations2-8 in this laboratory on alloys of noble metals with germanium and silicon. Nonequilibrium hep structures have so far been obtained by quenching Ag-Ge2, 3 and Ag-Si4, 5 alloys from the melt; an Au-Si alloy of approximately eutectic composition has been found6 to solidify in a noncrystalline state. EXPERIMENTAL PROCEDURES Twenty-nine alloys prepared from gold of 99.9+ pet and germaium of 99.999 + pet purity were investigated. The method of alloy preparation, the technique of rapid quenching from the melt, and the X-ray diffraction procedures were similar to those described previously.3, 4 Essentially, small sections (-20 mg) of thin alloy rods (-1 mm diam) were rapidly heated up to 1200°C in a graphite crucible under argon and the liquid alloy ejected by a helium blast onto a polished copper strip held on the inner periphery of a rotating wheel. The resulting thin foils were examined with or without the copper substrate in a G.E. X-ray Diffractometer or in a 114.6-cm-diam Debye-Scherrer camera. RESULTS Table I gives the constitution of liquid-quenched Au-Ge alloys, as estimated from intensities of X-ray reflections from the corresponding phases, viz., a (fee solid solution), ß (hep phase), ? (tetragonal phase), and diamond-cubic (dc) germanium. Despite unavoidable variations in the quenching rates, the proportions of the phases in different alloys were generally reproducible and stable at room temperature. There was no evidence of quenching strain, since the a doublets were well-resolved in high-angle Debye-Scherrer reflections in all cases except where stacking faults in the a fee and ß hep structures contributed to X-ray line broadening. The ß and ? metastable phases showed signs of decomposition to the equilibrium a phase and germanium only after a few hours annealing above 100°C. A few alloys around the eutectic composi-

Jan 1, 1965

By F. W. Glaser

Metal-carbon-boron powder mixtures were hot pressed and the resulting specimens were studied by X-ray diffraction. It was found that regardless of the starting combination of the metal, carbon, or boron powders, a metal boride phase was always the major component in these samples. In the absence of carbon the boride phase formed on hot pressing depended only on the amount of boron present. Two new phases of the system Ti-B were found. They are Ti2B and Ti2B5. The existence of a controversial face-centered cubic phase of formula TiB was confirmed. Electrical resistivities were measured for various boride phases. It was found that the diborides are generally better conductors than the monoborides of the same metal. THE carbides and borides of the transition elements have very high melting points, in the range 2500° to 4000°C, and are therefore of interest as high temperature materials. The literature on the stability or chemical reactivity of these carbides and borides is very scarce. Various investigators&apos;-" have demonstrated a relative instability of certain carbide phases in the presence of boron or boron-containing substances. In a recent publication, Glaserl demonstrated the stability of zirconium-boride (ZrB,) in the presence of carbon at temperatures in excess of approximately 2900°C, while during a preliminary investigation of boride phases, Steinitz&apos; concluded that the diborides are stable in the presence of carbon while the monoborides of the fourth and fifth group are not, forming diborides plus carbides instead. Nelson, Willmore, and Womeldorph" have elaborated on the reaction B,C + 2TiC = 2TiB, + 3C, which was known to occur because of a relative instability of B,C and the great tendency towards TiB, formation at relatively low temperatures (approximately 1200°C). A similar study, involving as starting materials TiO, and B,C and resulting in TiB,, was recently described by Honak4, who observed the beginning of an exothermic reaction of a Ti0,-B,C powder mixture, which, when preheated in a hydrogen atmosphere to approximately 950°C, was carried to about 1600 °C by the heat of reaction. To shed more light on reactions of this type (Metal-C-B), the final product apparently always resulting in a boride phase at the expense of a carbide phase," a systematic investigation was started * Boride phases of various metals, as reported to date, are listed in Table I. and the following is an account of some of the results that were obtained. Materials, Preparation of Samples, Testing Methods The raw materials employed for this work consisted of various carbide, boride, and metal powders. as well as of boron and graphite powders. In cases where commercial grades of carbides were considered unsuitable because of low purity or excessive amounts of graphitic carbon, such carbide powders were prepared by this laboratory. The procedure for the preparation of carbide powders (zirconium carbide, titanium carbide, tantalum carbide, and niobium carbide) consisted of mixing graphite and the respective metal hydride powders in stoichio-metric proportions and subsequent heating of such mixtures in a hydrogen atmosphere in carbon crucibles. The heating was by high frequency to temperatures ranging between 1700" and 2100°C. The resulting carbide was then comminuted and screened to the desired particle size. ZrB, and TiB, powders were produced by the electrolysis of fused salt baths, according to the method described by Andriex.. The borides of niobium, vanadium, tantalum, molybdenum, chro-ium, and iron were obtained by mixing the respective metal and boron powders in the desired proportions. Such metal-boron mixtures were heated in a high frequency furnace to form boride powders. For each metal-carbon-boron group (Tables I1 through XI) a metal, its hydride, carbide or boride were mixed with carbon, boron or boron carbide powders. The additions of carbon, boron or boron carbide powders to any of these metals or metal compounds were calculated to satisfy a particular carbide or boride phase that according to the literature (Table I) had definitely been established by X-ray diffraction work. Samples of powder mixtures were hot pressed in graphite molds that were heated by direct conduction. The specimen dimensions were approximately 2.5X1X1 cm. Hot pressing temperatures were measured optically and maintained for approximately 30 sec under a constant pressure of about 1.3 ton per sq in. Wherever possible, an attempt to obtain maximum specimen density was made by temperature variation. Electrical resistivity testing was done by measuring potentiometrically the voltage drop over a length of 1.5 cm for a current of 10 amp, at room temperature. To obtain electrical resistivities for specific carbide or boride phases, values were plotted as a function of the respective sample densities

Jan 1, 1953