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By T. A. Read, H. K. Birnbaum

STRESS-induced twin boundary motion in the AuCd ß'phase (52.5 at. pct Au 47.5 at. pct Cd having an orthorhombic structure (space group D h)' was discussed for the case of transformation twins.' Mechanical twinning, i.e., introduction of new twin orientations by the applied stress was sometimes observed in the ß' phase which contained two transformation twin orientations. (See Ref. 2 for experimental details.) The transformation twins terminated at the mechanical twin boundary; the two initial orientations having been twinned to the new orientation as shown in Fig. l. In a "hard" tensile instrument, formation of each mechanical twin was accompanied by a sharp drop in load (and an audible click) resulting in a typical serrated load-extension curve. The mechanical twins increased in width and additional twins were nucleated as the load increased until the entire specimen was occupied by the mechanical twin orientation. A restoring force2 caused the mechanical twin and transformation twin boundaries to return towards their initial configuration on releasing the load. In a stabilized specimen* the total strain introduced by mechanical twinning was recovered on releasing the applied stress. Initial and final microstructures were identical. After decreasing gradually to a critical width, mechanical twins suddenly disappeared accompanied by audible clicks and sharp increases in load (measured at a constant rate of crosshead motion in a "hard" machine). Mechanical twins could be stabilized in any position by maintaining the applied load for a sufficient time, indicating a decrease in restoring force with time similar to that observed for transformation twins.' Orientation relationships for mechanical and transformation twins were determined and are shown in Fig. 2. The mechanical twin, T3, was twin related to transformation twin, T,, by reflection in (1ll), and twin related to transformation twin, T2, by reflection in (111)T, . T, was related to T, by reflection in(ll1)T1 . The "average" mechanical twin composition plane lay several degrees from (1ll), . In every specimen which mechanically twinned, orientation T3 was most highly favored by the applied stress, i.e., it was the twin orientation which provided the greatest change in length compatible with the direction of applied stress

Jan 1, 1961

By J. J. Gilman

The dependence of ortho kink-band formation on crystal orientation, on temperature, and on the conditions at the ends of a specimen is described. Load-compression curves for crystals that kink are presented. The reversibility of ortho kinking is demonstrated, and motion of kink planes through crystals is shown. Finally, a phenomenological explanation of ortho kink-band formation is given. ORTHO compressive kink bands (Fig. la) were discovered for the case of metals by Orowan,1 and they have been discussed by Jillson2 and in some detail by Hess and Barrett." However, their history in nonmetallic crystals is considerably older. It is not known when they were first observed in mineral crystals, but descriptions of them may be found as early as 1885.* They were sometimes mistaken for twins.6 Mügge seems to have been first to realize their true nature and importance. He reviewed the knowledge of them for a wide variety of crystals: anhydrite, antimonite, kyanite, vivianite, gypsum, mica, graphite, molybdenite, barium bromide, cal-cite, sodium nitrate, and barite. Also, the "twins" observed by Bridgman7 in sapphire crystals resemble kink bands more than twins. The terms ortho and para kink bands have been coined in order to distinguish between two types of kink bands which seem to represent the same geometrical configurations, but which have different origins and sizes. Ortho kink bands are the usual type that form at low loads when zinc crystals are compressed parallel to the basal plane. Para kink bands form at much higher loads in highly deformed crystals. These terms refer to compression phenomena and are not to be confused with analogous tension phenomena such as tensile kink bands and deformation bands. Mügge described the geometry of ortho kink bands quite completely for kyanite crystals, Fig. lc. The geometry observed for cadmium by Orowan1 and for zinc by Hess and Barrett3 is very similar to that described by Mügge. (See Hess and Barrett for a complete discussion.) Mügge associated kinking with bending combined with basal slip. He pointed out that the kink angles

Jan 1, 1955

By Bernhard Blumenthal

A melting process was developed by which high purity electrolytic uranium crystals can be converted into sound ingots without serious contamination. Careful preparation of the crystals, melting in a high vacuum, and directional solidification led to a metal of better than 99.993 wt pct purity. The metallographic examination showed a substantially clean metal with only residual amounts of a second phase which responded to heat treatment. The density of high purity uranium is 19.05 g per cu cm. FOR the study of the fundamental properties of uranium and its alloys, a base metal of high purity is needed, and an investigation was started at Argonne several years ago to examine various methods of preparing high purity uranium. Early success was achieved by the fused salt electrolysis process in preparing high purity metal in crystal form,' and the problem became that of melting the crystals into ingots without contamination. The electrolytic crystals contained a few ppm of heavy metals, less than 15 ppm C, and large quantities of potassium and lithium in the form of trapped fused salt electrolyte. On melting in vacuo in magnesia or urania crucibles, the metal picked up considerable quantities of carbon and some iron, copper, and silicon but lost most of the potassium and lithium. It was not metallographically clean, and much research work was necessary to determine the optimum melting conditions for minimum contamination. This work has resulted in the development of a melting procedure by which sound ingots can be prepared with only a limited contamination by carbon. In the section entitled "Melting of the Electrolytic Metal," the equipment and processes that have been developed for melting the electrolytic crystals are described in detail. The "Metallography" section is devoted to the metallography of high purity uranium, and the section "Density of High Purity Uranium" contains the first property measurement on high purity uranium—a density determination. Melting of the Electrolytic Metal Apparatus—Melting and Vacuum Equipment: Uranium may be melted in either resistance or induction-type electric furnaces. Resistance heating was found to be preferable to induction heating for melting down the electrolytic crystals, for several reasons: Although induction heating permits fast heating, this advantage cannot be utilized when crucibles with limited thermal shock resistance are used. Temperature control of induction-heated melts is poor, and the stirring effect increases the rate of contamination of the melt from the crucible, slag, or the furnace atmosphere; liquation and controlled solidification are difficult to carry out. Because of these objections, resistance heating was chosen. A 10 kw Globar resistance furnace of low heat capacity permitted uniform heating of a uranium charge to its melting point in 45 min. The furnace was insulated with a removable aluminum shield so that, after melting and removal of the shield, the furnace could be quickly cooled to room temperature. The furnace was equipped with an apparatus for directional solidification; see Fig. la. The crucible was suspended in the center of the furnace; and when the melting was complete, the crucible was slowly lowered into the cold end of the furnace tube. Solidification took place upward from the bottom; and when the rate of lowering was small, a sound ingot free from sink holes and shrinkage cavities was produced. To remove the large quantities of gas contained in the electrolytic crystals, melting in a high vacuum is mandatory. While, from the standpoint of contamination, melting in a highly purified argon atmosphere appeared feasible, lithium and potassium removal proved less efficient under such conditions. Also, a vacuum-melted ingot has a cleaner surface than one melted in an inert atmosphere. The vacuum system must be of very high pumping speed and leak tight so that a vacuum of less than 2x10-7 mm Hg is attainable when the system is pumped down cold. The system must have a high pumping speed, so that the pressure maximum which may rise to 1x10-4 mm during heating can be quickly reduced

Jan 1, 1956

By D. K. Deardorff, Earl T. Hayes

An improved technique is described for the accurate determination of melting points of metals in the temperature range 1500' to 2500°C. The improvements consist of gradient heating and refinements in cavity preparation to obtain true black-body conditions. The melting point of hafnium, as determined by this method, is set at 2222°±30°C. This is over 200°C higher than one of the values reported earlier. Discrepancies in the previously reported values for the melting point of hafnium are explained. The melting points of zirconium and titanium are found to be 1855O±15OC and 1668°±10°C, respectively. Reproducibility of results was ±3OC. PRIMARILY, this report is being made to correct a 1975o±25oC1 value for the melting point of hafnium, and to present an improved method for melting point determination. As a matter of interest, the technique was also applied to determining the melting point of zirconium and titanium. There is a wide difference in the melting point of hafnium as reported by competent investigators. DeBoer and Fast' reported a value of 2230°±500C, and later work of McPherson, as reported by Aden-stedt, placed the melting point at 1975o±25oC. Zwikker" calculated that the melting point was 2430°C. In resolving this discrepancy, improvements were made in the procedure for determining the melting points of refractory metals in the range 1500" to 2500°C. Special techniques are necessary to determine the melting point of these metals accurately. Since the metals oxidize in air above 900oC, the determinations must be carried out in an inert gas or vacuum. Also, every known refractory contaminates them to some extent. In addition, the attainment of true black-body conditions is always difficult. DeBoer and Fast actually determined the melting point of Hf-Zr as-deposited iodide process alloys and extrapolated to 100 pet Hf. McPherson employed a tungsten wire suspension similar to that described by Schramm, Gordon, and Kaufmann,1 and used hafnium metal produced by the iodide process. There was no apparent reason for the wide spread in melting points, since the quality of the metal was very good and the work was under the direction of highly competent investigators. Retracing some of this work revealed one understandable error in the case of hafnium, and led to development of an improved technique for determining melting points of reactive metals in the range 1500" to 2500°C. Material Crystal bar hafnium, zirconium, and titanium were used for this work, with analyses as shown in Table I. The hafnium was of the highest purity that has been prepared, though not as pure as the zirconium and titanium. At the instigation of the Pittsburgh Area Office of the AEC, a special lot of Hfo2 was purified, converted to the tetraflouride, and bomb-reduced with calcium, all at the Iowa State College Laboratories. This crude hafnium metal was then refined by the iodide process at the Foote Mineral Co. All specimens were cut with a Sic wheel, ground, pickled, and dried carefully before use. Procedure General Tungsten Suspension Method—In this method, a small specimen is suspended by a fine tungsten wire in a slender, induction-heated tungsten or tantalum cylinder, all under vacuum. The melting point is determined by heating the specimen slowly while observing it with an optical pyrometer focused upon a portion of the specimen that radiates as a black body. This means that there should be no reflected radiation, either from colder objects near the pyrometer or from hotter surfaces. The induction heater would be the only hotter object and, if it is slender enough, the specimen will attain the same temperature. Often a slender hole is drilled into the specimen as a source of black-body radiation. This was the general method used by McPherson in obtaining his value of 1975°C for the melting point of hafnium. In trying to repeat this work, it was found that the points of contact between the

Jan 1, 1957

By H. D. Mellom

The direction of the optic or "C" axis of a uniaxial metal crystal can be found with the metallurgical polarizing microscope by examining two planes of section on the crystal. Complete orientation of the crystal can be determined if a twin crystal is common to both surfaces of section. Examination of Uniaxial Metals in Polarized Light-Rotation of a uniaxial crystal surface, properly prepared,' shows four extinction positions 90 deg apart, provided that the surface is not close to a basal section. One of these two extinction directions corresponds with the projection of the optic axis on the surface of section. This direction can be found if the 'sign of the rotationr2 for the metal is known. To find the direction of the optic axis, the crystal is rotated on the stage 45 deg from an extinction position. With a crystal of negative sign, to produce extinction the analyser must be rotated towards the optic axial direction and vice versa. All sections, other than basal, of a uniaxial crystal have the same sign of rotation and, since there is no change in sign with wave length, white light can be used for these measurements. U the sign of the rotation is not known the presumed direction of the optic axis can be determined assuming a positive sign. A third section cut normal to the presumed optic axis will show basal extinction if the assumption is correct. It is convenient to use the edge between the two planes of section as a fiducial direction common to both surfaces. The optic axis can be related to this direction. If the above measurements are repeated on the other plane of section and the angle between the two section planes measured, stereographic projection can be used to determine the orientation of the crystal optic axis. Determination of Complete Orientation—The complete orientation of a uniaxial crystal can be determined provided that a twin crystal is common to both surfaces of section, and the twin law is known. The orientation of the twin crystal optic axis can be determined in the same way as for the main crystal

Jan 1, 1962

By F. V. Lenel, G. S. Ansell, E. C. Nelson

IRMANN'S' discovery that extrusions of fine alumi-num-flake powders possess remarkable high-temperature strengths has led to the production of a new class of engineering materials whose properties are derived from a fine dispersion of oxide particles in a metallic matrix. A fundamental study of such materials and elucidation of the mechanism whereby the oxide parti-

Jan 1, 1958

By M. C. Lavine, H. C. Gatos, R. S. Mroczkowski

The structure of GaAs-Ge heterojunctions prepared by a back-melting process was studied by X-rav diffraction, melallographic, and electron-micro-analyzer techniques. The boundary region between the GaAs and the germanium was found to consist of a four-layered band. A discontinuity between the two inner bands was brought about by solidification occurring simultaneously from the GaAs and the germanium side. The junction material OH either side of the central discotinuity was found to have grown as a single Crystal with the orientation of the parent substrate. It was found essential to determine the GaAs-Ge phase diagram. This was accomplished by a series of alloys containing 0 to 40 wt pet GaAs. Over this range the diagram is a simple eutectic with the eutectic point at 74 wt pet Ge and 845 C. The limiting- solubility of Coils in germanium was found to bc less than 5 wt pet. The Kossel divergent-beam X-ray technique was used to study orientation relationships across the boundary lager. A great deal of effort has been recently expended on the electrical properties of heterojunctions. Little attention, however, has been given to the metallurgical variables which may be of importance in the formation of the junctions. The present work was undertaken to investigate some of the more pertinent metallurgical aspects of interface-alloyed GaAs-Ge heterojunctions. Such parameters as solidification processes, composition gradients, orientation effects, and the accommodation of crystal structures have been the objects of study in this program. EXPERIMENTAL PROCEDURES Phase Diagram. The phase diagram of the GaAs-Ge pseudobinary system was determined by means of cooling curves. A Bridgeman furnace was employed with the center section maintained at 1200°C and the upper section at 800°C. Melts varying from 0 to 40 at. pct in GaAs content were used. Each charge was incapsulated under a two-thirds atmosphere of argon in a carbon-coated quartz tube. Five min in the furnace hot zone provided rapid melting and homogenization. The capsule was then removed to the cooler portion of the furnace. Cooling times were of the order of 3 min. The resulting homogeneous microstructure indicated that equilibrium conditions were obtained by this procedure. Formation of Heterojunctions. Rectangular parallelepiped single-crystal samples were cut with the large faces of either the (l00), (110), or (111) orientation. The remaining four surfaces were also cut to specific orientations to permit accurate relative orientation of the GaAs-Ge pairs. The

Jan 1, 1965

By E. E. Schumacher

IN a laboratory devoted to the furtherance of the science of communication, the breadth and variety of the problems encountered are challenging to a metallurgist. In my own long association with the Bell Telephone Laboratories, our metallurgical group has dealt with a vast number of alloys, both ferrous and nonferrous, and with many materials in the border range of metallic properties., The objectives of our design engineers embrace every phase of telephony, and the properties they desire in metals are generally unusual, frequently unique, and often conflicting. Commercial alloys have not always been available with properties to meet a special requirement in the telephone plant, and many alloys have had, therefore, to be custom-made. Anyone who has had the opportunity to observe the results of tests of all types on numerous such alloys, both commercially and specially produced, or who has taken part in the development of an alloy of some necessary, but new or unusual combination of properties, could not fail to be impressed by the frequency with which a desired goal has been reached or adequately approached through the presence of decimal quantities of alloying elements. Nor could One fail to be struck by the occasions when, conversely, of a minute amount Of impurity element has rendered an alloy useful. It is no exaggeration to state that the content be- hind the decimal point often spells success or fail-ure, and the critical content is sometimes far to the right of the decimal point as I shall later illustrate. The great effects of small additions have long been recognized and have, in fact, been described by many investigators. The importance of the minute, however, cannot be overemphasized. Therefore, in this discussion, I propose to illustrate for you some of the startling instances, drawn primarily from our own experiences, wherein small, even vanishingly small, proportions of stranger elements in otherwise common aggregates have com-pletely changed the customary pattern of behavior. No one property has a monopoly as to being dis-proportionately affected by minor elements. Nearly all properties are affected, but there is time here to include only a selected few. I have chosen, there-fore, three of general interest: strength, magnetic, and electrical properties. I shall inquire into both the mechanisms and consequences of these disproportionate effects and try to share with you the fascination and challenge of this field. Strength Properties of Lead and Lead Alloys To illustrate the effect of the minute on strength properties I have selected lead and its alloys, with which we have worked extensively for many years. Lead has the advantages of being available in quantity in a state of high purity and of being fairly easy to melt and fabricate without contamination to any extent detectable by the spectrograph. It has the interesting disadvantages of recovery and recrystallization at room temperature. Whether I

Jan 1, 1951

By H. L. Luo, P. Duwez, C. C. Chao

BY rapidly cooling alloys from the liquid state, it is possible to obtain solid solutions beyond the equilibrium concentrations, provided that the components are miscible in the liquid state. Typical such cases include CU-Ag,1 Pt-Ag,2 CU-Rh,3 nickel-base alloys with silicon, germanium, and tin,4 and cobalt-base alloys with silicon, germanium, tin, aluminum, and gallium.5 In the present note it is shown that terminal solid solutions can be appreciably extended in A1-Mg alloys. The alloys whose compositions are given in Table I were prepared by induction melting in tantalum crucibles under an argon atmosphere. Both aluminum and magnesium were 99.99 pct pure. Weight losses after melting were less than 0.2 pct for the aluminum-rich alloys and less than 0.4 pct for the magnesium-rich alloys. From these weight losses it was estimated that the compositions of the ingots after melting were accurate within 0.3 at. pct for aluminum-rich alloys and within 0.5 at. pct for magnesium-rich alloys. As indicated in Table I, the compound AuCd10 in which there is an ionic contribution to the bonding. The authors thank Dr. J. H. Westbrook of the Research Laboratory, General Electric Co., for arranging the supply of materials and for his interest. This investigation was supported by the U.S. Atomic Energy Commission under Contract AT(30-1)-1002, Scope I. some of the alloys were also chemically analyzed and the results confirmed the limits of uncertainty based on weight losses. Rapid quenching from the liquid state was achieved by techniques previously described.8,7 The principle of these techniques consists of propelling a small globule of liquid alloy (about 20 mg) onto a copper substrate. The thin foils easily detached from the copper substrate were analyzed by X-ray diffraction using a Debye-Scherrer camera 114.6 mm in diameter and either nickel-filtered copper Ka or iron-filtered cobalt Ka radiation. Since the back-reflection doublets were resolved, the wavelengths used in the computatioas were CuKal = 1.54050A and CoKa1= 1.78892A. Lattice parameters were obtained by extrapolation of the Nelson-Riley function, and the tables compiled

Jan 1, 1964

By S. A. Mersol, C. T. Lynch, F. W. Vahldiek

The microhardness of single crystals of tungsten disilicide has been investigated by the Knoop method. The average random room-temperature hardness of the WSi, matrix was 1350 kg per sq mm. Hardness crnisotropy was noted with respect to plane and indenter orientation as determined by single-crq.stal X-rny studies. Annealing at 1600" and 1800°C decreased the average hardness to 1310 and 1230 kg per sq tnm, respectively, and produced a second phase identified by X-ray diffraction and electron-microprobe analysis to be wSio.7. Ball-impact experiwzents produced rosettes at 850°C. Optical and electron microscopy showed evidence of slip and cross slip and twinning produced by microhardness indentations. Prismatic (100), [001] slip was found and cor~elated with hardness data. THE present study was undertaken to investigate the hardness anisotropy of as-grown and annealed single crystals of tungsten disilicide. The existence of the silicide WSiz in the W-Si system has been well-established and its structure thoroughly investigated zachariasen2 found WSi, to have a tetragonal C type of structure, similar to that of MoSi, with lattice parameters a = 3.212A, Kieffer et al. studied the W-Si system and measured the density and microhardness (at a 100-g load) of both polycrystalline WSi, and WSi,.,. The values found were 9.25 g per cu cm and 1090 kg per sq mm for WSi,, and 12.21 per cu cm and 770 kg per sq mm for WSi0.7, respectively. According to Samsonov et a1.5 the microhardness of polycrystalline WSi2 is 1430 kg per sq mm (at a 120-g load). EXPERIMENTAL The WSi, single-crystal boules investigated in this paper were grown by a Verneuil-type process using an electric arc by the Linde Division of the Union Carbide Corp.6 The largest specimens were 8 mm in diameter by 16 mm long. The crystals had an average density of 9.01 g per cu cm with a tungsten • silicon content of 99.9 wt pct. The major impurities were: 87 ppm O, 41 ppm N. 54 pprn C, 500 ppm Zr, 50 ppm Na, and 50 ppm Mn. The crystals were silicon-poor, the average silicon content being 22.20 pct (stoichiometric value is 23.40 pct), and tungsten-rich, the average tungsten content being 77.70 pct (stoichiometric value is 76.60 pct). As-received single crystals were ground and analyzed by powder X-ray diffraction technique using Cu Ka radiation. Laue and layer line rotation patterns were obtained on cleaved sections of WSi, single crystals. Electron-microprobe traverses of representative crystals were done using a Phillips-AMR electron microanalyzer. Carbon replicas were used to prepare electron micrographs. This work was done with a JEM-6A electron microscope. Prior to the metallographic examination, the specimens were mounted in Lucite and then polished for short times on polishing wheels using 9-, 3-, and 1-p diamond-grade pastes. Finally they were fine-polished with Linde A powder for 24 hr on a Syntron vibratory polisher. The samples were etched with 4H 2 O:1HF:2HNO3, which is a medium fast-acting etchant. The combination 1HF:2HNO3:5 lactic acid is also a satisfactory etchant. Annealing runs for selected specimens were made at 1600" and 1800°C for 3 hr at 1.0 to 3.0 x 10-5 mm Hg. A Brew tantalum resistance furnace with WSi2 powder for setters was used. The WSi2 powder was the same as that used for the crystal growth. Temperatures were measured with a calibrated W, W-26 pct Re thermocouple and a microoptical pyrometer. Powder X-ray diffraction, emission spectrographic, and electron-microprobe analyses were done after the annealing runs. For microhardness measurements a Tukon Microhardness Tester Type FB with a Knoop indenter was used. Although measurements were taken at loads ranging from 25 to 1000 g, the 100-g load was chosen as the standard load. All measurements were taken at room temperature. Only indentations of cracking classes 1 and 2 were considered.' DISCUSSION OF RESULTS Powder X-ray diffraction analysis showed the as-received crystals to be single-phase WSi2. Laue and layer line rotation patterns obtained on cleaved sections of WSi2 single crystals proved them to be tetragonal WS 2 2 The results also indicated that the c axis of the crystal was oriented parallel to the boule or growth axis. Electron-microprobe traverses across the matrix of the as-grown crystals showed them to be homogeneous WSi,. Optical and electron microscopy of etched crystals, however, revealed that they contained minute amounts of the "golden" and the "blue" second phases as opposed to the "white" or WSi2 phase. These two second phases were concentrated in inclusion and etch-pit

Jan 1, 1965

By S. A. Mersol, C. T. Lynch, F. W. Vahldiek

The room-temperature Knoop hardness of single -crystal titanium diboride has been determined. Variations in microhardness have been correlated with orientation and structural imperfections in the crystals. Maximum differences in hardness values of 750 kg per sq mm with respect to orientation and 450 kg per sq mm with respect to substructure were found. The effects of etchants, load on indenter, and degvee of cracking of indentation were studied. The average hardness was determined as 2800 kg per sq mm on the as-grown crystals. Annealing increased the average hardness to 3375 kg per sq mm. ThE microhardness of polycrystalline titanium diboride approximating the stoichiometric composition TiB, has been studied by several investigators. The values for the Knoop hardness number range from 2710 kg per sq mm (Khn = 2710 kg per sq mm) to 3400 kg per sq mm.'-4 Samsonov et al.' reported a Vickers hardness number of 3400 kg per sq mm (Vhn = 3400 kg per sq mm) at 120-g loads for a 98 pct Ti + B material of 96 pct theoretical density. Recently Dunegan measured the Vickers hardness and obtained 2060 kg per sq mm with a 10-g load (Vhnlo = 2060 kg per sq mm), 1425 kg per sq mm for Vhnloo, and 1320 kg per sq mm for Vhnlooo. We have obtained a Khnloo = 3000 * 350 kg per sq mm for hot-pressed polycrystalline TiB2 that was 99.1 pct Ti + B and 96.6 pct of theoretical density. The present investigation has been made on single-crystal titanium diboride boules, and is believed to be the first such study on TiB, crystals of known orientation. EXPERIMENTAL The single crystals were prepared by a Verneuil-type process using an electric arc by the Linde Division of Union Carbide Corp. The largest specimens were 6 mm in diameter by 12 mm long. The crystals were of theoretical density (4.50 g per cm3) with a Ti + B content of 99.8 pct. The major impurities present were: 0.03 pct (by weight) 0, 0.06 pct C, 0.01 pct Si, 0.05 pct Fe, 0.01 pct Cr, and 0.01 pct Ni. The crystals were boron-rich, the average boron content being 31.S7 wt pct (stoichiometric value is 31.12 wt pct), and titanium-poor, the average titanium content being 68.2 wt pct (stoichiometric value is 68.88 wt pct). Titanium was determined gravimetrically by cupferron precipitation and boron by titration of H3B03 from a NaOH-fused specimen after removal of titanium with BaC03. The outer edge of the boule accounted for some of the excess boron. An electron-microprobe analysis showed that the titanium content increased from 60.0 wt pct at the edge to 68.5 wt pct at D, ~ 1000 p, where D, = distance from outer edge (see Fig. 5). X-ray measurements indicate that the boules are single crystals with considerable substructure. Upon etching this structure is seen in the form of a Widmanstatten-type rhombohedral line pattern in the basal (c) plane and a parallel "lines" pattern in the prism (a) planes. A typical WidmanstHtten-type pattern is seen in in Fig. 1. Electron microscopy reveals that the lines are less than 1 p across, which does not allow direct electron probe or microfocus X-ray determination of their composition. The identity of the second phase has not been established. It might be a higher boride or a nonstoichiometric TiB2. On annealing above 2200°C the precipitate is dispersed and dissolves in the matrix. The best etchants found were HF/HNO/HO, KF(CN)$NOH/HO (Murakami's Etch), and A large number of other etchants were previously Studied. The acidic etchant produced an acicular or "needles" pattern in the (a) planes which is an etchant effect on the "lines" pattern.

Jan 1, 1965

By E. P. Sadowski, R. J. Raudebaugh

A study of micro structural characteristics of a 30 pct Cr, 42 pct Ni, 26 pct Fe alloy has been correlated with its behavior in creep and rupture tests at 1400°, 1600°, and 1800°F. Nitrogen pickup observed at 1600" and 1800°F was associated withfissur-ing of specimens prior to fracture. IN studying the high-temperature characteristics of a nominally 30 pct Cr, 42 pct Ni, 26 pct Fe alloy certain anomalies were observed, namely: The alloy had greater stress rupture and creep strength at 1800" than at 1600° F for time periods beyond 550 hr. As placed in test, this alloy was nonmagnetic. After test, some specimens were found to be fairly magnetic. The gain observed in magnetic response was not solely a function of time and temperature of testing. This is a report of the observations and a recounting of the results of studies made in an attempt to explain the behavior of the material under investigation. MATERIAL AND PROCEDURE All test material was taken from a production size heat of an experimental alloy melted and processed using commercial practice and having the following percentage composition: C Fe Cr Ni Ti 0.05 25.83 30.13 41.59 0.41 Al Mn Si S Cu 0.42 0.75 0.45 0.011 0.34 The material was tested in two conditions, namely mill-annealed and grain-coarsened. Mill annealing, which consisted of heating to about 1900° F in a continuous furnace for about 30 min and air cooling, gave a grain size of ASTM 8 and finer. To produce the grain-coarsened condition, mill-annealed material was soaked for 2 hr at 2050°F in a laboratory furnace and air cooled. This resulted in an ASTM grain size of 4 to 5. Rupture tests were carried out at 1400, 1600, and 1800° F in accordance with ASTM procedures. In tests where only rupture life and elongation at fracture were determined, specimens of 0.252 in. diam and 1 in. gage length were used. In tests where creep curves were obtained, specimens were of 0.375 in. diam and 4 in. gage length. Several of the tests under relatively low stress were discontinued short of rupture. Three low-stress tests are still in progress at this writing. Structural changes which occurred during testing were determined from examination of the tested specimens by light microscopy and X-ray diffraction. Diffraction pattern recordings were made with a diffractometer using copper radiation on samples heavily etched in a picric-hydrochloric

Jan 1, 1960

By R. Wachtell

IN order to study better the phenomena at work in various phases of diffusion of the Fe/Si system when compounded and alloyed by powder metallurgy methods, several attacks have been planned. Electrical, hardness and other measurements have been discussed e1sewhere.l,2 The metallographic phenomena will be considered here. Concerning the metallography of the ferrosilicon alloys, some initial discussion is in order. As this laboratory discovered to its chagrin, there are many traps that lie in wait for even the experienced metal-lographer when he deals with alloys of high silicon content. Chief of these is probably the anomalous behavior of these materials under conventional etching techniques. We have before us, for instance, Corson's paper" of 1928, wherein is mentioned and illustrated (probably for the first time) the peculiar "barley shell" structure sometimes found in these alloys. Again in 19412 he same author reports the structure in 5-14 pct silicon irons. Houghton and Becker mention its presence' but make only tentative effort to explain it. In 1943, Hurst and Riley6 discussed the "barley shell" and declared it to be a false structure, resulting from the peculiar and characteristic film-forming propensities of the alloy. At about the same time, Wrazej7 proved correct the contentions of Hurst and Riley, establishing pretty well the mechanisms of formation and identifying the constituents involved. Hurst and Riley6 a1so describe in their paper (1943) a "cracked film" structure, also a pseudo-morph, the precise nature of which they do not suggest. Careful examination of such a film reveals that it not only cracks but begins to curl away from the surface of the metal and peel, much in the manner of flaking paint. The careful microscopist will be able to focus on the topmost edge of the flake, and, by optical sectioning, follow it down to the metal surface. We have been able, by oblique illumination of such specimens, to observe the raised segment of the film, the shadow that it casts, and the mating segment on the metal into which it fits. As late as 1946, Hurst and Riley8 found occasion to correct misinterpretation of this "cracked film" structure in a work by others on a 12 pct Si/Fe alloy. The "barley shell" and "cracked film" pseudo-morphs are not the only ones encountered in studies of these materials. A third pseudomorph, which may, however, have some structural significance is a fine striation of the surface. Closely spaced and parallel striae extend entirely across single grains. Each grain shows a different direction for the striae, and different closeness of the spacing. The regularity of this structure, and the fact that the striae do not cross from one grain to another, suggest that the structure is a film cracking controlled in some way by the crystal plane orientation of the surface being etched. In general it is wise for the metallographer dealing with these alloys to cultivate a feeling of uncertainty not only of the more complex questions of interpretation, but also of the basic question of whether that which he sees is indeed a true representation of the structure. We have taken, as a first test of the validity of a structure, its reproducibility "in situ" after repolishing and re-etching of the samples. Thus we knew the "barley shell" to be false (fig. 1) even before encountering the definitive works on the subject. The Metallography of Silicon Diffusion in Laminate Bars: As a first step in the study of the metallography of the diffusion process, we have pressed a series of laminate bars. These bars, or "sandwiches" as we have termed them, were pressed with covering layers of iron powder, and a "filler" of the Fe/Si master alloy under study. Bar dimensions were approximately %, x 3 x 1/4 in. The total

Jan 1, 1951

By Lawrence H. Van Vlack, Alfred S. Keh

The distribution of sulfur in iron was found to be dependent upon the time and temperature of the treatment as well as the chemical composition of the sulfide. With higher temperatures, the sulfide phase spreads more extensively between the iron grains. A complete network, however, does not form until approximately 1300°C (2372°F), ot which temperature the dihedral angle approaches zero. Silicon has little effect on this change. Aluminum, oxygen, and manganese all modify the temperature of sulfide network formation to a higher value. The microstructures which result and have significance in relation to hot-shortness are discussed. SULFIDE inclusions in steel were recognized by metallurgists before the turn of the century. Considerable attention has been paid to them, and numerous studies have been made on the working properties of steel as affected by them. It has been suggested that ho t-shortness of steel is due to the presence of an enveloping liquid sulfide film at the grain boundaries of the steel in the range of the hot-working temperatures. However, several facts have been observed in commercial steels which do not seem to be completely consistent with this explanation: 1) the frequent absence of continuous films of sulfide as observed under the microscope, 2) the absence of hot-shortness in some high sulfur and resulfurized steels, and 3) the limitation of hot-shortness to a definite temperature range. The role that manganese plays in steel to eliminate hot-shortness is well known, but the theory behind it is still under question. Moreover, the effect of de-oxidizers such as aluminum and silicon, and the effect of oxygen are still not clearly known. In this study, the effect of heat treatments and of additions of manganese, aluminum, silicon, and oxygen upon the distribution of sulfide inclusions in iron was investigated. Both the as-cast and the heat treated structures were studied, using microscopic and autoradiographic techniques. Part of the results were interpreted in terms of Smith's' concept of microstructure. Review of Literature Metallurgists used to consider sulfide inclusions as suspended particles. It was Wohrman' who first pointed out that sulfide inclusions are soluble in liquid iron. Many observed facts may be understood on the basis of this concept of solubility. Wohrman found that small castings contain small inclusions. and larger castings contain larger ones. Also, inclusion sizes are smaller near the chilled surface. For the same size ingot, the average inclusion size in a given heat of steel depends upon the rate of solidification. That is, with slow solidification, enough time for precipitation and agglomeration of these

Jan 1, 1957

By N. Brown, R. Kossowsky

Plain carbon steels with 0.48 and 0.95 pct C were quenched and tempered at 705°C to produce carbide dispersions with spacings on the order of 1 p. The morphology of the structure consisted of a carbide-dislocation network. The strengthening due to the dispersion was found to vary linearly with M-½ where M is the mean free-ferrite path determined by the entire network. No preyield microstrain preceded the upper yield point. After prestraining, the lowest stress at which dislocation movement could be detected was 104 psi; this frictional stress was independent of the dispersion and prestrains up to 6.5 pct. The Cottrell-Petch equation for grain-size strengthening was used to discuss the dispersion strengthening in this investigation. The results support an impurity mechanism for the upper yield point rather than one based on the Johns ton- Gilman theory. IT was pointed out by Gensamerl that the mean free path in the matrix is the important variable which controls the degree of dispersion strengthening. Gensamer's data on steels showed a linear relationship between the yield point and the logarithm of the mean free-ferrite path. The first theory was by Orowan,2 who suggested the mechanism of dislocations bowing between particles, with the resulting relationship that where a is the yield point, P is the distance between particles, and G is the shear modulus. The Orowan equation applies to that range of dispersion where P is large compared to the particle size and the particles are not coherent, so that the matrix is essentially free of internal stresses. As was pointed out by Orowan, there are different degrees of dispersion which, in turn, will influence the strength-controlling mechanism. However, in this investigation, we wish to confine ourselves to the region of coarse dispersion, where the Orowan mechanism should, intuitively, be applicable. It was soon evident that the data, which existed at about the time that the Orowan theory was proposed, did not agree with the Orowan equation in that the actual strengthening was always greater than the theoretical predictions. Thus, Fisher, Hart, and pry3 modified the Orowan theory by suggesting that the Orowan process took place in the microstrain region preceding macroscopic yielding and the subsequent, rapid work hardening in the form of residual dislocation loops around the particles largely determined the observed macroscopic yield point. The F-H-P theory stated that the increment of strengthening due to the work-hardening mechanism was proportional to f3/2 where f is the volume fraction of the precipitate. F-H-P used data by Shaw, Shepard, Starr, and Dorn4 on A1 3-5 pct Cu to support their theory. Roberts, Carruthers, and Averbach5 were the first to make a microstrain study of dispersion strengthening, and they found that for steel the Gensamer relationship was obeyed. Hayman and Nutting6 suggested that in the case of tempered steel 1) the ferrite grain boundary was the primary obstacle and 2) the carbide particles simply formed part of the grain boundary obstacles. They found that the strength varied as G"" where G is the ferrite grain size. Turkalo and LOW' determined the structure of quenched and tempered plain carbon steel using a replica technique; they concluded that the carbide particles did not necessarily lie in the ferrite grain boundaries. When they defined the mean free-ferrite path as being determined by both particles and the ferrite path boundaries, their data obeyed the Gensamer relationship. Meiklejohn and skoda8 showed that the strengthening from iron particles in mercury was a function of the particle size and the distance between particles. Dew-Hughes9 explained the Meiklejohn and Skoda results by the following theory: 1) grown-in dislocations which surrounded the particles were produced by the thermal stresses during the time the material was cooled to the testing temperature, and 2) the observed strengthening was associated with the cutting of the grown-in dislocations by the glide dislocations. Ansell and Lenel10 proposed a theory in which the glide dislocations pile-up or surround the particle until the number of dislocations in the pile-up is sufficient to yield or fracture the particle. The resulting theory says that the yield point varies as p-1/2, where P is the distance between particles. The Ansell-Lenel theory is almost identical to the

Jan 1, 1965

By J. H. Smith, H. W. Paxton

X HE recent investigation of Butler, McCabe, and paxton' on the vapor pressure of manganese gas in equilibrium with Mn-Fe alloys has been extended to include the iron-rich solutions and alloys in the two phase equilibrium region. The study of the phase equilibrium by Kurnakov and Troneva2 and the extensive reviews of the Mn-Fe system by Hansen,3 and Hellawe114 are not in agreement with regard to the location and extent of the equilibrium between the solid phases of fcc y and complex P in the temperature range above the eutectoid temperature of approximately 975Kwhere 0 decomposes to y and bcc a. An evaluation of the existing equilibrium data is confronted with the conflicting results presented by various investigators who used techniques requiring a quenching treatment to retain the high-temperature equilibrium or encountered temperature hysteresis during thermal analyses. The difficulty of suppressing the transformation of the manganese-rich concentrations of the y-solid solution during quenching has been well etablished.' To avoid this uncertainty the phase equilibrium was investigated at temperature by measuring the partial pressure of manganese gas in equilibrium

Jan 1, 1962

By E. R. Gilbert, D. E. Munson

Creep of copper under 75 to 1.50 kg per sq cm stresses at temperatures near the melting point was found to he a complex reaction controlled by three mechanisms acting in parallel. In order of appearance with decreasing temperature, the auerage activation energies, Qc , are 168, 79, and 24 kcal per mole. Stress dependence of the minimum creep rate was found to he an exponential for the two high-Qc processes and a power law for the low-Qc, process. Transition of control occurs from one mechanism to another. The relative transition temperature depends upon the applied stress, and the range oiler which the transition occurs depends upon the difference in the activation energies of the mechanisms. The creep behavior at high temperature is explained by the climb of dislocations through thermally or mechanically formed jogs, CREEP in pure fcc metals at temperatures in excess of one half the absolute melting point is normally controlled by dislocation c1imb.1,2 A climb model which seems applicable, according to an extensive analysis of data,3 was derived by Weert-man.4 This model assumes jog-saturated dislocations and predicts an activation energy, Qc, nearly equal to that for self-diffusion, USd. Although the requirement of jog saturation is restrictive, agreement between theory and experiment seems adequate. Many other theoretical treatments, including an early model by Mott,5 include detailed consideration of jog formation as an initial requirement for climb. These models predict activation energies for creep which differ from those of self-diffusion. Seeger6,7 postulates an observed activation energy related to the stacking-fault energy. Thus, Usd <Qc<5Usd + Uj where Uj is the jog-formation energy. Seeger incorporated qualitatively the influence of the relative numbers of thermal and athermal jogs. Expanding this concept, Shoeck8 explicitly states a function based on formation mode and relative numbers of vacancies: e r ci exp {-uf/k T} exp {-Um/k T} [ 1 ] where Uf and Um are energies of vacancy formation and migration, respectively. The concentration of jogs, Cj, depends upon the manner of jog production. For intersection jogs, Cj is not sensitive E. R. GILBERT, Junior Member AIME, formerly with De- to temperature; for thermal jogs, Cj is proportional to exp {—Uj /k T}. Schoeck regards each mode as a distinct mechanism; therefore, the mechanisms may act together.299 The diversity predicted by theory, surprisingly, has not been substantiated by experimental results. A significant investigation must include the extremes in stacking-fault energy. Extensive creep studies of aluminum10 and nickel,11,12 high stacking-fault energy metals, have been made. Comparable studies on a low stacking-fault energy metal, such as copper, have not. It is the purpose of this paper to report the results of an investigation of the creep of copper under conditions which favor thermal jogs. EXPERIMENTAL Cylindrical compression creep specimens (0.240 in. in diameter by 0.400 in. long) were machined from cold-drawn rods of electrolytic tough-pitch copper containing 0.0007 Mg, 0.002 Fe, 0.001 Ni, 0.0005 Ag, Cd < 0.005, and Pb < 0.005 wt pct impurities. Undetected spectrographically was a nominal 0.04 wt pct 0, which occurs as a Cu2O constituent distributed discontinuously at grain boundaries. Vacuum annealing at 900°C for 15 min produced a stable 0.03-mm average grain diameter. Testing was carried out using apparatus similar to that described by sherby,13 modified by enclosing platens and a push rod in a vacuum cylinder. Normally this arrangement resulted in pressures less than 10 Only a slight surface tarnish, less than 0.0005 in. in thickness, occurred during the test. The applied stress, corrected for atmospheric pressure, was maintained within 2 pct of the desired true stress by the addition of lead shot at fixed strain increments. Creep strain was measured with dial gages as a relative displacement of the upper and lower platens; accuracy of measurement was 0.0001 in. Two creep-test methods were used, the differential or cyclic temperature14 and the isothermal, to obtain creep data at stress levels of 150, 100, and 75 kg per sq cm over the temperature range of 620° to 1032°C. Minimum creep rates were used from both test methods; this was considered proper because comparable temperature tests or cycling to the same temperature gave the same creep rate, within experimental error. The cold vacuum test chamber, with the unstressed specimen in place, was heated to temperature by placing a preheated furnace over the chamber. Temperature equilibrium was attained within 30 min. For the cyclic tests, the stress was removed during the 5 to 10 min necessary to effect the temperature change and re-

Jan 1, 1965

By R. A. Swalin, A. H. Geisler

THIS study represents an investigation of the effect of annealing temperature on the texture of drawn molybdenum, iron, and an iron alloy containing 14 pct Al. The effect of temperature on annealing textures has not been widely investigated in the past. One of the obstacles in making such a study using usual X-ray diffraction techniques results from the

Jan 1, 1957

By A. C. Fraker, H. H. Stadelmaier

WIENER and Bergerl reported the existence of the nitrides Fe3NiN and Fe3PtNwith a cubic L1&apos;2 structure. The present note shows that a similar nitride Fe3PdN can be observed and that the composition range between the iron nitride Fe4N and the nitrides Fe3NiN, Fe3PdN, and Fe3PtN is occupied by single-phase cubic nitrides with structures intermediate between L1&apos; and L1&apos;2. The materials used in this investigation were electrolytic iron of 99.9 pct purity, carbonyl nickel of 99.7 pct purity, sponge platinum of 99.95 pct purity, and palladium sheet of 99.8 pct. purity. Binary alloys of the transition metals were induction-melted in a crucible of fused alumina, powdered, and screened to a particle size of - 200 mesh. The powders were nitrided at 600°C in a gas mixture of ammonia and hydrogen, and the crystal structure was determined with the aid of a Wilson powder camera. The results are compiled in Table I together with information from the paper of Wiener and Berger.1 The compound at the nominal composition Fe3PdN has a diffraction pattern with prominent reflections of a face-centered cubic lattice (h2 + k2+l2 = 3, 4, 8, 11, 12, 16) and weaker superstructure lines (h2 + k2 + 12 = 1, 2, 5, 6, 9, 10, 13, 14). With decreasing palladium content the superstructure lines become weaker and approach the intensity of those of Fe4N. No analysis was made of the intensities, but it can be assumed that the structure is that found by Wiener and Bergerl in Fe3NiN and Fe3PtN, namely L1&apos;2 (perovskite structure) with both palladium and the interstitial nitrogen fully ordered. The dependence of the lattice constant upon the solute atom (Ni, Pd, Pt) content can be explained tolerably well by simple lattice expansion due to subsequent work confirms this suggestion, it will be of significance to order-disorder theory. The author is grateful to Lester Guttman for discussions on this topic.

Jan 1, 1961

By J. S. Kirkaldy, D. G. Fedak

The extra degree of freedonz introduced by a second independetzt concentration in a tenzary system gives rise to the possibility of unstable planar phase interfaces in semi-infinite diffision layer couples. Diffusion paths which pass continuously in and out of too-phase regions of the phase diagram but which are not coincident with tie-lines give rise to super saturated zones which may be either adjacent to or isolated from planar phase interfaces, with corresponding nonisolated or isolated precipitates. Nonisolated precipitates are equivalent to morphological development of the interface. Selected couples in the ternary systems Fe-Ni-Cr and Cu-Zn-Sn have been examined nzetallographically and ifz the latter case both isolated and uo~zisolated precipitates have been observed. The diffusion path in one of the couples showing interface morphological developtnent has been examitled using electron probe microanalysis. IGNORING the spurious effects of preferential grain boundary diffusion and anisotropic crystal growth7 it can be unequivocally stated that it is impossible for a binary isothermal diffusion couple constructed of two unsaturated phases to sustain a nonplanar phase interface. Consider for demonstrative purposes the schematic phase diagram of Fig. l(a). Given an initial flat phase interface, semi-infinite boundary conditions and the internal boundary condition fixed by the constitution diagram, xB = xB(T), the unique solution of the diffusion equation requires a persistent flat interface. No other solution is possible. If, however, we anticipate the involvement of surface tension, then the internal boundary condition must be written xB = xB (T,Y) where r is the radius of curvature of a hypothetical nonplanar interface. The introduction of a new parameter such as r can be equivalent to the release of an internal constraint, so that a stable nonplanar interface may be possible after all. Consider a perturbation to the interface as indicated in Fig. 1(b). According to the Ostwald-Freundlich equation: which is applicable to a system of this constitution, the solubility of A in the matrix ß ,xB (Y), at a point of curvature r is given by where Va is the molar volume of the a-phase, R is the gas constant, and a is the interfacial tension. Since the right hand side is positive, xB(r) >xB cm, as indicated in Figs. 1(a) and 1(c). Accordingly, the perturbation will carry the interface concentration in a direction which is opposite to that which would sustain local equilibrium. Such a perturbation is necessarily moderated and the system is returned to its initial condition. This result is general for all binary couples since we can regard an interfacial perturbation as an attempt to precipitate within un-

Jan 1, 1962