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By B. D. Cullity, A. Freda

It is well known that deformation by cold drawing or swaging produces a kind of preferred orientation called fiber texture in metal wires. Such textures have been extensively studied by means of X-ray diffraction, chiefly by techniques in which the diffracted X-rays are recorded photographically. This paper describes some results obtained with a new method, involving the use of a diffractometer. X-RAY METHOD Since the pole figure of a material having a fiber texture is symmetrical about the fiber axis, there is no need to present the entire pole figure. Instead, texture data can be summarized by means of a curve showing the variation of pole density along any meridian of the pole figure, if the north-south axis of the latter is made parallel to the fiber axis. The position of a pole on the meridian is specified by the angle 9 between the pole and the fiber axis (wire axis). The X-ray method1 used here to measure pole density as a function of 9 actually consists of two methods, referred to as A and B. In method a, X-rays are reflected from the wire surface. Pole densities can be measured by this method from O = 90 deg down to a limiting value which is dependent on the Bragg angle 8 and is typically of the order of 30 to 40 deg. Method B involves diffraction from the wire cross section and covers a range of 0 values from 0 up to an angle sufficient to overlap the measurements made by method A. By a combination of the two methods, a complete pole-density curve can be obtained. However, method A alone, which requires no specimen preparation other than etching, can often reveal enough of the pole-density curve to provide complete information about texture components. Both methods require that the measured diffracted intensities be corrected for changes in absorption and diffracting volume with changes in 9. A11 intensities I reported in this paper are corrected diffracted intensities; they are propor-

Jan 1, 1960

By R. W. Kraft, F. D. George, F. D. Lemkey

From a conszderation of the geometrically possible ways in which an array of lines or linear features in three-dimensional space can depart from a statistically random arrangement, a system was developed to describe and measure different types of anisotrypic arrangements of lines in opaque solids. The technique was applied to a specimen of unidirectionally solidified Al-Cwll, eutectic which contained microstructural imperfections termed lamellar faults. Experimental data obtained from photomicrographs of different facets of a grain of this alloy were then analyzed using this system. It was deduced that the defects, which are essentially linear in natuve, were predominantly parallel to one another but not parallel to the nominal growth direction probably because of some lateral heat flow during solidification. InGOTS of lamellar eutectic alloys which have been unidirectionally solidified under growth conditions which prevent colony formation consist of large elongated grains in which the lamellae are approximately parallel to one another. Deviations from a parallel lamellar structure within a grain are caused by extra lamellae which form occasionally during the solidification process. Three-dimensional analyses of such specimens have revealed that one extra platelet produces two defects, termed 1amellar faults, which can and do grow more or less independently of one another.' The faults are analogous in many ways to edge dislocations in crystals. If a lamella of one phase is considered to be basically two-dimensional, e.g., a surface, then a fault can be defined as the locus of points which delineates an edge (or half the periphery) of one extra lamella as it grows and forms bridges and saddles in an otherwise parallel arrangement oi lamellae, Fig. 1. Were it not for the fact that bridges and saddles occur in the structure, the extra platelet could be defined as the unit defect and the concept of linear faults of positive and negative character around the periphery of the extra platelet would be unnecessary. If an extra platelet is merely inserted (actually it is grown in) between other parallel platelets, the fault line around the periphery of the extra platelet would circumscribe one and only one platelet, namely the extra one, and the faults would lie in a plane parallel to the lamellae. But when an extra platelet merges with adjacent ones and each side grows independently of the other, it is no longer topologically possible to delineate "the extra plate." A fault circuit which either closes on itself or passes out of the specimen still occurs although it no longer will be coplanar with the lamellae. In fact, it is usually not even Planar even if the lamellae are. Under these conditions a concept that a specimen contains N extra platelets of a given size per unit volume would be very artificial and consequently meaningless. Therefore, because of the topological complexity of the alloys, it was deemed necessary to define one-dimensional

Jan 1, 1962

By Stanley L. Lopata, Eric B. Kula

The problem of preferred orientation has been considered in quantitative phase analysis by X-ray diffraction techniques. The average intensity of a diffraction peak can be obtained by integration over the whole pole figure, or by random rotation of the specimen in the X-ray beam, or by a combination of both. For rod or wire, where the pole figure shows rotational symmetry, the intensity is measured along a radius of the pole figure. Where no symmetry is present, rotational symmetry is imparted by rapid rotation of the sheet around its surface normal during measurement. These methods are applied to the determination of aus-tenite in an Fe-31 pct Ni rod and Fe-28 pct Ni sheet. X-RAY diffraction techniques are frequently used for quantitatively determining the relative amounts of two or more phases present in a structure.1,2 Briefly, the method requires a measurement of the intensity of a Bragg reflection for each phase. The method is based on the principle that the intensity of a Bragg reflection, IA, is directly proportional to the volume fraction of the given phase present, CA: T KRACA/2µ [1] where K is a constant independent of the diffracting phase, RA can be calculated and depends on the Bragg angle ? and the particular reflection measured, and µ is the linear absorption coefficient of the mixture.' For a mixture of two phases A and B, IA = RAcA [2] IA/IB=RAcA/Rbcb and CA+ CB = 1 [3] This allows the volume fraction of a phase to be calculated directly from the measured integrated intensities of the two phases present. This method is frequently used for quantitatively determining retained austenite in steels, but can be used for the quantitative determination of in titanium alloys, as well as phases in other alloy systems. The intensities are generally measured by diffracting only from planes whose normals are perpendicular to the specimen surface, with the implicit assumption being made that this intensity is typical of reflections from planes whose normals lie at other angles to the surface. If the specimen contains a texture or preferred orientation of the crys- tallographic planes by virtue of previous mechanical working, however, this measured intensity is correct only if the intensity level fortuitously is one times random. Examination of pole figures for actual rolled sheet shows that the intensity level at the center of the figure is seldom at the random level, and can very easily differ by a factor of five or more, leading to a commensurate error in the calculated amount of the phases. Textures have been commonly found in materials which are cold-rolled or recrystallized after cold rolling. This problem is becoming of greater importance since structural materials in use today, such as the high-strength stainless steels, frequently are further strengthened by cold rolling, and hence contain textures. Although high-strength carbon-containing alloy steels generally do not contain textures, some of the newer thermomechani-cal treatments will impart a texture even in these cases. In the case of hexagonal metals, such as titanium, probably every wrought sample will show some anisotropy. Since the standard X-ray techniques are inapplicable in all these cases, there is a clear need for some method of taking this anisotropy into account. The specific problem is to obtain the integrated intensity of a Bragg reflection for each phase, averaged over all orientations of the sample with respect to the X-ray beam. This can be done either by purely mechanical means or by a mathematical averaging. Mechanically, a small spherical speci-

Jan 1, 1965

By D. C. Jillson

THE data on the critical resolved shear stress for zinc single crystals published by Rosbaud and Schmid1 showed many irregularities, as seen in fig. 1. There have been subsequent investigations,2,3 but no one has repeated this work with a wide range of orientations, with specimens of a higher degree of purity and perfection, and with improved experimental techniques in order to determine whether these variations were entirely experimental or whether they might have been due, at least in part, to some feature of the mode of deformation not known or understood. Mathewson4 and others have long felt that the glide observed in a [100]* direction between two (001)* planes in the zinc lattice might be actually the net effect of many consecutive unit glides or movements in two alternating [210] directions (fig.

Jan 1, 1951

By B. Ancker, E. R. Parker

The small-angle dislocation-boundary density of nickel and some of its alloys was investigated as a function of strength. It was found that the strength is a linear function of the density for pure nickel and for Ni-Ti alloys, but not for Ni-Co alloys. The results are interpreted in terms of dislocation theory. THE processes by which metals deform are still poorly understood, even though there has been a concentrated research effort in this field during the past fifteen years. Single crystals and individual grains in polycrystalline aggregates often develop a subgrain structure when deformed. Substructure formation occurs as a result of local plastic bending and its development is augmented by annealing at moderate temperatures.' Investigations made to date, largely qualitative in nature, have shown that the substructure is composed of a network of small-angle boundaries. The purpose of the present paper is to present some quantitative data on substructure (obtained by X-ray techniques) and to correlate the results with the strength of the metal in which the substructure existed. Nickel and some of its alloys were used for the experiments. The presence of certain kinds of foreign atoms greatly alters the formation and effect of the substructure. Polycrystalline tensile specimens were prepared from high purity nickel, two Ni-Ti alloys, and two Ni-Co alloys; the composition of these materials is given in Table I. Tensile specimens having a reduced section 1/4 in. diam by 3 in. long were machined from materials which were refined, cast, forged, and swaged in the authors' laboratories. The specimens were annealed for 1/2 hr at 1150°C, a treatment which produced grains having a diameter of approximately 0.5 mm. This rather large grain size was used to facilitate later examination of the substructure. Duplicate specimens of pure nickel and of each alloy were strained various controlled amounts in tension at room temperature and recovered for 1 hr at 800°C. (This strain induced prior to recovery will be referred to hereafter as the "prestrain.") One of each of the recovered specimens was pulled in tension at room temperature to obtain the stress-strain curves. The other bar was cross-sectioned in the gage length, carefully polished, and deeply etched to remove all cold-worked material. This specimen was then examined by the X-ray microscopic technique2 to determine the nature and magnitude of the substructure. The use of duplicate specimens permitted a direct correlation between tensile properties and substructure characteristics. X-ray microscopic studies were made with an apparatus constructed to have an average resolving power of about 4 microns. The monochromatic radiation reflected from grains which were oriented favorably was recorded on fine-grain photographic plates. These were examined at Xl00 which represents the true magnification, since there was essentially none in the X-ray system. This permitted detection of angular differences between subgrains of as little as 10 min. A very large number (over 100) of grains for each specimen was examined at this magnification and the number of substructure boundaries per grain was tabulated. An accurate relative measure was obtained under fixed experimental conditions, e.g., same grain size, resolving power, magnification, etc. The greatest source of error lay in counting the subboundaries, but this proved to be surprisingly small when repeated counts were made by different investigators. Furthermore, as the subsequent analysis indicates, whatever human error was introduced was not significant. Analysis of the Substructure Data When the number of substructure boundaries per grain was plotted against the frequency, skew (asymmetrical) distribution curves of the type shown in Fig. 1 were obtained. These are smoothed curves for the 1 atomic pct Ti in nickel alloy subjected to various prestrains and are representative of the curves obtained with all the metals. Since skew-distribution curves cannot be analyzed by elementary statistical methods, detailed investigations were undertaken in order to 1—test the consistency

Jan 1, 1955

By J. H. Waxweiler, L. C. Ikenberry, R. J. Bendure

Carbon which is present as insoluble carbides in austenitic stainless steels can be measured quantitatively by dissolving the steel in iodine-methanol and analyzing the residue for carbon. Severe sen-sitization was observed in Type 302 due to precipitation of only 0.003 pet carbon. Both cold work and the presence of delta ferrite caused a marked acceleration in rate of carbide precipitation. Carbide precipitation rates in 17-7 PH were stzulied for the austenite conditioning and also the aging heat treatment. CARBON and its compounds exercise a major influence on the properties of stainless steels and their response to thermal treatment. Sensitization in 18-8 type stainless steels has been the subject of numerous investigations throughout the years. Bain, Aborn, and utherford," and Binder, Brown, Frankss all studied the effects of heating austenitic stainless steels in the temperature range of 1000° to 1500°F. The primary purpose of most of these studies was the investigation of susceptibility to in-tergranular attack in acids due to these sensitizing heat treatments. Intergranular precipitation of carbides was always associated with intergranular attack but it was recognized2 that severe attack could occur with but minute quantities of precipitated carbide. Mahla and ielsen utilized the electron microscope to make a significant contribution in illustrating the appearance and method of growth of chromium carbides during sensitizing heat treatments. However, as they stated, their studies of residues could not be used to obtain a quantitative measurement of the amount of carbon which was actually precipitated. The aim of the present investigation was to devise a relatively fast, simple method for the quantitative measurement of carbides in stainless steel. EXPERIMENTAL WORK The initial investigations were made to determine the best means of separating carbides from the matrix. A number of dissolving media were tried using both chemical and electrolytic attack. Qualitative examination of the extracted residues by X-ray diffraction indicated that solution in iodine-methanol would furnish a good means of separation. Consequently, further work was pursued along this line. The method is quite simple. The sample in the form of millings or nibblings is dissolved in iodine-methanol solution at room temperature (6-g iodine, 25-ml methanol per g of sample). The insoluble residue containing the carbides is separated by suction filtration through an ultra-fine glass filter disc. This is a very fine filter medium that will retain particles as small as 0.1 to 0.2 p in diameter. After washing with methanol and drying, the filter disc and residue are placed in a conventional combustion carbon-tube furnace and the carbon determined gravimetrically. Using this technique it was found that reproducible insoluble carbon values were obtained. However, since such small amounts of insoluble carbon were obtained on Type 302 after sensitizatipn at 1250°F and 1500°F, the values were confirmed by a second method. In the second method the sample was dissolved with copper potassium chloride and filtered through a millipore paper. This treatment dissolves the matrix but leaves undissolved practically all of the carbon irrespective of how it is present in the steel. The amount of insoluble carbon present as chromium carbide is determined by calculation from the analysis of the residue for chromium and iron. The derivation of the formula used for this calculation is discussed later. The values obtained by the indirect copper-potassium-chloride method were in agreement with those obtained by the iodine-methano1 method. See Table I. It should be pointed out that the sensitivity of the direct combustion method is not too high when the amount of carbide present is small. This is due primarily to inherent blanks and to analytical errors such as weighing. For this reason it cannot be stated with any degree of certainty that there is a significant difference between values of 0.002 and 0.005 pct. Having confirmed that the iodine-methanol extraction gave a quantitative measurement of the precipitated carbides in Type 302, exploratory tests were conducted on Armco 17-7 PH stainless steel. Samples from commercial Heat 54807 were solution annealed at 2000°F, water quenched and heated at 1250" and 1500°F, and water quenched. The analysis of Heat 54808 is as follows:

Jan 1, 1962

By O. W. Simmons, L. W. Eastwood, C. M. Craighead

Eighty-four quaternary titanium-base alloys from the following systems were investigated: 1. Titanium-chromium-carbon-nitrogen. 2. Titanium-chromium-carbon with copper, vanadium, molybdenum, manganese, iron, or nickel. 3. Titanium-chromium-nitrogen with vanadium, molybdenum, or nickel. The tensile properties, minimum bend radii, hardnesses, response to heat treatment and aging treatment, and phase relationships of these alloys are reported. HIS paper is the third in a series describing J- titanium-base alloys. The first paper describingdescribed titanium binary alloys, and the second contained data on a large number of titanium ternary alloys. Eighty-four heats, comprising three groups of quaternary titanium-base alloys, are described in this paper as follows: 1. Titanium-chromium-carbon-nitrogen alloys. 2. Titanium-chromium-carbon alloys containing copper, vanadium, molybdenum, manganese, iron, or nickel. 3. Titanium-chromium-nitrogen alloys containing vanadium, molybdenum, or nickel. The methods of preparing, fabricating, and testing the alloys are described in the paper Binary Titanium Alloys. Quaternary Titanium - Chromiurn - Carbon -Nitrogen Alloys The properties of titanium-chromium-nitrogen and titanium-chromium-carbon ternary alloys have been described in the paper Ternary Alloys of Titanium. Because small amounts of nitrogen and carbon both appear to be somewhat beneficial, these additions have been investigated fairly completely in quaternary alloys containing chromium at the following levels of composition: 2.5, 3.5, and 5.0 pct chromium 0.25 and 0.5 pct carbon 0.1 and 0.2 pct nitrogen Mechanical Properties: The mechanical properties of these alloys are listed in table I. Fig. 1 and 2 show graphically a comparison of the tensile properties and hardnesses of quaternary alloys with those of the titanium-chromium binary alloys. Fig. 1 shows the comparison at 0.25 pct carbon, whereas fig. 2 shows the comparison at 0.5 pct carbon. It is concluded from these data that 0.25 pct carbon and 0.1 pct nitrogen are desirable additions to the 5 pct chromium-titanium base. It is further concluded that most of this improvement is produced by the nitrogen and that carbon is either harmless or mildly beneficial, at least up to 0.25 pct. Structure and Transformation Range: The effects of carbon and nitrogen on the transformation range of the titanium-chromium alloys were determined as part of the routine investigation. Additions of 0.25 pct carbon and 0.1 pct nitrogen had no appreciable effect upon the transformation range. However, the addition of 0.25 pct carbon and 0.2 pct nitrogen definitely raised the beta solvus temperature, as shown by fig. 3.

Jan 1, 1951

By Sidney Fischler

Large striation-free single crystals of bismuth have been grown from the melt by rapid freezing. Zone-refined bismuth, together with doping impurities if desired, is placed in a shallow flat-bottomed graphite boat and melted in air with a propane hand torch. The torch is then withdrawn in a manner which causes the melt to freeze direction-ally. Crystallization, which resuires only a few minutes, usually results in the formation of a single crystal even when a seed crystal is not used. Crys -tals of any desired orientation may be grown by using oriented seeds. Undoped crystals grown by this method have residual resistivity ratios greater than 200. THE growth of large single crystals of bismuth by either the Czochralski or horizontal zoning technique is not entirely satisfactory. Specifically, difficulties are encountered in producing single crystals of the required dimensions in all desired orientations, and striations caused by low-angle polysynthetic twins are frequently present in the crystals. In addition, both methods are time-consuming and require special apparatus of some complexity. A simpler method has now been developed for growing large striation-free bismuth single crystals of desired orientation in a short time. Fig. 1 shows a typical setup consisting of a rectangular graphite boat which contains zone-refined bismuth, a 1/8-in.-thick flat quartz plate which covers the entire inner bottom of the boat, and three additional quartz plates about 1/4 in. thick which are used to separate the bismuth from the graphite everywhere except at a small area at the left of the boat. The graphite boat is 1 in. high, and its sides and bottom are about 1/8 in. thick. The quartz plates should be smooth and clean. The graphite boat is heated from the right with a propane torch, as shown in Fig. 1, until the bismuth is completely melted. The melt has the shape of a triangle with a narrow neck at the apex farthest from the torch. The melt is frozen direc-tionally by gradually moving the torch toward the right, away from the boat. The bismuth in contact with the graphite, at the left end of the neck, freezes first. The freezing interface then moves down the neck into the main bulk of material, where it develops a convex shape ideal for the continuation of single-crystal growth. The interface continues to move through the melt until the entire bulk is solid. The entire procedure may be completed, in air, in a matter of minutes. The technique described almost always yields a single crystal whose basal plane is nearly perpendi,cular to the bottom of the graphite boat. In earlier experiments, in which the bottom of the melt was in direct contact with the graphite boat, single crystals were grown with basal planes parallel, perpendicular, or at some intermediate angle to the bottom of the boat. At times the orientation of the bulk of the material differed from the orientation of the material in the narrow neck. In these cases, a nucleation site initiated the growth of a differently oriented crystal, and the thermal conditions favored the new orientation over the initial one. The thermal conditions depend on a number of factors, including the heating technique, the placement, shape, and thickness of the quartz plates, the thickness of the walls and bottom of the graphite boat, and the quantity of bulk bismuth employed. All of these factors, plus the initial orientation and the presence and effectiveness of nucleation sites, will determine the orientation of the final large single-crystal slab. When a crystal of specific orientation is desired, an oriented section of a rapid-freeze crystal is shaped by spark cutting and grinding for use as a seed. To grow a doped crystal, the desired impurity is placed in the graphite boat together with the bismuth chunks and seed. Crystals doped with mercury, cadmium, lead, and selenium have been grown. The rate of freezing is so great that the distribution coefficient of any impurity approximates unity. On a gross scale, therefore, impurities should be more homogeneously distributed in rapid-freeze crystals than in Czochralski or zoned crystals. Because of the possibility of constitutional supercooling, however, it is quite possible that impurities are not homogeneously distributed on a microscopic scale in the rapid-freeze crystals. Generally the single crystal slabs which have been prepared are initially 5 to 7 mm thick. Thicker crystals may be obtained by using one of these slabs as a seed. The slab is placed in a graphite boat resting on a large aluminum block, either air- or

Jan 1, 1964

By Robert Parker

Measurements of the latent heat of transformation from a to 0 titanium were made using a pulse-heating technique. Rapid heating was accomplished electrically by discharging a low-inductance capci-tor hank through a thin-strip specimen. Both high-purity titanium and titanium containing 2700 ppm of 0 were studied. An electrical-resistivity technique and a plot of energy input us temperature rise were used to monitor the phase change. It was concluded that: 1) the latent heat of transformation is 800 i 48 cal per mole; 2) phase change occurred in less than 200 sec; and 3) the percentage phase change that occurs is linearly a function of the energy available for phase change. In addition, this experiment provides a simple and inexpensive method for determining specific heat of electrical conductors over large temperature ranges. The chief advantages of this technique over past pulse-heating experiments are: 1) static measurements of capacitor-bank voltage capacitance and specimen resistance can be made quickly and accurately to determine the energy supplied to the specimen; 2) the specimen is heated so rapidly (10 µ sec) that there is no appreciable heat loss during the heating. THE purpose of this paper is to illustrate the use of the capacitor bank as a tool for basic experimental investigations and to determine the rapid phase-change kinetics of titanium. The capacitor provides a number of unique properties: 1) rapid heating rates in the order of l09oC per sec can be obtained; heating rates as slow as l04oC per sec are available by merely adding a given amount of inductance to the circuit; 2) the energy input can be precisely determined and can be varied from several joules to millions of joules, depending on the bank capacitance and voltage; 3) the specimen can be heated uniformly through the thickness and length or non-uniformly, as desired; 4) the apparatus is simple and inexpensive and can be easily altered; 5) the form of the specimen can be wires, sheets, foils, rods, and cylinders; 6) there is great precision in the timing of events; 7) this technique may be readily used in conjunction with rapid cooling.

Jan 1, 1965

By Dol Duwez, R. H. Willens

A technique is described by which metastable alloy phases can be obtained by very rapid cooling from the liquid state. The results obtained so far have lead to; 1) extension of solubility limits beyond their equilibrium values, 2) new phases not found under equilibrium conditions, and 3) amorphous alloys. MANY new nonequilibrium phases have been found in alloys rapidly cooled from the liquid state. The technique used in these investigations essentially consisted of propelling a small molten globule of an alloy against the inside surface of a rotating copper cylinder. Since these original experiments were performed, the technique has been greatly improved. The purpose of this paper is to describe in detail, for the first time, a simple and efficient apparatus for producing foils of alloys which are rapidly cooled from the melt and to summarize some of the observations made on these alloys. I) FAST COOLING APPARATUS In the original fast cooling apparatus used by Duwez, Willens, and Klement,' the alloy was melted in a graphite container having the shape shown schematically in Fig. 1. A shock tube was attached to one end of the graphite part to supply the pressure for ejecting the molten globule. The globule was received on the inside of a rotating cylinder located about 0.5 in. from the exit of the graphite container. The globule trajectory made an angle of about 15 to 25 deg with the tangent to the cylinder at the point of

Jan 1, 1963

By A. E. Powers, J. F. Libsch

THE simultaneous influence of time and temperature upon the tempering process in steel is well known. Hollomon and Jaffe¹ have expressed the effect of time and temperature upon the progress of tempering by a single parameter, T (c + log t); where T is the absolute temperature; t, the tempering time; and c, a constant characteristic of the steel and treatment. It was found that when hardness data for plain carbon and low alloy steels were plotted vs. this parameter, a master tempering curve was obtained which satisfied all the data for a given steel and a given hardening treatment. Roberts, Grobe, and Moersch² successfully applied the tempering parameter to various tool steels, including high speed steel for tempering times ranging from 6 min to 100 hr. The influence of tempering time and temperature upon the physical properties of high speed steel has also been extensively investigated.'" Recently, the use of induction heating for very rapid tempering treatments has received attention. The authors% have shown that tempering treatments of only a few seconds duration may be successfully applied to plain carbon and low alloy steel. However, the influence of such short tempering treatments in high speed steel has not been explored. Since high speed steel is a highly alloyed steel exhibiting physical reactions during tempering that are often relatively slow and large in magnitude, it is an interesting steel for the study of tempering reactions. It was the purpose of the present investigation, therefore, to study the early progress of reactions leading to secondary hardness in high speed steel, type 6-5-4-2, and to determine if induction-tempering treatments of a few seconds duration are able to produce results equal to those of more conventional long-time tempering cycles. Experimental Details The chemical compositions for two different heats of W-Mo high speed steel used in this study are given in Table I. Table II summarizes the heat treatment given various groups of specimens cut 21/2 in. long and centerless ground to ¼ in. diam. Hardening treatments consisted of preheating to 1600°F (870°C) for V2 hr and then oil quenching from 2250°F (1230°C). Heat I was subjected to the high temperature in a salt bath for 3 min, while specimens of heat II were austenitized in a carbon-block muffle furnace for 10 min. Specimens of series E were immediately refrigerated to — 318°F (— 195°C) in liquid nitrogen following the hardening treatment to transform the major portion of the retained aus-tenite. Specimens of series F and G were subzero cooled to — 100°F (-75°C) in a dry ice-alcohol mixture. Induction tempering was performed with a Lepel high frequency converter operating at a frequency of approximately 300,000 cycles per sec. Heating cycles were controlled and recorded for each specimen using a chromel-alumel thermocouple, percussion welded to the surface of the specimen, and a modified high speed instrument with a timer previously used and described.' Temperature oscillation of the specimens, while being held at the tempering temperature, was ± 15°F (8°C). Tempering Treatments: After hardening, specimens were tempered between 900" and 1500°F (480" and 815°C) to study the progress of retained austenite decomposition and secondary hardening. The specimens tempered in salt for 1 hr (series C, G, and I) were considered to represent conventional treatment and are used for comparison with rapid tempering treatments. Several induction-tempering cycles were employed: 1—flash tempering, i.e., induction heating to the tempering temperature and immediately cooling (series A, E, and F); 2—induction heating for 60 sec at temperature (series B); and 3—double-induction tempering by two successive heating and cooling cycles (series D). All specimens tempered by induction were heated at an initial rate of 450°F (250°C) per sec except those of series E, which were heated at an initial rate of 1300 °F

Jan 1, 1952

By J. H. Wernick, S. Geller

A number of new AB, compounds, in which A is a rare earth or yttrium atom and B is Al, Mn, Fe, Co, or Ni, having the cubic MgCu, structure (Laves phase) are reported. In most of the compounds, the interatomic distances are normal; the lanthanide contraction is observed. However, in the compounds CeB2, B = Fe, Co, Ni, Ce appears to have a valence substantially greater than three; in CeAl2, Ce appears to be trivalent. These results are interpreted as vesulting from an energetically more favorable situation for the 4f electron of Ce to be lost to the 3d band of Fe, Co, nnd Ni. Evidence for the existence of homogeneitj! ranges in these compounds is presented. THE purpose of this paper is to report on the crys-tallographic investigations of a large number of new compounds, Table I, with the MgCuz structure (cubic Laves phase). It is a part of a continuing study of the magnetic properties and structures of binary systems involving the rare earths. In a previous paper' some of the magnetic properties of such systems were reported briefly. We have also reported2 a number of compounds of the CusCa type. The compounds were prepared by induction melting in Argon atmosphere of mixtures of stoichiomet-ric amounts of the constituent elements.* For most of the compounds fused alumina or silica crucibles were used as the containing vessel; graphite crucibles were used to prepare CeA12, NdAl2 and TbA12. X-ray powder photographs were taken with CrKa radiation and the use of Straumanis-type Norelco cameras of 114.6 mm diam. The compounds AB2 belong to space group Fd3m-Ol with 8A in a): (000; 01/2 '/2;3) + 000-; b % XA and 16B in dl: (000; oVa %; 3) + 5/87; % %; 7e % Vs\£>. Because the intensities of the reflections on the powder photographs of the different compounds are similar, relative intensities were computed only for DyMnz for CrKa radiation, Table 11. In the calculation of intensities, the Thomas and umeda3 atomic scattering factors, corrected for dispersion,4 were used. The agreement between observed and calculated intensities leaves little doubt that DyMnz has the MgCuz structure. Crystallographic data for the ABz compounds are given in Table JII. Twelve of these compounds have been reported by other authors. A comparison of the published results with our work is given in Table W.* ers of elements of lower purity, to peritectic reactions in some cases, and to the presence of homogeneity ranges in these alloys, so that melting losses would still result in single phase alloys. The possibility of existence of homogeneity ranges in these compounds was investigated systematically only for GdMnz. At room temperature, the solid solution range is between

Jan 1, 1961

By S. E. Haszko

SEVERAL new XGa2 intermetallic compounds, Table 1, where X is a rare earth, having the AlB, structure were prepared as part of a continuing study of the magnetic and structural properties of rare-earth alloys.' In this note, the crystallographic results of these compounds are presented. Chart, Eq. 121 for the data in Fig. 1 can be adequately represented by The corresponding standard free-energy change is AF; = -4.576 T log K = -20,900 + 10.20T [5I At low nitrogen concentrations the ratio of mole fraction to weight percentage is a constant so that the mole fraction X of nitrogen is given by X = 0.0392 [%NJ. At low values of X, the activity and the concearation approach each other, hence K, = 0.0392 K,, where K, is expressed in terms of the mole fraction of nitrogen. Hence Eqs. [4]. and [5] become

Jan 1, 1962

By R. E. Cech, J. H. Hollomon

KURDJUMOV and Maksimova reported experiments with manganese steels and high carbon steels' and with an Fe-Ni-Mn alloy' in which mar-tensite was formed isothermally over a range of temperatures. They found in some cases that mar-tensite formation could be suppressed by rapid quenching to liquid nitrogen temperature. From their microstructural observations of martensite formed isothermally, they concluded that the rate controlling step is nucleation rather than growth. Kulin and Cohen,3 in an attempt to reproduce these experiments, found that with a steel having the same composition as that reported by Kurd-jumov and Maksimova, the transformation to martensite was essentially complete above the temperature range of Kurdjumov and Maksimova's isotherms. The possible reasons for this disagreement were not considered. Recent papers by Das Gupta and Lement4 and Kulin and Speich5 report the formation of isothermal martensite in a high chromium steel and in an Fe-Cr-Ni alloy, but neither paper can be considered a verification of the original Kurdjumov and Maksimova results. Further, in neither case were the authors able to suppress the formation of martensite entirely. Because of the important bearing the Kurdjumov and Maksimova results have to an understanding of the mechanism of martensite reactions it was felt that an experimental investigation directly concerned with checking the validity of their results was in order. This paper describes the results obtained on the isothermal transformation over the temperature range from —79" to —196°C of an alloy of iron, nickel, and manganese. Experimental Apparatus A 15 lb heat of an alloy containing 73.3 pct Fe, 23.0 pct Ni, and 3.7 pct Mn was melted by induction and cast under argon. The ingot was forged to 1-in. bar and a portion rolled to 1/16x1 1/2-in. strip. This strip was pack-homogenized 300 hr at 1100" in a helium-filled sealed iron tube. The composition after homogenization was 73.2 pct Fe, 22.94 pct Ni, 3.73 pct Mn, 0.05 pct C, and 0.015 pct N. The strips were cut to 1/2-in. width for dilatometer and metal-lographic specimens. Only the center portion of the 11/2-in. strip was used in the present investigation. The dilatometer employed was similar in design to one described by Flinn, Cook, and Fellows." A concentric fused auartz rod and tube assembly with hooks for holding the specimen was mounted so as to transmit the specimen dilation to a 1/10,000 in., 1/10 in. travel dial gage. The dilatometer proper was mounted by means of extension arms to a counterweighted sliding member on a vertical standard. This method of mounting permitted rapid transfer of the dilatometer from the austenitizing furnace to the quenching bath and low temperature chamber. A small electrical vibrator on the dilatometer kept frictional effects of the quartz members at a minimum. The austenitizing unit was a vertical, molybdenum-wound, hydrogen atmosphere furnace maintained at a constant temperature ±3°C by means of constant power input. A 12-in. stainless steel jacketed copper liner having 1/2-in wall thickness acted to equalize the temperature in the hot zone of the furnace. This liner, closed at the bottom end and open at the top to permit entrance of the dilatometer and specimen, was kept filled with dry nitrogen gas. A chromel-alumel thermocouple was placed inside the tube to determine the temperature. The 4-in. dilatometer specimens in the chamber varied less than 1/2° across the specimen length except for a 1 1/20 drop at the end nearest the open end of the furnace. The low temperature isothermal holding bath was a double Dewar arrangement similar to one described by Turnbull7. The outer bath was filled with a refrigerant at a temperature lower than the desired holding temperature. The inner bath was filled with Freon "11" or "12" or a mixture of both, depending upon the holding temperature. This inner bath which tended to be cooled by the outer bath was kept at a constant temperature by introducing a small amount of heat with a manually controlled electric heater. Stirring was accomplished by bubbling dry air through the bath. A Leeds and North-rup type K potentiometer was used to measure the inner bath temperature as indicated by a five element copper-constantan thermopile. The bath temperature was maintained within ±0.2°C of the desired temperature by occasionally adjusting the heater current so as to keep the Leeds and Northrup galvanometer at zero deflection with a constant setting of the potentiometer. Isothermal tests were usually continued for 300 to 400 min and another reading made at approximately 1000 min if the bath, unattended overnight, had not deviated in temperature more than 5°C. Transformation curves are drawn dashed (Fig. 1) through the time region where temperature was not controlled precisely. Experimental Procedure Dilatometer specimens of 1/2x1/16-in. strip were cut to 41/2-in. length and holes were drilled for the quartz hooks with proper spacing to give a 4-in. measured length. A thermocouple consisting of 0.012-in. diameter chrome1 and alumel wires was spot welded to the specimen and threaded between the dilatometer rods to binding posts near the dial

Jan 1, 1954

By R. F. Mehl, R. F. Bunshah

A fast amplifier technique has been developed for the measurement of the rate of propagation of martensite in an Fe-29.5 pct Ni alloy. The time of formation of one plate of martensite is 3x10 sec and the rate of propagation is 3300 ft per sec approximately. IT has been known for some time that the plate-like structural unit of martensite forms from austenite with great rapidity. Wiester1 and Hane-mann, Hofmann, and Wiester took motion-pictures of the transformation as it occurs in a 1.65 pct C steel; they demonstrated that a single plate formed fully in the time interval between successive frames, viz., 1/20 sec, thus setting an upper limit. Forster and Scheil,3 using an Fe-Ni alloy with 29 pct Ni, recorded the sonic characteristics of the process electrically, upon an oscillograph, setting the upper time limit at 0.002 sec. Forster and Scheil,~ measuring the change in electrical resistance in the same alloy upon a cardiograph, set a limit of 0.02 sec. Forster and Scheil5 later, employing the same alloy, improved their technique, reporting an upper limit of 7.10 sec. In studying signals of such short duration, it is an important question whether the frequency response of the electrical system used is high enough compared to frequency of the pulse measured, or, put differently, whether the system is able to reproduce without distortion the signal arising, in this case, from the formation of a single martensite plate. Forster and Scheil (referring only to their last paper) obtained signals of a frequency of 30 kilocycles (hereinafter kc); this was about the frequency response of the equipment used; thus, if the signal had a frequency higher than 30 kc, it would still appear as a signal of frequency 30 kc. All of these results thus provided upper limits only. Recent developments in electronics have made available equipment with very high frequency response, very high sweep-speeds, high gain, etc. The electrical characteristics of such equipment, used in the present study, are given in Table I. Such equipment offers obvious attraction in the study of the rate of propagation of a martensite structural unit—and perhaps of other structural alterations proceeding at a very high rate. This paper reports an attempt to develop a technique employing such equipment to measure the time of propagation of a martensite structural unit and the variation of this with temperature, with the mode of formation—athermal and isothermal—in both polycrystalline and single-crystal samples; and from such measurements to obtain the rate of propagation. As will be seen, the results obtained are useful theoretically. Materials All data presented here are for an Fe-Ni alloy of the following analysis: 29.5 pct Ni, 0.027 pct C, 0.135 pct Mn, 0.094 pct Si, balance Fe. There were several reasons for choosing this alloy: 1—it is substantially the one used by previous investigators; 2—it exhibits both the athermala and the isothermal' mode of formation of martensite, both studied in detail by Machlin and Cohen; 3—the subzero temperatures of transformation in this alloy are experimentally very convenient; 4—it exhibits the "burst phenomenon";" 5—the change in electrical resistance upon the formation of martensite, a decrease, is great, approximately 50 pct.' The polycrystalline specimens were in the form of wires of 0.025 in. diameter; the single crystals were 1x1/4x1/4 in. Experimental Methods Electrical Apparatus: Fig. 1 is a schematic drawing of the electrical circuit used. The principle used in these measurements is the same as that used by Forster and Scheil." A small direct current, about 1 to 2 amp, is passed through the sample. When a martensite plate forms, the resistance of the sample changes and a high frequency signal is generated. This signal is picked up by the probes attached to the sample, fed into the bank of amplifiers and thence to the vertical deflection plates of a cathode-ray oscilloscope. The signal itself triggers the oscilloscope trace which flashes across the tube face and is photographed by means of a 35 mm movie camera at the end of a light-tight hood. The camera has no shutter. As soon as the signal flashes

Jan 1, 1954

By P. G. Shewmon, F. N. Rhines

THE determination of the self-diffusion coefficient of magnesium has been made possible recently by discovery1-1 of a radioactive isotope, Mg28 having a half-life of 21.3 hr,1 and subject to manufacture in useful quantity. In the present research this material was condensed from the vapor phase upon a surface of high purity magnesium. The progress of diffusion of the tracer atoms into polycrystalline magnesium was followed by machining layers and measuring the change in the intensity of radiation as a function of the distance of each layer from the surface. The self-diffusion coefficient was found to be 2.1 X 10-8 sq cm per sec at 627°C, 3.6 X 10-9 sq cm per sec at 551°C, and 4.4 X 10-10sq cm per sec at 468°C; the activation energy is about 32,000 cal per mol. Experimental Procedure Since there was no other published measurement of a diffusion velocity in any magnesium-base material, is was necessary to employ a number of new experimental techniques. The short half-life of Mg28 made it necessary to complete the entire experimental procedure within three or four days. This meant that the work had to be done where a cyclotron was readily accessible and that all operations, prior to the diffusion heat treatment, had to be so designed as to minimize their time requirements. Unusual problems were imposed also by the chemical reactivity of magnesium, its high vapor pressure, and the fact that no satisfactory method for elec-trodepositing magnesium on magnesium is presently available. Finally, the machining and handling of the easily air-borne radioactive-magnesium chips involved certain health hazards, resulting in the need for further experimental restrictions. Preparation of Mg28 The Mg28 was produced in the Carnegie Institute of Technology syncrocyclotron by the neutron spallation of chlorine.5 his involved bombarding a 2 gram crystal of high purity NaCl with a beam of 350 mev protons for a period of 2 hr, after which the crystal was dissolved in warm water and the Mg28 was concentrated and purified by chemical means (see Appendix). About 50 microcuries of Mg28 thus were obtained in the form of magnesium oxinate (8 hydroxyquin-olatc?), which was ignited in air to produce MgO. This in turn was reduced to magnesium metal vapor, by the method of Russell, Taylor, and Cooper," in the vacuum apparatus shown schematically in Fig. 1. Here the essential part is a tantalum ribbon, slightly dished to receive the MgO. The ribbon, pre- viously outgassed at high temperature, is heated to about 1700°C by passing an electric current through it, whereupon tantalum oxide is formed, magnesium vapor is released almost instantaneously, and condensed partly upon the diffusion sample. Diffusion-Sample Preparation: Hot-extruded magnesium rod, 21/32 in. round was used in making the diffusion specimens. The magnesium analyzed as follows: 0.004 pct Al, 0.027 pct Fe, 0.040 pct Mn, 0.0004 pct Cu, 0.0002 pct Ni, and less than 0.01 pct Ca, 0.0004 pct Pb, 0.0011 pct Si, 0.001 pct Sn, and 0.001 pct Zn. A brief study of the crystal texture of this material revealed a sharp fiber texture with the (001) plane roughly parallel to the extrusion axis. Cylindrical samples 1/2 in. long by 5/8 in. were machined from this rod, the end faces dressed on 3/0 emery, and lightly etched with 20 pct HC1 in water. These samples then were annealed for at least twice the intended time of diffusion, at the intended diffusion temperature, in order to stabilize the grain structure at about 1 mm average diameter. The annealing treatments were conducted in argon in the same apparatus and in the same manner as the subsequent diffusion treatments, which will be described presently. Thus, a strain-free plane surface was produced, but there remained a layer of MgO which had largely to be removed before the layer of Mg28 was deposited. Most of this layer was taken off by two light passes over 3/0 emery paper. The balance of the oxide and a thin layer of metal were then removed by etching 5 to 10 min in 4 pct nital (4 pct HNO3 and 96 pct ethyl alcohol) made with absolute alcohol. There followed immediately three quick rinses in: 1-49 1/2 pct methanol, 49 1/2 pct acetone, and 1 pct formic acid, 2-50 pct methanol and 50 pct acetone, and 3-pure benzene. This procedure is essentially that of Sturkey.7 The resulting surface, which was of almost elec-tropolished brightness, remained plane and was free of cold work. It could be kept clean by storing under benzene, or in a desiccator; short exposure

Jan 1, 1955

By H. S. Cannon, F. N. Rhines

The application of a static load to a copper powder compact during sintering at an elevated temperature accelerates the rate of sintering in such a way that a given load induces the same proportional increase in rate for all times of sintering. It is shown that sintering under a load is like creep under a fixed load in that the stress required to accomplish a given degree of densification is proportional to the logarithm of the sintering time. VIRTUALLY all of the theories of sintering that have been put forward within recent years have contained the assumption that the chief driving force of the process is the surface tension resident in the exposed surfaces and internal pores of the compact, or powder mass. An externally applied load might be expected to provide a somewhat equivalent driving force for sintering. It is known that the application of a compressive load, during sintering, hastens densification, but it is by no means clear whether sintering under the influence of surface tension alone and sintering under an applied load are fundamentally similar processes. The present study was undertaken in an effort to obtain an answer to this question and with the hope that the answer may contribute to an understanding of the mechanism of sintering. It has been found that there is a close relationship between sintering and creep processes. Copper powder compacts were sintered in hydrogen at 1000 °C under dead loads of 0 to 165 psi and the progress of sintering was observed by means of density measurements. Using a commercial reduced oxide copper powder, see Table I, cylindrical compacts Yz in. x x in. high were made at a pressure of 12,500 psi. For sintering, these were placed in a cylindrical graphite container into which they fitted snugly. A compressive axial load was applied to the compact through a graphite rod of similar diameter, which rested upon the upper end face of the compact; iron weights, in suitable amount, were placed upon the graphite rod to provide the load. This assembly was encased in a vertical silica tube within a resistance-type furnace capable of maintaining a substantially constant temperature within the working zone. After various predetermined times at temperature, the samples were removed from the furnace and their density was measured by conventional means. The heating-up and cooling-down times have been included in the computed sintering time, using the assumption that the sintering rate doubles for each increase of 10 °C in temperature. The experimental results are presented in Table 11. Upon metallographic examination of the specimens, it was found that there was no apparent change in the shape of the pores as a result of loading, i.e, no flattening, Fig. 1. Substantial contraction away from the container walls was observed in all samples sintered with loads of 10.5 psi and less.

Jan 1, 1952

By H. S. Cannon, F. N. Rhines

IT. H. Hausner( Sylvunia Electric Products Inc., Bay-side, N. Y.)—The results reported by the authors are interesting because they contribute some information on the principles of sintering and also because of a certain practical value for the manufacture of powder . metallurgical parts in production. Pressed powder compacts are frequently sintered in a boat in several layers so that the lower layers are actually sintered under the weight of the upper layers and, therefore, are frequently deformed. The authors gave information on this deformation process by determining the effect of the load on the density. The experimental setup of the authors, however, has a disadvantage—that the sintering occurred in a cylindrical container so that the deformation was stopped in the direction of the diameter. The authors should have rather called this a study of hot-pressing than of sintering. I want to call the attention of the authors and others interested in sintering under load to the experimental work described by Dawihl and Rix.' These authors sintered cylindrical compacts under a lengthwise ten-sional load of 0 to 25 g and found that the shrinkage in length and diameter changes considerably with the load, but that the volume shrinkage is affected very little by the load. H. S. Cannon (authors' reply)—We wish to thank Dr. Hausner for calling attention to the experimental work of Dawihl and Rix. The fact that they found volume shrinkage little affected by varying tensional loads during sintering seems to support the contention that the forces due to loading are equivalent to and additive to the forces which cause densification. This contention is the basis of our comparison of sintering and creep.

Jan 1, 1952

By R. W. Heckel, H. W. Paxton

The growth of grain boundary films and Widmanstiitten plates of proeutectoid cementite in hypereutectoid steels ulas considered quantitatively in three plain carbon steels and one "pure" iron-carbon alloy. Growth was found to take place at a slower rate than would be predicted on the basis of a model invoking rate control by long range diffusion of carbon. The films in the pure alloy grew faster than those in the plain carbon steels. A possible explanation is suggested. In 1904, Nernsst' and Brunner2 postulated that ionic crystals growing in supersaturated aqueous solutions grew at rates determined by diffusion through the solution to the growing crystal interface. The gradient down which the diffusion took place was assumed to be fixed by the supersaturated bulk solution concentration, the concentration dt the crystal surface and some arbitrary distance. The flux of material to the crystal was described as Flux where D is the diffusion coefficient, A the crystal surface area, Ci the bulk solution concentration, Co the saturated solution concentration, and 6 the thickness of the diffusion zone. It is obvious that the expression is an adaptation of F+ick's3 first law of diffusion (steady state). This was the initial formulation of what will be defined here as "diffusion controlled growth," i.e., growth limited by the long range diffusion of a component either to or from the advancing interface assuming both phases at the interface to be in equilibrium. This idea of equilibrium interface concentration implies an infinite reaction rate at the interface. Volmer, Zener,' and 11schner6 have derived expressions for the diffusion controlled growth of spherical precipitates in metallic systems. These analyses agree that the growth in a given direction should be parabolic with time. Brown and Hawkes7 observed that graphite, formed both from the decomposition of cementite and from the decomposition of eutectoid austenite in cast irons, grew according to a parabolic relationship and, therefore, concluded that the laws of diffusion controlled growth, as defined above, were obeyed. They, among others, failed to consider, however, that parabolic growth is a necessary, but not sufficient, condition for diffusion controlled growth. Mehl and Dube"' have pointed out that equilibrium may not be realized because the concentration at the interface cannot change instantaneously from that of the bulk to the saturation value. Furthermore, they propose that an equilibrium solute concentration in the parent phase adjacent to the precipitate would remove the thermodynamic activity gradient across the interface and, thus, stop the growth process. Frisch" has also emphasized the importance of interface concentrations that vary throughout the growth of a new phase. The variation in activity across an interface, as discussed by Onol' and Cahn and Hilliard,12 is of fundamental importance to this question. Ham13 considered the data of Turnbull14 on the growth of barium sulfate in aqueous solutions. He concluded that the data could be analyzed in terms of a radiation-type boundary condition. Flux across interface = K (c, - C,) [21 where C, is the concentration in the solution at the interface, Co the saturated solution concentration, and K the proportionality constant. The edgewise growth of Widmanstlitten plates, frequently referred to as the rate of lengthening, was first considered by Zener.'' Basing a mathematical analysis on steady-state (constant velocity) growth he concluded that where G is the growth rate, D the solute diffusivity, Ci the concentration in the bulk, Co the equilibrium interface concentration, C, the concentration in the precipitate, a a constant, Y the radius of curvature of the edge of the plate, and rC the critical radius, i.e., the radius which would increase the activity of solute at the interface at equilibrium under the influence of capillarity to that of the bulk. Zener1'

Jan 1, 1961

By Elmars Ence, Harold Margolin

The Ti-Fe and Ti-Fe-0 systems were re-examined because of the controversy regarding the existence of Ti2Fe, and to consider all available data points to the existence of Ti,Fe. The Ti-Fe-0 system contains two ternary compounds: E, corresponding to Ti2Fe; and y. SEVERAL authors have investigated the constitution of titanium-rich alloys of the Ti-Fe system. The most extensive study of titanium-rich portions of the Ti-Fe system was by Van Thyne, Kessler, and Hansen,' who investigated this system up to 50 pet Fe and used the highest purity materials available for their alloy preparation (iodide titanium was used for alloys up to 20 pet Fe). According to Van Thyne et al., phase relationships in the region of investigation are governed by phases a,ß, and TiFe. No evidence of an intermetallic compound Ti,Fe was found. although earlier. works by Laves and Wallbaum,' and Duwez and Taylor' reported its existence. Rostoker's study on the occurrence of TilX phases" confirmed the findings of Van Thyne et al. Rostoker proposed that the compound Ti Fe found by Duwez does not exist but is actually a ternary compound, Ti1Fe3O, originated by inadvertent oxygen contamination. The partial isothermal section of the ternary Ti-Fe-0 system as reported by Rostoker' seems to confirm this explanation. In the course of phase diagram work conducted at New York University, however, certain irregularities were observed to be associated with ternary systems of the type Ti-Fe-X. The ternary phase 6, observed in the systems Ti-Fe-Mo' and Ti-Fe-V." was found to be structurally identical to the com-pound Ti2Fe as reported by Duwez and Taylor, and to Rostoker's compound Ti1,Fe2O. Since the amount of oxygen possibly present in the Ti-Fe-Mo and Ti-Fe-V systems could not account for the observed amounts of the compound, it appears that the d phase in these systems is not Ti,Fe,O. On the other hand, considering the amounts of the 6 phase, particularly in the Ti-Fe-Mo system, the location of the phase was uncertain. It was felt that the introduction of the Ti2Fe phase would alleviate some of the inconsistencies of the two systems. Because of the uncertainties relating to the phase Ti2Fe, or Ti1Fe2O, it appeared worthwhile to re-examine the Ti-Fe and Ti-Fe-O systems in the vicinity of these compounds with highly sensitive metallographic techniques and with X-ray methods. Experimental Procedure Alloy Preparation—For the study of the binary system four alloys were prepared, containing 26.9 (30 wt pct), 34.0 (37.5 wt pet), 41.2 (45 wt pet), and 56.3 (60 wt pet) atomic pet Fe, and for the ternary Ti-Fe-O system 18 alloys were prepared in the composition range of 1.6 to 16.9 atomic pet 0 and 24.3 to 55.7 atomic pet Fe. The materials used for the Ti-Fe alloys were iodide titanium (99.98 pet Ti, 0.002 pet 0) and Ferro-vac-E iron (99.95 pet Fe, 0.0052 to 0.0072 pet 0). For the Ti-Fe-O study the materials used were sponge titanium containing less than 0.06 pet 0, Ferrovnc-E iron, and Baker's analyzed TiO2. The nominal weight and atomic percentages of the ternary alloys prepared are shown in Table I. Melting—Charges of 10 to 15 g were melted in a nonconsumable arc furnace in argon atmosphere according to the technique described elsewhere.'" Since there was practically no weight loss through the various stages of melting and oxygen losses have not been observed previously, it appeared justifiable to use nominal composition for interpretation of data. Iodide titanium control buttons, of the same mass as the Ti-Fe charges, and melted under the same conditions, were used to check hardness pickup during melting. A hardness increase of 2 to 3 Vhn was found. Heat Treatment—The specimens were annealed in quartz capsules under argon. To avoid contact between the specimen and capsule material all specimens were wrapped in titanium sheet. The annealing times are given in Table 11. The attain-

Jan 1, 1957