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By R. G. Hudson, L. Yang

In electrochemical separation of metals, it is necessary to control the potential applied between the electrodes so that only the desired electrode reactions can occur. A knowledge of the minimum potential needed for a given electrode reaction to take place continuously is therefore pertinent to the success of the process. In molten chloride systems at high temperatures, activation polarization is believed to be small,1,2 so the minimum potential needed for the electrolytic decomposition of a molten metal chloride MC1n (n = valence of the metal ion) is therefore essentially the equilibrium electrode potential E of the metal-chlorine galvanic cell in the solution. If the decomposition products are pure M and chlorine at a pressure of p atm, E*Eo-rt/nf in [i] Since EO (the standard electrode potential of MCln) is a constant at a given temperature and R (the gas constant), and F (the Faraday's constant) are all constants, the variation of E with the electrolytic bath conditions is largely determined by the effect of the latter on the activity of the metal chloride amcln (standard state = pure MC1,). Measurement of E and study of how amCin varies with temperature and concentration, the nature of the solvent, and that of the coexisting solutes are therefore pertinent in the selection of optimum operating conditions for the electrochemical separation of metals. Because of its low melting point, high decomposition potentials of its components, and the low solubility of heavier metals in it, the LiC1-KC1 eutectic melt is a useful solvent and has been used in the electrodeposition of Mo,3Ce,4 and u5 metals. Senderoff and Brenner3 measured the potentials of Zn/Zn++, Fe/Fe++, Cu/Cu+, Mo/Mo+++, and Ag/Ag+ against an Ag/AgCl (pure) reference electrode at 600°C and a concentration of 4.1 mole pct of each chloride and came to the conclusion that complex ion formation occurred in some of the melts. Laitinen and LiUe gave an electromotive force series at 450°C for a number of {Metal/Metal ions (unit activity)) electrodes in LiC1-KC1 eutectic melt, the {pt/pt++ (unit activity)) electrode being chosen as the reference. Walker and Danly7 studied the ther-modynamic properties of NiC12 in LiC1-KC1 eutectic melt over a wide range of temperatures and concentrations by measuring the electrode potential between a nickel electrode and a chlorine electrode in the melt. Since potential measurements made by using the chlorine reference electrode are of more direct significance both in the study of electrochemical separation and in the evaluation of the thermodynamic properties of the melt, it is therefore decided to measure the electrode potentials of cells of the type {M/MCin (mole fraction = N) in LiC1-KC1 eutectic melt/Cl2 over a range of N values and temperature for a number of metal chlorides. The results should indicate the order by which the various metals plate out on the cathode and how this order is affected by temperature and the concentration of the metal chloride in the solu-

Jan 1, 1960

By C. A. Siebert, D. J. Girardi

PRACTICALLY every metallurgical process in-volves, at some stage or another, the contact of a metal with a gas. Because of this, gas-metal re-actions are of great practical importance and have been subjected to numerous investigations. It has long been recognized that the presence of cuprous oxide in liquid copper has an important effect on its casting characteristics and .physical properties. This has led to the desirability of determining the quantitative relationships between cuprous oxide in the liquid copper and the usual gases present during processing. Such data are not only of practical importance but also of a theoretical interest in that they provide a means of applying the principles of physical chemistry to problems in applied metallurgy. The object of this investigation was to determine the quantitative relationship at equilibrium between the gases carbon monoxide and carbon dioxide and the cuprous oxide content of liquid copper. The reaction and its equilibrium constant may be expressed as follows: 2 Cu (liq) + CO2 (g) = Cu2O + CO (g) acu2o X aco/a2 cu X aco2 Here the underlined symbol Cu2O designates cuprous oxide dissolved in liquid copper and (a), the activity of the various constituents. The activities of carbon monoxide and carbon dioxide are assumed to be directly proportional to their re- spective partial pressures and the activities of cuprous oxide and copper are taken to be proportional to their respective mol fractions. The equilibrium constant may then be expressed as follows: K2 = Ncu2o X Pco [Ncu]2 X Pco2 where N = mol fraction The solubility of carbon monoxide and carbon dioxide in liquid copper is a rather controversial subject.'-' A review of the literature is given by 0. W. Ellis.' It appears that carbon monoxide and carbon dioxide are both only very slightly, if at all, soluble in liquid copper. Volskii and Slobodskoi5 determined the equilibrium between carbon monoxide, carbon dioxide, and liquid copper at 1100°C. Their results show a higher cuprous oxide content for a given CO/CO2 ratio than found in the present work. The system hydrogen, water vapor, and liquid copper was investigated by Allen and Hewitt8 over the temperature range of 1090" to 1350°C. Equilibrium values for the system carbon monoxide, carbon dioxide, and liquid copper calculated from these data agree reasonably well with the present work. Materials and Procedure The copper used was oxygen free, designated as OF and analyzed as 99.98 pct copper with no appreciable amount of any one impurity based upon spectrographic determinations. The carbon dioxide was cylinder gas averaging about 99.95 pct CO2 with air as the major impurity. The carbon monoxide was produced by passing carbon dioxide over heated charcoal. Alundum combustion boats were used as containers for the liquid copper. No appreciable reaction took place between the copper and the alumina. This was evident from the fact that the copper sample after a run did not ordinarily adhere to the bottom of the boat and that the

Jan 1, 1951

By Benny J. Nelson

AS a part of the fundamental research program of Aluminum Research Laboratories, some data were obtained on the ternary system Mg-Al-Mn. As very little information on the magnesium corner of this diagram has heretofore been published, it seems desirable to make available the values found for the liquidus and solidus surfaces of this system. Procedure The settling procedure was used for the determination of the liquidus compositions. Metallo-graphic examination of quenched samples, and stress-rupture upon incipient melting, were used for the solidus determinations. The settling procedure has been described in a previous paper.' Briefly, this method involved saturating the alloy with manganese at a temperature substantially above that at which the sohbility was to be determined, then cooling the melt to the latter temperature, and holding it at that temperature for a substantial period of time. Samples for analysis .were carefully ladled from the upper portion of the melt at hourly intervals during the holding period. After the ladling of each sample, the melt was stirred to redistribute some of the manganese that had already settled, because it appeared that when the latter particles of manganese again settled, they aided in carrying down more of the manganese and thus hastened the attainment of equilibrium. The melts were prepared and held in a No. 8 Tercod crucible holding approximately 4 lb of metal. The manganese was added either in the form of a prealloyed ingot (Dow M) containing about 1.5 pct Mn or by the use of a flux (Dow 250) containing manganese chloride. In calculating the flux additions, it was assumed that the manganese introduced would be equal to 22 pct of the total weight of the flux. Temperatures were measured with an iron-constantan thermocouple enclosed in a seamless steel tube, the lower end of which was welded shut. This protection tube also served as a stirring rod. The samples ladled from the upper portions of the melts at the various intervals were analyzed for aluminum, manganese, and iron. When making the alloys which were to be used for the determination of the solidus, 2½ in. diam tilt mold ingots were cast, scalped to 2.0 in. in diam, and extruded into ? in. diam wire. The principal impurities in the melts for this investigation were iron and silicon; their total not exceeding 0.03 pct. Portions of the wire, approximately 2 in. in length, were enclosed in stainless steel capsules for protection from the atmosphere. Bundles of these capsules, with a dummy capsule containing an iron-constantan thermocouple, were heated inside a large steel block (acting as a heat reservoir) in a closed circulating-air type electric furnace. At ap- propriate times, the capsules were removed and quenched in water. The wires were examined metallographically to determine the temperature of initial melting. Short times at temperature were used at the beginning for wire specimens of all alloys to obtain quickly the approximate temperatures at which melting could be first observed. When approximate solidus temperatures had thus been determined, equilibrium heating was attempted. This equilibrium heating consisted of an 8 or 16 hr period at a temperature, about 50 °F below the lowest temperature at which melting occurred when short heating cycles were used, followed by further heating for 1 hr periods at consecutive 10" higher temperatures. The theory for the method of stress-rupture at incipient melting has been well covereda and its limitations are recognized. Thus, if the interfacial tensions are such that the first minute quantity of liquid is "bunched up" at the grain boundary junctions instead of spreading out along the grain boundaries,³ temperatures higher than the solidus are required before melting will be manifested by rupture of the specimen. This point will be elaborated later. Specimens of the wires with a reduced section (approximately 1/16 in. diam) were suspended vertically in a tubular furnace. The setup used is shown in Fig. 1. The clamp holding the specimen was made from alumel thermocouple wire and the thermocouple was thus completed across the specimen by attaching a chrome1 wire to its lower end. Temperatures were read from a Speedomax recorder used in conjunction with a calibrated thermocouple. The small weight attached to the specimen and a vibrator attached to the furnace tube, to aid in distributing the molten constituent along the grain boundaries, were used to bring about rupture at a temperature closely approximating the solidus. The specimens were heated at a rate of about 5°C per min. The rupture of the specimens was indicated both by sound and by the action of the recorder. An argon atmosphere containing a small amount of SO² was used for protection of the specimen. The assembly was taken out of the furnace immediately following rupture and the specimen removed. Some of the broken specimens were examined metallographically and will be referred to later. Results and Discussion Fig. 2 shows a set of typical time-composition curves for liquid samples of the Mg-Al-Mn alloys used for the settling tests. The data as presented

Jan 1, 1952

By R. M. Hudson

Equilibrium constants have been determined for the jeaction C (in steel) + H2O = CO + H2 as a function of carbon content (0.013 to 0.74 ujt pct) and temperature (1200° to 1800°F) by using a flow system in which the water-vapor content of various gas mixtures containing hydrogen, carbon monoxide, and nitrogen was systematically controlled. The equilibrium partial-pressure ratio K1 = PcoPH2/PH2O was obtained for two low-carbon steels and four Fe-C binaries from a series of runs made at each temperature. By comparison of Ki with the analogous ratio for the reaction C (graphite) + H2O = CO + H2, the activity of carbon in steel (ac) was comDuted for each temperature and carbon content studied. From these data values of the equilibrium constants for two other carburization-decarburiza-tion reactions of importance were estimated. The chemical compositions of the two steels and four Fe-C binaries used are given in Table I. Specimens of Steel E were cut from a 0.035-in.-thick commercially box-annealed and temper-rolled lot of a low-carbon low-metalloid steel. Steel F was a commercially hot-rolled rimmed steel that was cold-reduced at the Laboratory to a 0.059-in. thickness. Fe-C binaries A, B, C, and D were vacuum-melted and poured, hot-rolled, pickled, and cold-reduced at the Laboratory to furnish specimens 0.035 in. thick. Sheet specimens were cut from all lots to approximately 3/4 by 3 in. by gage and were vapor-degreased with trichlo-roethylene and weighed before heat treatment. Heat treatments were carried out in a hinged-top Hevi-Duty electric furnace with a Vycor-glass tube. When the furnace was at the desired

Jan 1, 1965

By Larry Kaufman

A useful method for estimating the entropy of NaCl type compounds has been developed by combining the Debye theory of specific heat with the Lindemann formula. This method pemnits calculation of the entropy from a knowledge of the atomic weight, atomic volume, and melting point of the compound. Comparison of the calculated and observed entropies of 35 NaCl type compounds over a wide temperature range yields good agreement thus establishing the usefulness of the method. Application of the method to other crystal types is also discussed. KUBASCHEWSKI~ has suggested a useful method of estimating the enthalpy of formation of inter-metallic compounds from a knowledge of the crystal structure, lattice parameter, and heats of sublimation of the component elements. This method is quite valuable in predicting the tendency for compound formation in a given system and is of particular value in considering systems for which no experimental data are available. In order to supplement the enthalpy calculation afforded by Kubschewski's method, it would be of value to be able to estimate the entropy of compounds so that the entropy and free energy of formation could be evaluated. Although a considerable body of such data determined by calorimetric and equilibration techniques exists at present, the difficulty of making such measurements (particularly at high temperatures) and the ever increasing number of new compounds which are of interest make the desirability of such a computational method clear. This paper deals with a method for calculating the entropy of a compound from a knowledge of the atomic weight, atomic volume, and melting point. Since these data are generally easier to determine and therefore more readily available than thermo-dynamic data, they should provide an additional tool in estimating the stability and reactivity of compounds at high temperatures. I) THE ENTROPY OF PURE METALS As a first approximation, the specific heat of a pure metal can be represented by a sum of terms where C,($/tI is the vibrational specific heat calculated on the basis of the Debye theory, lo4 TC, (0/T) is an approximation to the Cp - C, correction. T is the electronic specific heat, and Cp (p) is the magnetic specific heat. Although the Debye theory does not provide a complete description of the vibrational specific heat, i.e., the Debye temperature 0 is not a constant, it does provide a good approximation when used properly. Blackman6'7 has written several reviews on this subject discussing the different values of 0 determined by various methods (specific heat, X-rays, elastic constants, electrical resistivity, and so forth), as well as the observed temperature dependence of 0 with T. In the present case the 0 in Eq. [I] naturally refers to the specific heat 0. The variation of Debye 0 (as determined from specific heat measurements) with temperature is observed to exhibit its most radical behavior at very low temperatures. At higher temperatures the 8 values are observed to become more nearly constant. Blackman6'7 has presented the variation of 0 with temperature for several metals. His tabulation is reproduced in Table I. Reference to Table I shows that above a certain temperature level, 0 remains fairly constant. The value of 0 which is most useful in computing the specific heat, and hence the entropy, is the value of 0 in the range 0/6 « T « 0. In 1910, Lindemann' suggested a relation for estimating Debye 0 values. According to this formula, where V is the atomic volume, T is the melting point in 'K, M is the atomic weight, and CL is a constant. Mott and Jones9 present this formula and state that there is at present no satisfactory theoretical explanation, but that it does have an empirical basis. By using the valuesof 0 in Table I along with known1' values of M, T, and the molar volume VM, where the Lindemann constant can be evaluated. This

Jan 1, 1962

By G. B. Craig, G. L. Montgomery

A simple technique has been developed which reveals glide dislocations in zone-refined aluminum as etch pits. The technique has been tested quantitatively by making dislocation counts on aluminum single crystals deformed in pure bending. The method has the advantage of being insensitive to the orientation of the surface under examination. A number of etchants which will reveal dislocations in annealed aluminum have been reported.'-' These etchants apparently depend on decoration of the dislocations by an impurity (probably iron) and cannot, therefore, be used to reveal fresh glide dislocations. In addition, it is unlikely that screw dislocations will be etched because of their small attraction for impurity atoms. Amelinckx6 has reported that glide dislocations located along slip traces can be revealed by fluo-rated aqua regia. This work indicated that dislocations themselves, irrespective of decoration by impurity atoms, are responsible for the etch pits. However, as a result of additional etching experiments, Wyon and Marchin7 concluded that dislocations alone are not sufficient to explain the localization of etch pits on slip lines in aluminum using chemical etch-pitting techniques. In addition, their results show that glide dislocations are revealed only on surfaces approximately parallel to a (100) axis and that not all slip systems are attacked. These observations have been confirmed by the present authors. The technique described in this article is insensitive to the orientation of the surface, and reveals dislocations on all glide systems. Fig. 1 shows a deformed region in the neighborhood of a scratch. Dislocations can be seen along three slip traces. In the present work, single crystals of zone-refined aluminum are polished for from 2 to 5 min in a chemical polishing reagent held at 100°C, as described by Herenguel and second:' H2Sol 60 pct H3PO4 30 pct HNO3 10 pet The samples are then coated with a thin oxide film by annealing in air at 450°C for 1 hr and cooling in the furnace. After deformation, dislocation etch pits are developed by immersing the specimen in the same polishing solution which is now held at a temperature of 50°C. The formation of small hydrogen bubbles indicates that etching has occurred. The sample is immediately removed from the solution. Prolonged etching results in enlargement of the pits and, in regions of high dislocation density, the overlapping and loss of resolution of individual pits. It is the experience of the authors that contact with iron will ruinously contaminate the reagent. With careful etching, dislocations separated by half a micron can be resolved optically and the method has been shown to be sufficiently sensitive to reveal dislocations produced during the pre-easy glide region of deformation, Fig. 2. A few examples of glide-dislocation configurations, as revealed by this technique, are shown in Figs. 1 to 4.

Jan 1, 1964

By J. Bulina, J. T. Brown

SINCE the inception of vacuum cold-hearth arc melting of high-temperature alloys (about 4 years ago) it has been theoretically reasoned and qualitatively found that better mechanical properties are often obtained from alloys which have been melted by this practice. Mechanical properties are emphasized in this evaluation because of their importance in design. Vacuum-arc melting has an intrinsic advantage in that it creates an ingot grain structure which is easier to forge and thus leads to greater yields for the finished product. The greater yield naturally gives the process an economic basis when the cost per pound of alloy is rather high. However, the question stressed here is, does the process create an alloy with inherently better properties? There are several theoretical reasons for suspecting that it does. One is the vacuum refining which should permit lowering the content of harmful residuals with high vapor pressures. A second is the vacuum process which should preclude the formation of harmful distribution of oxides, nitrides, or carbo-nitrides, by keeping the amount of such constituents formed to a minimum. A third is the vacuum-arc cold-hearth method which prevents the possibility of pick-up of harmful elements from a refractory crucible since none is used. Fourth, the rapid solidification rate, induced by having a cold mold and only small amounts of molten metal at any one time, minimizes the segregation of elements like titanium, aluminum, and molybdenum from the top to the bottom of the ingot. This rapid solidification rate also decreases the tendency of the above elements to form undesirable phases by segregating locally. Another less frequently considered vacuum-arc melting advantage is the inherently finer-grained ingot structure which can be produced. Finer initial grain size leads to better forgeability, and means that after relatively small reductions to billets or final forged products a finer-grained material is obtained, which, by virtue of its finer grain size, would usually have greater stress-rupture ductility.

Jan 1, 1960

By J. P. Hirth, G. M. Pound

The flux of evaporation J from {loo) , {111), and (110) planes of silver single crystals and from poly crystalline specimens was measured at 744°C under a high vacuum using a time average weight loss method. The evaporation coefficient a was computed from the measured flux and vapor pressure data through the relationship J = aPe/21rmkT. The evaporation coefficients for the close packed planes and the poly crystalline specimens were found to be about 0.4 providing tentative confirmation of the theoretical limiting law for evaporation of metal crystals, a = 73(P/Pe) + % . The (110) planes faceted to form (111) planes, resulting in a greater evaporating area and larger a . The evaporation morphology of the plane surfaces and the surfaces of spherical single crystals is described in some detail. RECENTLY a detailed theoretical analysis of the mechanism and kinetics of metal crystal evaporation was made.1'2 A mechanism for vaporization involving the recession of monatomic ledges across a metal surface was deduced, and evaporation rate equations were derived by considering the dynamics of these ledges. In most metals, the process of evaporation is one of dissociation of atoms from ledges, surface diffusion, and desorption. The gross vaporization rate is controlled by the surface diffusion and desorption rates. It was shown that the monatomic ledges, under the action of surface diffusion gradients, tend to assume a terminal velocity and spacing. Further, it was shown that crystal edges and grain boundaries are a ready source for monatomic ledges. In the case of evaporation of large, perfect crystals with clean, low-index surfaces, the kinetic treatment1 leads to a limiting law for crystal evaporation: where J is the gross flux of vaporization, a is the evaporation coefficient, Pe the equilibrium vapor pressure, m the atomic mass, k Boltzmann's constant, T the absolute temperature, and P the actual pressure of vapor over the crystal surface. Imperfections such as grain boundaries, pores, cracks, and macroscopic steps should have the effect of causing positive departures from this limiting law. Such an imperfection is a source for monatomic ledges, which, near their origin, are closely spaced and thus give rise to higher evaporation rates there. However, under many experimental conditions,' this effect should be negligible due to the relatively low concentration of imperfections. In general, a more important cause of positive departure should be screw dislocations, and the effect of these has been considered in some detai1.2,3 Monatomic ledges emanate from points where the surface is intersected by a dislocation with a screw component normal to the surface. The dynamics of ledges arising from a screw dislocation are similar to those of ledges arising from a crystal edge. There is a perturbed region near the origin where the monatomic ledges are, under certain experimental conditions, closely spaced and the

Jan 1, 1960

By R. E. Cech

Particles -3 to 30 µ diam of a 29.5 pet Ni, 70.5 pet Fe alloy after being melted and solidified while falling through a hydrogen atmosphere were found to contain a distortion-free body-centered-cubic phase. In most particles, the amount of body-centered-cubic phase approaches or equals 100 pet. It was shown by a series of experiments that this phase forms directly from the liquid. It is pointed out that a mechanism previously suggested in recent papers on the theory of the liquid-solid transformation can account for the solidification of body-centered-cubic phase in a temperature range where face-centered-cubic is the stable phase. IN the course of an operation to produce powders for a study of the martensitic transformation in Fe-Ni alloys, an unexpected phenomenon was observed. Powders of a 29.5 pet Ni alloy, which has a subzero M, temperature, were found to contain considerable amounts of a body-centered-cubic phase at room temperature. The experiments to be described were designed to elucidate the mode of formation of this unexpected phase. Experiments and Discussion An alloy containing 29.5 wt pet Ni, balance iron, was vacuum melted from high purity electrolytic iron and carbonyl nickel shot and solidified in the crucible. The metal was converted to oxide, ground to 100 pet <37µ, pressed to a cake, and homogenized 20 hr at 1350°C. To insure homogeneity the grinding, pressing, and homogenizing heat treatment was repeated a second time. The oxide was then ground to 100 pct <37µ and hydrogen reduced at 500°C. The reduced powder was given a spheroidizing treatment to produce spherical cast particles. The treatment consisted of dispersing the particles in hydrogen and allowing them to drop through a vertical hydrogen furnace set to a temperature higher than the liquidus of the alloy. This operation, commonly known as shotting, has been described more fully elsewhere.&apos; The powder produced by this method was found to contain approximately equal parts of body-centered-cubic and face-centered-cubic phases. A magnetic separation technique, previously described,&apos; was applied, and it was found that particles were either completely face-centered-cubic or nearly completely body-centered-cubic. The microstruc-

Jan 1, 1957

By T. S. Plaskett

A striated structure perpendicular to the growth axis was observed by the copper-decoration tech-nique in dislocation-free, .float-zoned silicon crystals. The striations, which were spaced about 100 p apart, fitted the relationship d = f/u , where d is the spacing, f is the growth rate, and u is the crystal rotation rate. Each stria was resolved into an UNDOPED silicon crystals pulled from quartz crucibles by the Czochralski technique usually exhibit a striated structure perpendicular to the growth axis.&apos;-&apos; This structure has been attributed to oxygen segregation, with the oxygen being introduced from the quartz crucible. If the crucible is rotated, the level of oxygen contamination has been reported as high as 10° atoms per cu cm.10 These striations are similar to solute striations commonly observed in doped Czochralski-grown crystals. The periodic nature of the striations is caused by a periodic variation in the growth rate",12 which is attributed mainly to thermal gradients in the melt.13 A finer striated structure14 attributed to constitutional supercooling is sometimes observed between the coarse striae. The oxygen striations have been observed by infrared transmission techniques,&apos; by the copper-decoration technique,&apos; by X-ray diffraction microscopy,6-8 and by 9 p absorption measurements3 on crystals pulled from the melt both with and without dislocations. In this investigation float-zoned dislocation-free crystals were examined by the copper-decoration technique. The level of oxygen for float-zone material is less than 1016 atoms per cu cm the lower limit of detection by 9 p absorption measurement. EXPERIMENTAL TECHNIQUE The crystals were grown by the float-zone process with the rf heating coil outside of the quartz envelope containing the silicon. All float zoning was done under an atmosphere of purified helium. The Dash technique15 was used to grow the crystal dislocation-free. This involves growing the crystal initially with a diameter between 2 and 3 mm and at array of starlike precipitates of copper. The strucLure was not .found at the surface tor a depth of about 1.5 mm, or in a region of similar width ahead of a dislocation network. The structure is postulated to consist of vacancy clusterings or dislocation loops. very rapid rates, about 20 mm per min, for a distance of about 3 cm. The diameter of the crystal is then increased to the diameter of the source of silicon, which in this case was about 19 mm. Because of the arrangement of the apparatus, the zone was passed downward rather than upward, contrary to the standard float-zoning practice. Also, the source was rotated rather than the seed. ziegler17 has made dislocation-free crystals by a similar technique but has passed the zone upwards. The starting material was zone-refined and had a p-type resistivity of 150 ohm-cm. The major impurity was boron; the total impurity excluding the boron was reported by the supplier (Dow-Corning) to be typically less than 2 x 1013 atoms per cu cm. The crystals were examined by the Dash copper-decoration technique18&apos;19—a method in which about 10" atoms per cu cm of copper are diffused at a temperature between 900" and 1000°C into silicon which is then quenched to room temperature. On quenching, the copper precipitates on crystalline defects which are then visible when viewed by transmission infrared microscopy. The photomicrographs shown were taken either of the infrared image tube screen or directly on infrared film. All sections prior to decorating were chemically polished and, for some sections, given a sirtlZ0 dislocation etch-pit examination. After decorating, the samples were mechanically polished. RESULTS A photomicrograph, taken in transmission of a decorated cross section, is shown in Fig. 1. The portion of the section shown is near the surface of the crystal. The entire cross section showed no dislocation etch pits after being given a Sirtl etch treatment. It is seen that the copper precipitated randomly. Each precipitate, as has been reported by others, was found to have a starlike structure.

Jan 1, 1965

By Jack Washburn

The c axis indentations in zinc crystals were shown to undergo 100 pct strain recovery on heating. The mode of deformation and the details of the polygonization and collapse of indentations were found to be consistent with a number of predictions of dislocation theory. IT has been observed that indentations of the type made by Jillson in zinc crystals can, under certain conditions, be macle to shrink in size and even completely disappear on heating the crystal, Fig. 1. The details of the process are interesting in that they afford a simple experimental illustration of several predictions of dislocation theory. The indentations were made as shown in Fig. 2 by dropping a 0.2 g indenter with a hemispherical point on a cleavage face of a. zinc crystal. Such indentations can be made through relatively great thicknesses in the c-axis direction. The material close to the struck surface, where mechanical twinning and complex distortion occurred, was removed by cleavage. The remaining crystal, which was about 1 cm thick, contained a dent that was accurately round and almost the same size on both top and bottom surfaces. The indentation had the form of a flat cone of half angle approximately 89°, having a rounded apex and a rounded intersection with the unde-formed part of the crystal. The mode of deformation was primarily slip on the three equivalent systems in the basal plane. Six types of dislocation loop (plus and minus in each of the three directions) as shown in Fig. 3 are required by the geometry of the indentation. The maximum shear strain in the radial direction amounted to approximately 2.5 pct. Slip was distributed in the usual, microscopically nonuniform, manner among the possible parallel slip planes. When an indentation was made close to a (2110) surface, three separate sets of interpenetrating slip bands appeared as shown schematically in Fig. 3. A small biaxial tension strain in the plane of (0001) is also a geometrical necessity for the formation of an indentation. If the sides of the dent are tilted l° with respect to the undeformed surrounding crystal, and the surround-

Jan 1, 1957

By Larry Kaufman, Martin Schatz

The effect of pressure on the extension of the ? loop in the FeSi system has been determined by means of metallogvaphic studies and hardness measurements performed on a series of high-purity Fe-Si alloys containing 7.5, 11.0, and 13.9 at. pct Si, respectively. These mensurements, performed at 42 kbar and temperatures up to 1200oC, indicate that the ? loop is expanded to about 10 at. pct Si at 42 kbar as opposed to a maximum extension of 4 at. pct Si at 1 atm. Comparison of the experimental results with thermodynamic predictions of the pressure shifts yields satisfnctory results. DURING the past few years, several studies have been performed in our laboratory1-&apos; in order to determine the effect of high pressure on phase equilibrium in pure iron and iron-base alloys. The purpose of these studies has been to elucidate the effects of high pressure experimentally and to compare the observed results with predicted pressure effects derived on the basis of known thermody-namic and volumetric data at 1 atm. These studies have included work on pure iron2,5,7 as well as Fe-Ni,1,5 Fe-cr,l,5 and Fe-c4-6 alloys. In addition, Tanner and Kulin3 have reported results of pressure studies on two Fe-Si alloys containing 2.0 and 6.25 at. pct Si. At the time of this latter study, no detailed information was available concerning the difference in volume between the a (bcc) and ? (fcc) phases in the Fe-Si system as a function of silicon content. In order to compare their observations with calculated pressure shifts, Tanner and Kulin were forced to assume that silicon had no effect on the difference in volume between a and ? iron. The resulting discrepancy between their calculation of the a/? phase boundary at 42 kbar and the observed results led them to the conclusion that silicon additions probably decrease the difference in volume between a and ? iron. Recently: Cockett and Davis8,9 have reported de- tailed studies of the lattice parameters of a series of Fe-Si alloys at temperatures ranging from 20" to 1150°C. These measurements, performed on alloys in the bcc and fcc range, show that silicon does indeed decrease the difference in volume between a and ? iron. By correcting the calculations of Tanner and Kulin in line with the observed effect of silicon they were able to show improved agreement between computed and observed pressure shifts.&apos; The present measurements were undertaken to provide additional corroboration of this effect, by extending the range of composition, in addition to exploring a situation where large extensions of a ? loop could result in impingement of the ? field with an ordered bcc phase (based on Feo.75Sio.25). I) EXPERIMENTAL PROCEDURES AND RESULTS The alloys investigated were obtained from Dr. F. Kayser of M.I.T. They were prepared at the Ford Scientific Laboratory by vacuum melting electrolytic iron and high-purity silicon. The melts were poured under an argon atmosphere into hot-topped steel molds. Subsequently the ingots were hot-worked down to 1/2-in.-diam rods. Three alloys containing 7.5, 11.0, and 13.9 pct Si were studied. Carbon, regarded as the principal impurity, analyzed at, or below, 0.001 wt pct for all of the alloys. Prior to pressure-temperature treatment, the rod was annealed for 24 hr in vacuum at 1000°C, water-quenched, and subsequently machined into 0.100-in.-diam by 0.100-in.-long specimens. Subsequent to machining, the specimens were again annealed and then examined metallographically. They were found to exhibit a clear coarse-grained ferrite similar to Figs. 10 and 110 of Ref. 1 and Fig. 2 of Ref. 3. Subsequently, specimens of each alloy were equilibrated at 42 kbar at various temperatures in supported piston apparatus.1,3,4,6 Three specimens, one of each alloy, were wrapped in platinum and exposed simultaneously. The pressure-temperature cycle consisted of increasing the pressure from ambient to 42 kbar at 25oC, heating rapidly to the desired temperature, holding for 15 min, and quenching to 100°C, followed by slower cooling to 25°C and pressure release. The temperature was measured with a Pt/Pt-13 pct Rh thermocouple which was not corrected for pressure effects. Subsequently, specimens were examined metallographically and by

Jan 1, 1964

By M. J. Spendlove, H. W. St. Clair

An automatic surface-follower mechanism was used to measure the surface temperature and the rate of evaporation of molten zinc while undergoing distillation at low pressure. At pressures of 50 to 100 microns Hg, the rate of evaporation may be 60 to 80 pct of the theoretical maximum corresponding to the measured temperature. The rate decreases rapidly at pressures above 100 microns Hg, to only 7 pct of the theoretical maximum at 2000 microns Hg. Measured temperature gradients at the surface are in agreement with theoretical gradients calculated from the heat of vaporization, rate of evaporation, and thermal conductivity of molten zinc. THIS paper covers a series of careful measurements of the rate of evaporation of molten zinc at low pressures, undertaken as a part of a general investigation on the separation of metals by distillation.&apos; Preliminary results on the rate of evaporation of zinc were described in a previous publication.&apos; The results of the preliminary study are as follows: 1—The observed rates of evaporation of zinc were as high as 0.60 g per sq cm per min, corresponding to 75 pct of the maximum rate as determined by the temperature of the metal. 2—The evaporation rate was decreased from 0.30 g per sq cm per min at a pressure of 25 microns to only 0.09 g per sq cm per min when the pressure was increased to 5 mm. 3—The temperature of the surface during rapid evaporation may be 35" to 100°C less than the main mass of metal. Observed temperature gradients at the surface are consistent with those calculated from the heat of vaporization and the thermal conductivity of liquid zinc. In the preliminary experiments, the rate of evaporation was determined from the loss in weight of the metal in the crucible during the test. This was compared with the weight of the condensate. Obviously, this method was subject to error because of some uncertainty as to the period of evaporation that should correspond to the loss in weight. Furthermore, this method did not permit taking into account probable variations in the rate of evaporation during the test. A more serious shortcoming was that the temperature at the evaporating surface was not known, so the maximum theoretical rate of evaporation could not be calculated accurately. To improve the precision of rate measurements and to allow accurate measurements of the surface temperature, a refined experimental technique was developed whereby the temperature and position of the surface could be measured accurately. This was done by an automatic surface-follower mechanism that gave a continuous record of the surface temperature and the amount of zinc evaporated. These data permitted an instantaneous determination of the actual rate of evaporation and the surface temperature throughout the test. Pure zinc (99.99 pct Zn) was used in all experiments. Description of Distillation Furnace The details of construction of the distillation furnace are shown in Fig. l. The furnace is enclosed in a large quartz tube open at one end. The open end of the quartz tube is sealed by means of a water-cooled head. The entire enclosure is vacuum tight. The connection between the water-cooled head and the quartz tube is sealed by means of a rubber gasket which fits tightly against the smooth end of the tube. The metal is heated by induced current set up in the molten zinc and the walls of the graphite crucible by means of a water-cooled induction coil outside the quartz tube. High frequency current is passed through the induction coil from a 6 kw mer-

Jan 1, 1952

By W. A. Krivsky, R. Schumann

Relaxation calculations were made to find the temperature distribution in the refractory wall surrounding a tuyere pipe. One set of boundary conditions approxiwated those of a standard copper converter while a second set represented a hypothetical,thermally insulated tuyere pipe. The pvonounced cooling effect of the tuyere is related to the problem of tuyere accretions and tuyere bunching. PARTICULARLY severe service conditions are encountered in the tuyere zone of a basic copper converter during the matte blow. High-temperature gradients result from the juxtaposition of the hot charge and the tuyere pipe carrying cold air. Refractory temperatures near the tuyere mouth vary rapidly with time when the blast is turned on or off. One of the most serious problems is the formation of magnetite accretions around the tuyere mouth, which occurs very rapidly under some operating conditions. As a result, continual punching of the tuyeres is necessary to maintain a satisfactory rate of converting. Thus, if some means could be devised for keeping the tuyere pipes fully open without necessity for punching, a considerable saving should result from reduction in labor cost, maintenance of full capacity operation, and reduction in refractory costs. Recent quantitative studies of the FeO-Fe,O,-SiO, system1 have shown that the solubility of magnetite in iron-silicate slags decreases rapidly with de- creasing temperature in the range of copper converting temperatures. Data taken from the previous work are shown in Fig. 1. Thus the formation of "horns" and accretions around the tuyere mouth can result simply from the freezing out of magnetite on the cooler areas of refractory immediately adjacent to the tuyere pipes. One might expect to find a correlation between the severity of magnetite formation and the area of refractory surface which is cooled below the freezing point of the slag (corresponding to magnetite precipitation). Magnetite itself has a very high melting point (1597°C) so that the accretions are for all practical purposes infusible. The accretions might be dissolved by increasing the silica content of the slag and at the same time raising the temperature, but this approach is not feasible because it results in rapid corrosion of the basic refractory lining of the converter as a whole. In order to understand better the problems involved in tuyere design and operation, it seems essential to have at least a semiquantitative idea of the thermal conditions in the tuyere zone. Accordingly, this paper presents results of thermal calculations for an idealized copper converter tuyere. These calculations follow the so-called relaxation

Jan 1, 1960

By Don M. Jackson, Robert W. Howard

Techniques for the epilaxial growth of single -crystal silicon carbide films on silicon were developed. The vapor-phase decomposition and bydrogen reduction of silicon tetrachloride (SiC14) and Propane (C3H8) resulted in clear films of silicon carbide, lip to seveval microns in thickness. The growth took place in a horizontal . silicon epilaxial reactor at 1100°C (pyrometer) at a rate of- 3000Å per minute. Electron diffraction and X-ray diffraction studies demonstrated that the films were single-cyrstal, ß -phase, or cubic silicon carbide. SiO2 film were used to mask areas of the silicon sur-lace in order that the silicon carbide might be grown in controlled geometries. Both n- and p-type films were grown on p-type silicon waters.. Heavily doped silicon films of the same conductivity type as the silicon carbide films were deposited over the silicon carbide in order to affect better probe contact to the structures. when n-type silicon carbide mesas were grown on p-type silicon substrates the de vollage-current relationships between films and substrates were that of junction diodes. These diodes showed a sensitivity to while light ill that the incident light increased forward- and reverse-satro,ation currents, P-type silicon carbide mesas grown on p-type silicon were ohmic rather than rectifying in their voltage -current relationship. No conclusions could he reached concerning heterojunc-tiou rectification in the structure. SILICON carbide is a semiconductor with many interesting properties. It decomposes at temperatures above 2200°C.1 It occurs in two general crys-tallographic forms—hexagonal (a Sic) and cubic (ß Sic)—with the cubic form having a forbidden-gap energy of 2.32 ev and the hexagonal form (specifically the 6H polytype) a gap energy of 2.86 ev.3 It behaves as an extrinsic semiconductor at temperatures approaching 5003C. It has been shown to have a high resistance to radiation damage4 and p-n junctions formed in Sic have been shown to radiate visible light under forward- or reverse-bias conditions. Epitaxial silicon carbide on silicon carbide has been successfully grown through the use of a variety of techniques, such as gaseous cracking of SiCL4 and CC4, nearly all of which require a deposition temperature above 1500°C.6 This paper will cover very recent work on the gas-phase deposition of highly ordered films of silicon carbide on high-quality silicon single-crystal substrates. The films have been shown to ex- hibit junction-rectification properties when geometrically isolated regions are electrically biased with reference to the silicon substrate. There will be no discussion of the mechanism of heterojunction rectification, but the methods of film fabrication, geometry control, and structural evaluations will be covered in detail. Electron diffraction, X-ray diffraction, and diode electrical properties were used to characterize the films and the junctions. GAS-PHASE DEPOSITION OF Sic The techniques for the deposition of silicon carbide films were a logical outgrowth of the standard silicon epitaxial process. The major premise followed was that, for any film to nucleate in an ordered fashion where there is considerable mismatch in lattice parameters (in this case 22 pct), an extremely clean, damage-free substrate surface must be presented to the gas stream. Thus a standard gas-phase HCl etching step was used to prepare the substrates for growth. A minimum of 5 µ of substrate-surface material was removed prior to the deposition of Sic overgrowth films. The techniques used for growing silicon carbide films were those of growing silicon alone, with the added injection of a hydrocarbon gas into the hydrogen and silicon tetrachloride gas stream. The hydrocarbon gases used thus far have been research-grade (99.99 pct) methane (CH4) and propane (C3H8). Propane ultimately gave the best results. The gas flows were controlled through a panel shown schematically in Fig. 1. A hydrogen main stream of 30 liters per min passed through the horizontal quartz-tube epitaxial reactor, while SiC14, C3H8, HC1, and doping gases were injected as side streams. The

Jan 1, 1965

By K. G. Wikle, J. G. Klein, W. W. Beaver

Consolidation of hydride process, electrolytic, calcium reduced, and comminuted thorium powder, as well as saw chips and lathe turnings, by vacuum hot pressing and by cold pressing-vacuum sintering was studied. The mechanical properties of the consolidated material in the extruded form are compared with those of wrought castings. AT present there little little industrial use for thorium metal, although it has some important though small scale applications in electronic equipment. Despite its high inelting point—about 1750°C —a low modulus of elasticity, 11.4xl0 si at 20°C;&apos; relatively low mechanical properties coupled with a high density, 11.7 g per cu cm; and an unusually high chemical activity with normal atmospheres limit any structural applications. The metal is utilized as an alloying element principally in magnesium. Pure thorium finds utility as electrodes in gaseous discharge lamps such as the high intensity mercury lamp&apos; because its low work function and high electron emissivity provide lower starting potentials and more uniform operating characteristics than other available materials. The metal is also found in photoelectric tubes used for the measurement of the ultraviolet spectrum." Thorium metal has been used in germicidal lamps of the cold cathode type as sputtered coatings on nickel in order to provide a low work function surface and a low starting voltage. Other applications have involved the radioactive properties of thorium for the production of ionized particles." The potential value of thorium is much greater than its present use pattern because of possible utility in the field of nuclear power. Th may be converted through nuclear reaction to a fissionable element U which should be capable of acting similarly to U in the g&apos;eneration of atomic power. Thorium has been reported to be about three times as plentiful as uranium in the earth&apos;s crust, placing it in the order of abundance of lead and molybdenum." Thus, it is of interest in augmenting the potential supply of fissionable material for nuclear power. Because of its high melting point, thorium is usually produced as a powder through the calcium reduction of its oxide or thermal reduction of halides by sodium, magnesium, and calcium. It may also be produced in flake form by electrolysis of fused alkali or alkaline earth chloricles and fluorides. Therefore, powder metallurgy assumes importance in the fab- rication of thorium metal shapes. Furthermore, it is rather difficult to obtain pure thorium by melting, as the molten metal reacts readily with graphite as well as oxide, carbide, and nitride refractories. These contaminate the melt with oxides, carbides, and metallic impurities." The current investigation was undertaken to examine the fabrication of thorium by powder metallurgy methods which have been used for the commercial production of beryllium and other metals.&apos; A sparcity of data concerning the comparative cold and hot compaction of thorium powders of different derivation existed. Therefore, all commercially available types were examined along with other experimentally produced thorium powders in order to round out the comparison of consolidated thorium powders with melted reguline metal. Review of the Literature By heating a mixture of ThC1, with potassium, Berzelius made the first thorium metal as an impure powder in 1828. Improvements in the basic process, increasing thorium assay to 99 pct, were made by several investigators including Arsem," Lely and Hamberger10 and Von Bolton." Calcium reduction of Tho, to make powders was investigated by Berger," Huppertz,&apos;" Kroll," and Kuzel and Wedekind.&apos;" A thorium powder produced by this method using a CaC1, fluxing agent assayed 99.7 pct, as reported by Marden and Rentschler.&apos;" Compacted and sintered, this product was found to be ductile, and could be fabricated into wire and sheet. Improvements of the calcium reduction process were made later" wherein CaCl, was eliminated from the reaction, producing metal assaying 99.8 pct Th. Further work by Lilliendah118 howed that a coarser metal could be obtained by the substitution of ThC1, or ThOC1, for oxide with consequent advantage of stability to atmospheric reaction. Reports on the technology of thorium developed in Germany during World War II have been made by Espe."&apos; Thorium powder of 99.5 pct Th was obtained by reduction of the oxide by calcium. Screening to —200 mesh, compacting with about 20 tsi, and sintering in vacuo at 1320" to 1360°C for 3 hr resulted in a porous sinter cake. The sinter cake was sufficiently ductile to be worked into bar, wire, and sheet which could be employed as electrode materials.

Jan 1, 1957

By H. H. Klepfer, M. E. Snyder

UNTIL very recently, the commercial availability of the rare earths as metals has been very limited. Fabrication of mill products from these metals has not been studied in most cases. This note reports the results of the development of fabrication techniques for thulium. Thulium has a melting point of about 1550°C and a hexagonal close-packed crystal structure. It oxidizes in air to give a black oxide (TmzO,). The procedures for producing thin thulium foil were developed on one ingot weighing about 220 g and were subsequently applied in processing about a pound of metal to foil. Several alternates for the various fabrication steps were investigated and will be discussed. The metal fabricated was in the form of commercial chill-cast ingots 1 in. in diam and 2 in. long weighing approximately 220 g. Impurities in the ingots were reported by the vendor to be 4000 ppni tantalum, 2000 ppm calcium, 200 ppm nickel, 100 to 200 ppm iron, 100 ppm europium, and less than 100 ppm copper, lutetium, and ytterbium. In addition to these impurities, several salt-like inclusions as large as l/8 in. in diam were revealed along the center line of the one ingot sectioned. Preliminary tests indicated that small wafers cut from the as-cast ingot would not fabricate readily by rolling. Forging of copper-jacketed wafers was therefore attempted. At 1550"F forging was satisfactory but an apparent reaction of copper with thulium demanded investigation of lower temperatures. Therefore, the remainder of the test ingot was forged at 1450°F—with only minor cracking. All ingots forged were inserted into copper tubes of 1 in. 1D and 0.125-in. wall thickness. The jackets were sealed by flattening the ends of the tubes and welding under helium. Heating time at 1450°F was 30 min. Press forgings of 1/8 in. per pass were used, followed by 10 min reheats. When the ingots had been squared and reduced to 0.250 in. in thickness, the original copper jacket was stripped off and replaced by a new jacket in preparation for hot rolling. Hot rolling at 1450" F without edge cracking was readily accomplished after forging. Excellent results were obtained with 10 pct reductions of thickness followed by 5 to 10 min reheats. After reduction of thickness from 0.250 to 0.100 in. the copper jacket was removed. It was found, in fact, that hot rolling in air was possible. A tenacious black oxide similar to that seen on zirconium was formed during 3 min reheats at 1450°F. Reduction in air to 0.010 in. foil was possible taking 10 pct reductions per pass. The oxide coat formed during hot rolling in air could best be removed by sand blasting and pickling. Common pickling solutions containing polar solvents were found to attack the metal too rapidly and a concentrated nitric-hydrofluoric acid mixture attacked neither the oxide nor the metal. The most satisfactory pickling solution was 52 vol pct concentrated nitric acid-48 vol pct glacial acetic acid. After forging to 0.250 in., vacuum annealing and cold rolling was found to be another satisfactory alternate to hot rolling in copper jackets. After forging. a hardness of Rockwell B76 was found. Annealing in vacuum (2 X l0-5 mm of Hg) for 1 hr at 1200"F did not alter this value. Annealing at 1470"F for 1 hr brought the hardness down to R;]63. With cold rolling (5 pct per pass) the hardness returned to about R,1:76 after 20 pct reduction and edge cracking became noticeable. However, cold rolling to a total of 40 pct reduction in thickness (RI,,83) was possible before edge crack propagation became serious. Good surface finish was obtained, and the metal loss due to oxidation was minimized by cold rolling and vacuum annealing. Using this procedure the yield of 1.25 in. wide by 0.010 in. thick foil from a 1-in. diam ingot was about 40 pct.

Jan 1, 1961

By H. W. Weart, S. Sarian

The effect of temperature, impurities, and capillarity on the morphology of solid particles in a liquid matrix is investigated. For the NbC-liquid iron system, at least, it is found that those particles which are dissolving will assume a spherical shape; all others will take the equilibrium shape. An aniso-tropic-isotropic transition in the interfacial tension is observed at about 1725°C and is believed to be due to cooperative roughening of the interface on an atomic scale. The presence of boron as an impurity lowers the transition temperature. It is the purpose of this paper to present some observations on the shape of particles in a system consisting of solid particles of nonuniform size in a liquid matrix and to propose an explanation for these observations. Although qualitative only, these observations, which were made largely in the course of a study of particle-growth kinetics,&apos; are presented separately to draw attention to them. They reveal the effect of particle size, temperature, and surface-active impurities on particle shape in the solid NbC-liquid iron system, and are unique in demonstrating for the first time the reversible effect of temperature and impurity concentration on the anisotropy of the energy of the solid-liquid interfaces involved. Some of the changes are readily demonstrated with the few particles whose shapes can be clearly shown in a single photomicrograph, but others are more subtle, and their recognition depends upon some experience and the study of a much larger sample. For this reason, a certain amount of judgment must be acknowledged to underlie their detection, but every effort was made to maintain proper scientific objectivity in evaluating the observations. The observations will be presented first, and then an explanation rationalizing them will be offered. Brief reminders of the relevant theories will be inserted into the discussion at the appropriate points. EXPERIMENTAL The observations were made of the shape of NbC particles dispersed in liquid iron. All specimens were prepared by pressing NbC powder and car- bony1 iron powder treated according to standard powder-metallurgical procedures. Except as noted, all specimens were heated to 1900°C for 1/2 hr. In investigating the effect of temperature, specimens were quenched from 1900°C to various lower temperatures, and held long enough (48 hr) to allow shape changes to be completed. Boron, in the form of amorphous powder, and oxygen, in the form of niobium oxide powder, were added as impurities to some specimens. Nitrogen was introduced into some specimens as an impurity by heat treating the samples under a partial pressure of nitrogen gas. Effect of Temperature. It was first observed that increasing temperature had a marked effect in rounding the corners and edges of the particles. At 1900°C the particles are all very nearly spherical and remain so even when the holding time is extended to 48 hr. As can be seen from Fig. 1, there is virtually no evidence of the presence of low-index planes. At 1300°C, however, most of the particles are prismatic with very slightly rounded edges and corners, Fig. 2. (In making these shape comparisons, attention should be focused on the larger particles, even though the smaller ones would be expected to be the first to change shape, owing to the shorter diffusion distances involved. The cause of the mixed particle shape will be explored later.) It was established that these structures are re-

Jan 1, 1965

By M. Semchyshen, A. P. Coldren, W. G. Scholz

A survey of the Fe-Mo-B system was made to determine the extent to which boron might affect the microstructure and strength properties of iron-rich Fe-Mo alloys. Seventeen vacuum-induc tion melted ingots were prepared with approximate mo-lybednum contents 01 2, 4, 9, and 19 wt pct and approximate boron contents of 5, 14,50, and 630 ppm by weight. Phases extracted from annealed specimens and identified by X-ray diffraction analysis included E Fe3MO2, Mo2FeB2, and an oxide hazling a diffraction pattern practically identical to that of MnFez04. Three types of alloy strengthening were observed: a) Fe-2 pct Mo-B alloys were hardenable by a bainitic-type transformtion of austenite to ferrile; b) Fe-4 pct Mo-B alloys were hardened by the solid-solution mechanism; and C) Fe-9 pct Mo-B and Fe-19 pct Mo-B alloys were hardened by precipitation of the E Fe3Moz phase. Boron was observed to have a very strong hardening effect in the Fe-2 pct Mo alloys and a mild strengthening effect on the 1200°F properties of the Fe-4 pct Mo alloys. In the Fe-9 pct Mo alloys there was no consistent effect of boron on strength at either room temperature or 1200° F. The Fe-19 pct Mo alloys were so brittle that meaningful tensile or creep-rupture data could not he obtained. No dispersion hardening from borides was recognized in any of the alloys. THE high-temperature properties of iron-rich Fe-Mo alloys were studied by Reiter and Hibbard.&apos; They found that the strength of the alloys increased markedly as the molybdenum content was raised from 0 to 15.9 pct by weight. The authors explained the various degrees of strengthening in terms of a) solid-solution strengthening, b) strengthening from a martensitic-type transformation, and C) precipitation strengthening. The objective of the present investigation was to conduct a survey of the iron-rich portion of the Fe-Mo-B system to determine whether the beneficial effects of molybdenum on the strength of iron can be enhanced by the presence of boron. It was thought that perhaps fine dispersions of molybdenum- rich borides would be found which could improve the high-temperature strength of ferritic alloys beyond the strengths attainable by molybdenum additions alone. Also, it was considered worthwhile to study the effects of boron in alloys which are essentially free of carbon. Most, if not all, of our present knowledge about boron in high-temperature alloys pertains to alloys containing carbon and carbides. EXPERIMENTAL PROCEDURES Alloy Preparation. Seventeen 8-lb ingots were prepared from electrolytic Plastiron, unalloyed molybdenum chips, and boron powder by a vacuum-induction melting procedure employing hydrogen as the main deoxidizer. The alloys were melted in alumina crucibles using 35-lb melts, with each melt being split four ways to achieve varying boron contents within four Fe-Mo base compositions. The results of chemical analyses conducted on chips machined from the chilled ends of the ingots are presented in Table I. The boron analyses were performed by a spectrographic method utilizing primary standards prepared from boric acid. They indicated that in ingots to which no boron was

Jan 1, 1964

By A. Yoshikawa, T. Sugeno

To detenmine the origin of the sharp fatigue limit in many ferrous metals, S-N curvces were determined in push-pull fatigue at 18.6 kc per sec at room temperature and - 67°C for various kinds of iron. These differed in carbon and nitrogen content and grain size (25 and 2500 grains per sq mm). Sharp knees appeared in the S-N curves of fine-grained specimens with high interstitial content at both temperatures. In coarse-gruined specimens, knees were not sharp .regardless of the interstitial content and the testing temperature. It is concluded from these results that the controlling factov for the appeavance of the sharp fatigue limit is not strain aging but grain size. It has already been mentioned that iron and some iron alloys show fatigue curves characterized by a more or less sharp knee and a fatigue limit.&apos; Explanations of the origin of this characteristic behavior of fatigue in iron appear to fall into two broad classes. One is concerned with interstitial impurities through the aging or locking effect and the other is based on intrinsic plastic properties of iron. Levy and sinclair2 and Lipsit and Horne3 have suggested that the fatigue limit observed generally in iron and steel is due to interstitial solute aging during the course of fatigue. On the other hand, Thompson and wadsworthl have suggested that the fatigue limit may be analogous to the sharp yield point, in the sense that alternating stresses above the fatigue limit are necessary to break down the dislocation locking. However, there are some experimental results which cast doubt on the general validity of these suppositions.4-8 Ferro and Montalenti&apos; have recently suggested that the typical form of the fatigue curve of iron seems to come from an intrinsic plastic characteristic of the iron structure. This suggestion is based on the experimental result that pure iron containing an insufficient amount of interstitial solute to alldw appreciable strain-aging effects still showed a fatigue limit at room temperature. Although systematic investigations on the effect of grain size are scarce, it is thought that this effect is important in relation to the sharp fatigue limit. In fact, sensible differences were observed by Oates and Wilson4 in fatigue behaviors of a low-carbon steel depending on grain size. Ideally, the fatigue limit will be defined as the upper limit of stress amplitudes at which the life is infinite. This ideal fatigue limit can never be determined because of the limited time available for experiments, and moreover scattered data points often leave the construction of S-N curves somewhat arbitrary. And, it is usually stated that a material has a fatigue limit when its statistical S-N curve becomes apparently parallel to the N axis with a more or less sharp knee. The meaning of the fatigue limit and the sharp knee will be discussed in detail in relation to the construction of S-N curves. The origin of this sharp knee was investigated by determining S-N curves of iron for various interstitial contents, testing temperatures, and grain sizes. EXPERTMENTAL PROCEDURE The material as received was a commercial-grade low-carbon steel in the form of 4-mm-diameter rods with the following composition in weight percent: C, 0.09; Si, 0.017; Mn, 0.43; P, 0.006; S, 0.012; Cu, 0.11. Five different groups of specimens were tested. The grain sizes and interstitial contents in each group are given in Table I. The specimens of group C-02 were decarburized by annealing for 100 hr at 700°C, followed by an additional 100 hr at 850°C in wet hydrogen. Specimens of groups N-2 and N-9 nitrided in a mixed atmosphere of ammonia and hydrogen after an initial treatment identical to group C-02. The specimens of group C-9 were prepared in a manner similar to those of group C-02, except heat treatment was carried out in a vacuum. After polishing their surface with emery paper, the specimens of all groups were given a final anneal for 30 min at 930°C and air-cooled in an evacuated quartz tube. For group C*-9 this

Jan 1, 1965