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By Claude H. P. Lupis, John F. Elliott

WAGNERL described the thermodynamic behavior of a metallic solution consisting of the solute components 2, 3,... at very dilute concentrations in the solvent, 1, in terms of the excess partial molar free energy, or alternatively the logarithm of the activity coefficient. In the Taylor expansion of FiE, by neglecting the second- and higher-order differential terms, he obtained the excess partial molar free energy as a linear function of the composition:

Jan 1, 1965

By W. W. Mullins

The development, by the mechanism of volume diffusion, of a grain boundary groove on an interface separating a solid phase and a saturated fluid phase is calculated under the following assumptions: 1) isotropy of the interfacial free energy 's, 2) applicability of the Gibbs-Thompson formula relating curvature and chemical potential, 3) a nearly planar groove surface. The groove profile is found to have a .fixed shape and linear dimensions that are proportional to (time)1/3. The constant of proportionality is evaluated and involves ?s in such a way that the latter my be evaluated from grooving kinetics. A grain boundary that intersects an interphase interface will tend to produce a groove along the line of intersection (Fig. 1). The ultimate motivation for the formation of the groove is the reduction in interfacial free energ that occurs as the grain boundary contracts. Smith has discussed the way in which the free energy of the interfaces determines their dihedral angles at the groove root and has given a detailed analysis of the effect of the dihedral angles upon the nature the microstructures developed in polyphase alloys. Herring2 has extended the considerations to the case in which the free energy of an interface depends on its orientation and has thus added torque terms to the condition for interface equilibrium. Although the grain boundary and the interphase interface may quickly achieve the proper equilibrium angles where they intersect, as determined by the balance of equivalent tensions and torques,"2 the groove will generally continue to grow (though with decreasing speed). Thermodynamically, this continued growth reflects the fact that the lowest free energy cannot be achieved until the entire grain boundary has been eliminated; mechanistically, the growth occurs in response to inhomogeneities of chemical potential that are caused by the constraint of fixed dihedral angles at the groove root. The latter point is illustrated in Fig. 1, which shows that atoms* near the groove are on the convex side of a cylindrically curved surface and therefore have a higher chemical potential than atoms on the flat interface beyond. Thus atoms tend to leave the curved groove shoulders and force the groove to deepen. When both phases are solid additional chemical potential gradient may arise from whatever strain fields are necessary to insure continuity of material across the interface. In the discussion that follows, this complication due to strain will not occur since we always assume one phase to be a fluid and the other a solid. The transport processes that may operate to permit growth of the groove are surface (interface)* diffusion, volume diffusion in the phase on either side of the interface, and evaporation-condensation when one phase is gaseous. (Also a melting-freezing process is responsible for groove enlargement on a solid-melt interface as discussed in Sec. VI). The theory of grain boundary grooving has been discussed in detail for the processes of surface diffusion and of one type of evaporation-condensation.3'4 The theoretical predictions for the mechanism of surface diffusion have been confirmed in an experimental study of grooving kinetics in copper.5 In this paper, the theory of groove development by volume diffusion will be presented. Volume diffusion can occur either by atomic migration in the solid S or by diffusion of S atoms in the fluid phase, or solvent, adjacent to the groove surface. As discussed in Sec. V, volume diffusion in the solid should be less important than surface diffusion in forming a groove. Apart from verifying this statement then, the main interest of the present analysis centers on diffusion in an external solvent. The solvent may be either a liquid (e.g., water, or-

Jan 1, 1961

By D. J. Rose, W. M. Armstrong, A. C. D. Chaklader

The wetiability of single crystals of MgO by specimens of vacuum-cast iron was studied using the sessile drop technique in vacuo at 1550ºC. Formation of FeO at the liquid-vapor interface caused the contact angle (6) to decrease from 117 to 65 deg during the first minute. After cooling, all specimens possessed a peculiar annular interfacial deposit of Fe,O,. Within the annulus the interface showed no sign of chemical attack. Chemical reaction occurred where the iron was alloyed with Ti, V, Cr, Nb(Cb), Ta, cmd Zr. Vnriation of 6 with alloy concentration was studied. Although vanadium md chromium improved the wettability of MgO by iron, the effect of Zr, Ti, Nb, and Ta was indeterminable because the 8 derived from sessile drop considerations was that of the metal against a restrictive peripheral volume of liquid oxide wetting the substrate. INTEREST in the nature of metal-ceramic interactions has been stimulated by progressive development of metal-ceramic combinations. One of the more valuable methods used for bond investigation is the sessile drop technique. Recent attempts to improve wettability through solute additions to the drop have revealed that the solute may a) react with the oxide creating new compounds at the solid-liquid interface, b) adsorb at the interface in a monolayer formation, or c) distribute throughout the drop uniformly causing no wettability variation. This work, part m in a series1,2 of investigations of interface reactions between metals and ceramics, embraces a study of the Fe-MgO system. MgO possesses a large negative free energy of formation (-173 kcal per g mole 0, at 1550ºC) compared to FeO at this temperature (-73 kcal per g mole 02). Hence the electronegativity difference between the drop and substrate allows selection from an extensive group of metal solutes with intermediate electronegativity differences that would, in the absence of chemical reactions, be expected to adsorb pre- ferentially at the solid-liquid interface. However, the high oxygen pressure of MgO in vacuo creates an oxidation environment which influences the solute behavior. Recent studies of the Fe-MgO systemJ74 have been restricted by chemical reactions that obscured observations. In view of the current interest in liquid-phase sintering of metal-ceramic combinations under partial oxidizing conditions,5 a more comprehensive study of the Fe-MgO system seemed beneficial. EXPERIMENTAL PROCEDURE 1) Specimen Preparation. Optical-grade MgO single crystals and Ferrovac vacuum-cast iron were used in all experiments. A spectrographic analysis of the iron indicated 0.008 C, 0.05 V, 0.005 Mo, 0.01 Ti, 0.002 S, 0.01 Cr, 0.001 Mg, 0.01 Si, 0.003 A1 (wt pct). The MgO crystals were cleaved along (100) planes into plates approximately 20 by 20 by 1 mm. Following annealing in vactto at 1100°C for 3 hr, surface irregularities were removed by a chemical etch in phosphoric acid at 100°C. The iron rod was machined into small cylinders approximately 0.250 in. in diam and length and these were carefully cleaned in organic and acidic solutions to remove surface impurities. All specimens were weighed to the nearest 0.1 mg. 2) Apparatus. The apparatus described in detail by previous authors1,2 was modified for this work. A Polaroid Land Camera was incorporated into the optical system so that drop measurements at ten-second intervals after melting were possible with ultra-high-speed self-develop ing film. 3) Procedure. The iron cylinders were placed upright on the MgO plates and positioned in the molybdenum susceptor. After furnace assembly the system was pumped down to a vacuum of 1 µ and flushed with H2 at 800°C. The system was again evacuated to 5 x 10-5 mm of mercury and the temperature raised to 1550°C within 2 min. The molten drops were photographed at appropriate time intervals. TWO min after melting the iron vapor pressure caused gas discharge, or corona, within the tube. The specimens were then slowly cooled to room temperature, sectioned and examined metallographi-cally. X-ray identification of interfacial reaction products was attempted by the powder diffraction technique. Alloying elements (Ti, V, Cr, Zr, Ta, Nb) were added to the iron in various concentrations up

Jan 1, 1963

By A. E. Jenkins, L. A. Baker, N. A. Warner

The electromagnetic levitation technique has been successfully applied to rate studies of the de-carburization of liquid Fe-C alloys from 5.5 to zero pct C at 1660°C using gas mixtures containing 1 to 100 pct CO2 with CO or helium as diluents. The rate was shown to be independent of carbon concen -tration in the melt and only slightly affected by total pressure. The results were shown to be consistent with control either wholly or predominantly by gaseous diffusion. As well as avoiding side reactions with the crucible, levitation allowed precise calculation of reaction rates since the geometry of the system was accurately known. It also enabled interpretation using the established mass-transfer correlations for a gas flowing past a single sphere, The existence of gaseous-diffusion control and the extremely high rates of reaction reported for gas flowing past small spheres are of great practical significance in relation to current investigations into the use of spray-refining techniques in continzcous steelmaking. PREVIOUS work on the kinetics of carbon removal from liquid iron has generally involved the use of oxygen as the gaseous oxidant. Investigations have been made on commercial steelmaking furnaces and in addition laboratory-scale experiments have been attempted with a slag phase present. The general lack of agreement and the obvious difficulty in interpreting the results is at least partly due to the complexity of the systems used. This work has been reviewed by carter1 and more recently by Bods-worth2 and ward.3 The gas-liquid metal reactions have also been studied with slag-free melts by the conventional technique of blowing oxygen onto the melt contained in a crucible. The reaction between the oxidant and carbon could be followed by sampling either the melt or the gas phase. Side reactions with the crucible material have precluded direct interpretation of the results. Fujii, 5-7 using O-A mixtures, attempted to correct for the crucible reaction and reported that the reaction between carban and oxygen occurred at the melt surface and was controlled by gaseous diffusion for carbon contents greater than a critical value and by carbon-diffusion control in the metal for carbon contents less than the critical. The critical value was given as approximately 0.15 pct C. However, whereas this was the case for oxygen contents in the gas stream up to 10 pct, between 10 and 20 pct the rate was found to be approximately independent of oxygen content, and Fujii attempted to explain this observation by postulating the presence of an oxide "film".7 Turkdogan et a1.' have studied iron fuming during decarburizing and suggested gaseous-diffusion control of the reaction above a critical carbon concentration, given as approximately 2 pct C. Filippov9-12 reports in a similar manner to Fujii but makes no mention of the crucible reaction with the carbon in the melt, even though this would be expected with the temperature, crucible material (MgO), and carbon contents used. Both CO2 and O2 were used as oxidant gases. In the present work the decarburization of liquid Fe-C alloys was studied using an adaptation of the electromagnetic levitation technique, and thus the complexities introduced by crucible side reactions were eliminated. In this technique a single sphere of the Fe-C alloy is freely suspended and heated by electromagnetic induction in a flowing stream containing the oxidant gas. The full range of carbon levels up to saturation was studied, the temperature of the melts in this instance being held constant at 1660°C. The oxidant was CO2 which was used either pure or in association with a diluent gas, the total gas pressure being maintained either at 1 atm or at a reduced pressure. Thus the reaction under study was:

Jan 1, 1964

By R. D. Pehlke, P. D. Goodell, R. W. Dunlap

The rate of dissolution of steel bars in molten pig iron has been measured experimentally in the temperature range 2300° to 2650° F. The rate of solution is shown to be a .function of bath composition, temperature, and stirring. A kinetic model based on carbon diffusion in the liquid phase has been derived to fit the experimental results. THE rate of scrap melting has long been an important variable in steelmaking operations. With the advent of the oxygen steelmaking process and the accompanying shorter heat times, the rate of scrap melting has now become one of the rate-limiting factors in steel production. As observed in commercial practice, the solution rate is influenced by the compositions of liquid and solid, the temperature, agitation, and time. However, no definitive work has been done on the Fe-C system, and there is very little information in the literature regarding the relative effects of these variables in steelmaking systems. A number of questions have been raised in regard to scrap utilization in basic oxygen steelmaking operations. Consideration has been given to the optimum size and shape of scrap, and to the use of preheated scrap as a means for decreasing the pig-iron requirement in oxygen blowing. The determination of an optimum scrap practice for a specific installation depends to a large extent upon the economics of the scrap market and also upon the behavior of scrap in the vessel. The present research was undertaken as a preliminary study in evaluating the behavior of steel in a pig-iron bath under various conditions of temperature, composition, and agitation, as might be encountered in oxygen converter operations or in any steelmaking operation where scrap behavior is an important process variable. Related studies have been carried out on non-ferrous systems. The solution rate of solid aluminum in a molten A1-Si alloy has been studied.' Furthermore, the increasing use of liquid metals has created considerable interest in studies related to dissolution of a solid in a liquid, or mass transfer taking place between a solid interface and a liquid metallic phase.2-10 In an effort to clarify the relative importance of factors influencing the dissolution of scrap in Fe-C alloys, this paper presents the results of a study of the rate of dissolution of a low-carbon steel cylinder in a molten Fe-C bath at various bath compositions, temperatures, and conditions of agitation. An attempt has been made to determine the mechanism of solution, and a model has been derived to fit the experimental results. The rate of heat transfer between the molten pig-iron bath and the solid-steel cylinder has also been studied. EXPERIMENTAL PROCEDURE A molten bath of pig iron (nominal composition 4.2 pct C, 0.5 pct Mn, 0.8 pct Si)* was held in a 200- *In view of the fact that the composition of the pig-iron bath was slowly changing with time because of steel dissolution and reaction with the surrounding atmosphere, intermittent samples of the liquid bath were taken throughout each experiment, Table 1. lb induction melting unit. The internal diameter of the furnace was 8-1/2 in. and the bath depth approximately 14 in. Cold-finished 1020 steel cylinders of 1/2-, 3/4-, 1-, 1-1/2-, and 2-in. diameters were cut into 1 -ft lengths for use as test specimens. One end of each bar was machined to fit a hand-driven, mechanical stirring device which rested on top of the furnace. This fixture permitted 7 to 8 in. of the bar to be immersed in the melt. The steel cylinders were cleaned to free the surface of grease and oxide. The rods, at ambient temperature, were immersed into the molten pig-iron bath. The melt temperatures studied in this investigation were 2300°, 2500°, and 2650°F. Different agitation conditions were achieved by operating with a) the power to the furnace, giving induction stirring; b) induction stirring plus mechanical stirring, using hand rotation with a chain and sprocket assembly at approximately 200 rpm; c) mechanical stirring alone with the power off; and d) the power off and no mechanical stirring resulting in minimum agitation, i.e., only that caused by natural convection currents in the bath. The samples were immersed in the melt for prescribed times ranging from 30 sec up to 6 min. Immersion times were measured with a stopwatch and temperature control of the melt was achieved by power adjustment following temperature measurement with an optical pyrometer. The measurements with the calibrated pyrometer were checked out to within less than 10°F with simultaneous thermocouple measurements. Following immersion, the bars were water-quenched and the diameters were measured with a micrometer at several positions on the reduced area, and by volumetric displacement. The dissolution rates calculated from the

Jan 1, 1965

By N. A. Gokcen, M. N. Dastur

In metallurgical process industries a knowledge of true melting and casting temperatures is very essential for increasing the operating efficiency as well as improving the quality of the finished product. Optical pyrometers are widely used for the purpose because of their cheapness and ease of operation. It is an unfortunate fact, however, that this mode of measuring temperature entails a source of severe error inasmuch as with all non-black bodies, the temperature observed is always below the true temperature owing to deficient emissive power. Hence, the need for emissivity correction arises. In 1913 Bidwell1 obtained emissivity values of 0.36-0.48 for a temperature range of 1520-1800°C. He measured the true temperature of molten iron by sighting into a carbon cavity immersed in the iron under an atmosphere of hydrogen. In 1914 Burgess2 determined the emissivity of a smooth surface of liquid iron free from oxide to be 0.37 under an atmosphere of hydrogen. He claimed to have attained an accuracy of within 5°C at 1500°C with an optical pyrometer using monochromatic light. He reported that no difference appeared to exist between the emissivities of pure iron and of steels containing considerable percentages of carbon, nickel or manganese; nor was there any appreciable variation of emissivity with temperature from 1530 to 1571°C. Wensel and Roeser3 found an elnissivity value of 0.4 for temperatures above 1375°C for cast iron. The average emissivity value for carbon steels reported by Leiber,4 odd,5 Hase6 and Umino7 range from 0.4 to 0.45. Knowles and Sarjant8 made an extensive study of emissivity of molten iron and steel over a wide range of workshop and laboratory conditions. Observations of true temperature were made with immersion thermocouples, and correlated with apparent temperatures indicated by optical pyrorneter readings taken on the pouring stream as it passed over the lip of the crucible. The emissivity of Armco iron was shown to be of the order of 0.4, rising slightly with increasing temperature. But for plain carbon steels containing up to 1.0 pct carbon the emissivity was shown to fall slightly with increasing temperatures. The results of their investigation are in agreement with those of Guthmann.9 Naeser,10 Spencer," Todd5 and Hase6 in regard to the relation between emissivity and temperature for plain carbon steels. However, Goller12 has found a rise of emissivity with increasing temperature for the same steels. Results of several investigations are given in Fig 1. Measurement of True and Apparent Temperatures The necessity for an accurate optical temperature scale for liquid iron in studies of gas-metal equilibria led to this investigation. Experimental measurements were carried out in the induction furnace used by Dastur and Chipman13 in their studies of the reaction of hydrogen with oxygen in liquid iron. The furnace as modified for this work is shown in Fig 2. The atmosphere of the furnace was a mixture of purified argon with hydrogen containing a small controlled amount of water vapor. Gases were preheated by means of a platinum-10 pct rhodium resistance coil. The stream of hydrogen sweeping across the surface of niolten iron vir-

Jan 1, 1950

By B. M. Larsen, T. E. Brower, J. W. Bain

A mass of circumstantial evidence is presented to indicate that the main source of alloy losses in open-hearth tapping is oxidation by air, with the steel apparently reacting with an amount of oxygen equivalent to about 30 times its own volume of air. The effect is erratic from heat to heat, depending largely on turbulence and distance of free fall of the stream of liquid metal. THIS paper reports another phase in a general study of oxidation in the open-hearth, various aspects of which have already been discussed in previous papers.1,2,3 It is based on experiments conducted on commercial, basic open-hearth furnaces at various times when opportunity offered, beginning about 1938, in an attempt to determine the mechanism and measure the amount of air oxidation during the flow of the tapping stream through the air into a ladle. This effect is closely connected, not only with alloy losses and control of chemical analysis, but possibly with various questions regarding inclusion content and steel quality and with a true picture of the deoxidation of steel in general. Thus it is believed that discussion of a large part of the results to date may be of assistance to others who are studying problems related to deoxidation or to steel cleanliness and quality. The literature reveals that strangely little attention has been devoted specifically to the effect of pouring steel through open air. Bardenheuer and Henke4 in 1939, in a paper devoted largely to overall losses of manganese in the basic open-hearth, reported that for a number of heats tapped into a tilted ladle held just beneath the runner so that the maximum distance of drop of the stream through the air was held to less than about 30 in., there was in most cases almost no loss of manganese from furnace to ladle. Hultgren,5 in 1945, showed that in high-carbon basic electric steel deoxidized in the furnace and containing no large inclusions as it left the furnace, large silicate inclusions, often rich in manganese oxide, were present in samples dipped from the pool in the ladle; also that the rise in content of large inclusions (above 0.01 mm dim) was greater when the metal stream was made to spray or flow in a turbulent manner. Our interest in this phase of oxidation in the open-hearth was first aroused by a study of alloy efficiencies on ladle additions. The frequency curves of manganese efficiency (net recovery of metal added in ladle) given in fig. 1 are of interest in this connection. Lowest, and also most variable (47 to 95 pct), recovery was in low-carbon heats having no coal additions to the ladle, progressing upward to a maximum average and least variable (82 to 97 pct) recovery in high-carbon heats killed with aluminum and silicon. Both coal additions and higher carbon contents in the tap stream, as well as additions of silicon and aluminum, help to lower the amount of manganese loss. But the most interesting point is the amount of manganese oxidized in many heats tapped at 0.50 to 1.06 pct carbon. It has been shown in a previous paper3 that the amount of oxygen in the tap stream in such heats is very small, usually below 0.01 pct, too small to oxidize any appreciable amount of manganese. Also, the fact that manganese should be oxidized at all in a metal solution containing such more stable oxide-forming elements as aluminum and silicon is an indication of some rather intense oxidizing effect such as could be caused by reaction with the air during ladle filling. Oxygen Balances from Tap Stream to Teeming Stream: Data in the form of alloy efficiencies, such as those shown in fig. 1, are not a proper measure of such a general oxidation effect. Having available

Jan 1, 1951

By W. M. McKewan

Magnetite pellets were reduced in flowing hydrogen at pressures up to 40 atm over a temperature range of 350° to 500°C. The rate of weight loss of oxygen per unit area of the reaction surface was found to be constant with time at each temperature and pressure. The reaction rate was found to be directly proportional to hydrogen pressure up to 1 atm and to approach a maximum rate at high pressures. The results can be explained by considering the reaction surface to be sparsely occupied by adsorbed hydrogen at low pressures and saturated at high pressures. PREVIOUS investigation1,2 have shown that the reduction of iron oxides in hydrogen is controlled at the reaction interface. Under fixed conditions of temperature, hydrogen pressure, and gas composition, the reduction rate is constant with time, per unit surface area of residual oxide, and is directly proportional to the hydrogen pressure up to one atmosphere. The reduction rate of a sphere of iron oxide can be described3 by the following equation which takes into account the changing reaction surface area: where ro and do are the initial radius and density of the sphere; t is time; R is the fractional reduction; and R, is the reduction rate constant with units mass per area per time. The quantityis actually the fractional thickness of the reduced layer in terms of fractional reduction R. It was found in a previous investigation2 of the reduction of magnetite pellets in H2-H,O-N, mixtures, that the reaction rate was directly proportional to the hydrogen partial pressure up to 1 atm at a constant ratio of water vapor to hydrogen. Water vapor poisoned the oxide surface by an oxidizing reaction and markedly slowed the reduction. The enthalpy of activation was found to be + 13,600 cal per mole. It was also found that the magnetite reduced to meta-stable wüstite before proceeding to iron metal. The following equation was derived from absolute reaction-rate theory4,8 to expfain the experimental data: where Ro is the reduction rate in mg cm-2 min-'; KO contains the conversion units; Ph2 and PH2O are the hydrogen and water vapor partial pressures in atmospheres; Ke is the equilibrium constant for the Fe,O,/FeO equilibrium; Kp is the equilibrium constant for the poisoning reaction of water vapor; L is the total number of active sites; k and h are Boltzmann's and Planck's constants; and AF is the free energy of activation. Tenenbaum zind Joseph5 studied the reduction of iron ore by hydrogen at pressures over 1 atm. They showed that increasing the hydrogen pressure materially increased the rate of reduction. This is in accordance with the work of Diepschlag,6 who found that the rate of reduction of iron ores by either carbon monoxide or hydrogen was much greater at higher pressures. He used pressures as high as 7 atm. In order to further understand the mechanism of the reduction of iron oxide by hydrogen it was decided to study the effect of increasing the hydrogen pressure on rebduction rates of magnetite pellets. EXPERIMENTAL PROCEDURE The dense magnetite pellets used in these experiments were made in the following manner. Reagent-grade ferric oxide was moistened with water and hand-rolled into spherical pellets. The pellets were heated slowly to 550°C in an atmosphere of 10 pct H2-90 pct CO, and held for 1 hr. They were then heated slowly to 1370°C in an atmosphere of 2 pct H2-98 pct CO, then cooled slowly in the same atmosphere. The sintered pellets were crystalline magnetite with an apparent density of about 4.9 gm per cm3. They were about 0.9 cm in diam. The porosity of the pellets, which was discontinuous in nature, was akrout 6 pct. The pellets were suspended from a quartz spring balance in a vertical tube furnace. The equipment is shown in Fig. 1. Essentially the furnace consists of a 12-in. OD stainless steel outer shell and a 3-in. ID inconel inner shell. The kanthal wound 22 in. long, 1 1/2, in. ID alumina reaction tube is inside the inconel inner shell. Prepurified hydrogen sweeps the reaction tube to remove the water vapor formed during the reaction. The hydrogen is static in the rest of the furnace. The sample is placed at the bottom of the furnace in a nickel wire mesh basket suspended by nickel wire from the quartz spring. The furnace is then sealed, evacuated, and refilled with argon several times to remove all traces of oxygen. It is then evacuated, filled with

Jan 1, 1962

By N. A. Gokcen, John Chipman

SILICON is the most commonly used deoxidizer and an important alloying element in steelmak-ing; hence a detailed study of this element in liquid iron containing oxygen is of considerable interest. The equilibrium between silicon and oxygen in liquid iron has been studied by a number of investigators but generally with inconclusive or incomplete results. The variation of the activity coefficients of silicon and oxygen with composition is entirely unknown. Published investigations deal with the reaction of dissolved oxygen with silicon in liquid iron and the results are expressed in terms of a deoxidation product. For consistency and convenience in comparison of the published information, the deoxidation product as referred to the following reaction is expressed in terms of the percentage by weight of silicon and oxygen in the melt in equilibrium with solid silica: SiO (s) = Si + 2 O; K'l = [% Si] [% 012 [I] Theoretical attempts to calculate the deoxidation constant for silicon in liquid iron from the free energies of various reactions yielded results which were invariably lower than the experimental values. Thus, the deoxidation "constants" calculated by McCance,1,2 Feild,3 Schenck, and Chipman were of the order of 10, which is below the experimental values by a factor of more than 10. Experiments of Herty and coworkers" in the laboratory and steel plant resulted in an average deoxidation constant of 0.82x10 ' at about 1600°C. The technique employed in their investigation was crude and the reported temperature was quite uncertain. The concentration of silicon was obtained by subtracting silicon in the inclusions from the total. Since at least some of the inclusions resulting from chilling must represent a fraction of the silicon in solution at high temperatures, such a subtraction is not justifiable. Results of Schenck4 for K'1 from acid open-hearth plant data yielded a value of 2.8x10-5, which was later revised as 1.24x10 at 1600°C. Similarly Schenck and Bruggemann7 obtained 1.76x10-5 at 1600OC. The discrepancies and errors involved in the acid open-hearth plant data as compared with the results of more reliable laboratory techniques were attributed by these authors to the lack of equilibrium and the impurities in liquid metal and slag, and are sufficiently discussed elsewhere." Korber and Oelsen" investigated the relation between dissolved oxygen and silicon in liquid iron covered with silica-saturated slags containing varying concentrations of MnO and FeO. The deoxidation products obtained by their method scatter considerably, and their chosen average values of 1.34x10, 3.6x10-5, and 10.6x10-5 1550°, 1600°, and 1650°C, respectively, represent the best experimental results which were available until quite recently. Darken's10 plant data from a steel bath agree approximately with their data at 1575° to 1625°C. Zapffe and Sims" investigated the reaction of H2O and H2 with liquid iron containing less than 1 pct Si and obtained deoxidation products varying by a factor of more than 20. Inadequate gas-metal contact and lack of stirring in the metal bath should require a longer period of time than the 1 to 5.5 hr which they allowed for the attainment of equilibrium. Furthermore, their oxygen analyses were incomplete and irregular and confined to a few unsatisfactory preliminary samples. Their results did indeed indicate that the activity coefficient of oxygen is decreased by the presence of silicon, although they made no such simple statement. They chose to attempt to account for their anomalous data by the unlikely hypothesis that SiO is dissolved in the melt. Hilty and Crafts" investigated the reaction of liquid iron with acid slags under an atmosphere of argon, making careful determinations of silicon and oxygen contents at several temperatures. Despite erroneous interpretation of the data at very low silicon concentrations, their data represent the most dependable information on this equilibrium that has been published. In the range 0.1 to 1.0 pct Si, their data yield the following values for the deoxidation product: 1.6x10-5, 3.0x10- ', and 5.3x10 at 1550°, 1600°, and 1650°C, respectively. The purpose of the work described herein was to study the equilibrium represented by eq 1 as well as the following reactions, all in the presence of solid silica: SiO2 (s) + 2H2 (g) = Si + 2H2O (g);

Jan 1, 1953

By N. A. Gokcen, J. Chipman

A revised treatment of the authors' published data eliminates the complex relation previously proposed between concentration of silicon and activity coefficient of oxygen in liquid iron. Revised values of the thermodynamic properties of the liquid solution are presented. IN a recent experimental study of the reaction SiO2 (s) = Si+ 2O; Kf, = [% Si] [% O]² [1] the authors' found a substantially constant equilibrium product in liquid iron at 1600°C of 2.8x10-5 They also reported extensive data on the reactions: SiO2 (s) + 2H2 (8) = Si- + 2H2O (g); K'2= [% Si] (H2O/H2 [2] and H, (g) + 0 = H2O (g); K'3 =( H2 O [31 (H2) [%O] From the results on reaction 3 and earlier data of Dastur² on this same reaction in the absence of silicon, they determined the activity coefficient of oxygen, f0, on the basis of the definition K3 = (H2O)/ (H2)f0 [% 0] where K, is the equilibrium constant and f0, is taken as unity in the pure Fe-0 system. Similarly values of fsi were deduced from results on reaction 2. In a more recent study" of analogous reactions in the system Fe-A1-0, it was found impossible to reconcile the results on reaction 3 with Dastur's data; accordingly the latter were ignored and the equilibrium results were extrapolated to find a value of K, at zero concentration of aluminum. This procedure failed to locate the cause of the discrepancy but it did yield reasonable values of activity coefficients. It also avoided introduction of the complex empirical relation between the oxygen activity coefficient and the concentration of the added element. The same type of discrepancy exists for system Fe-Si-0.' In the earlier paper an attempt was made to fit both sets of data by a single curved line (Fig. 6 of ref. l), the form of which is contrary to the theoretical requirement of a finite slope at infinite dilution. In the light of experience on the Fe-A1-0 system the discrepancy must be recognized as one which can be resolved only by more refined measurements. Accordingly Figs. 6 and 10 are retracted. It is pointed out also that until the discrepancy is resolved Figs. 7, 8, and 11 are subject to some uncertainty. Qualitatively the following conclusions still appear valid: 1—The activity coefficient of oxygen is reduced by addition of silicon. 2—In dilute solutions the activity coefficient of silicon increases with its concentration. 3—With respect to equilibrium in reaction 1, the above effects are approximately compensating. The discussion of K'1 in the previous paper requires no revision. It was pointed out that the constancy of the product [% Si] [% 0]² ndicated a compensating effect of the activity coefficients of silicon and oxygen. Therefore, as a very good approximation, K1 = K'1 and the following average values are suggested both for K, and K', at the temperatures 1550°, 1600°, and 1650", respectively, 1.0x10-", 2.8~10-" and 5.5 ~lo-'. Revision of the thermodynamic treatment is necessitated by the recent appearance of new data, based on a combination of combustion and solution calorimetry,' which yields for the heat of formation of low-cristobalite from the elements, the value —209,330 ±250 cal per mol at 25°C. This is about 4000 cal larger than the value previously accepted. The new value for cristobalite is used, together with Kelley's tables of high-temperature heat contents" and entropies and with Korber and Oelsen's' heat of fusion of silicon to obtain the following equation for the standard free energy of cristobalite in the temperature range 1700" to 2000°K: Si (1) + 02 (g) = Si02 (crist.); ?F° = -217,700 + 47.OT [4] The free energy of solution of 0, in liquid iron is:8 O2 (g) = 20 (in Fe); AF° = -55,860 - 1.14T [5] and these two equations are combined to give: Si (1) + 20 = SiO, (crist.); AF° = -161,840 + 48.14T [6] ?F°1873 = -71,700 cal. From the experimental value of K, = 2.8x10-5, Si + 20 = SiO2 (crist.); ?F°1873 = -39,000 cal. [7] The combination of Eqs. 6 and 7 yields the free energy change when liquid silicon dissolves in iron to form the dilute solution of unit activity (1 pct). Si (1) = Si; ?F°1873 = -32,700 cal. [8] The heat effect in this process according to Korber and Oelsen' is an evolution of 28,500 cal per gram

Jan 1, 1954

By J. O. Edstrom, G. Bitsianes

Parabolic rate constants were determined for the formation of wiistite by the solid state reaction between magnetite and iron. The reaction was diffusion controlled and inert marker studies indicated that the mass transport through the wiistite layer was accomplished by means of iron migration. Relationships between rate constants and self-diffusivities are discussed. The transport capacity for iron through dense wustite layers was found to be sufficient to carry on reduction, even in gaseous reduction processes. REDUCTION of iron oxides has been the subject of many investigations. Most of the work, however, has been done on rather crude material and there have been difficulties in correlating the data on a quantitative basis. The reverse process of the oxidation of iron has been studied more thoroughly and usually by starting from the pure metal. As a result, the mechanisms of oxidation are now fairly well known.It has generally been found that the oxidation of metals follows the so-called parabolic 1aw;Q hat is, when the oxides are formed as dense layers, their thicknesses are proportional to the square root of the reaction time. This parabolic law follows directly from Fick's first law of diffusion The flux, J, is defined as the quantity of diffusing substance passing per unit time through unit area of a plane at right angles to the direction of diffusion. This flux is proportional to the concentration gradient of the diffusing substance, and the factor D, known as the diffusion coefficient, is introduced as the proportionality factor with dimensions of (length)'/time. If the flux is measured as the increase of thickness of the oxide layer (Ax) per unit time (t) and the assumption is made that the concentrations of the reactants at the phase boundaries of the layer are independent of time, Eq. 1 reduces to Integration of this expression yields the parabolic function ax = k . t. [31 The product of the diffusion coefficient and the concentration difference across the oxide layer is included in the constant k. When iron is exposed to a highly oxidizing atmosphere, the oxide phases are formed normally in a topochemical fashion, i.e., the interfaces between the phases maintain parallel positions to the original surface of the specimen. Above 570°C, the phases are orientated usually in the order of iron, wiistite, magnetite, and hematite, a condition in conformance with the requirements of the Fe-0 system. Below 570°C, the wiistite phase is unstable and has not been found as a normal constituent in the oxidation products of iron. There is a strong probability that continuous layers of the solid oxides are formed as the specific volumes of the phases increase in direct order from iron to hematite.' , ' Regarding mass transport through dense solid oxide layers, Wagner' has postulated that diffusion in oxides generally may be interpreted as migration processes of ions and electrons. There must be normally a concentration gradient across the growing oxide layer for diffusion to occur. This condition requires that there be deviations from the ideal stoi-chiometric composition of the oxide and accordingly deviations from the strict order of an ideal lattice. Such "lattice defects" include interstitial ions, cation and anion vacancies, quasi-free electrons, and electron holes and are decisive for all migration processes.' Cations, anions, and electrons all may have some mobility in the oxides but the movement of one of the particles generally far exceeds that of the others. Previous work', 1,2,7 has shown that during the oxidation of iron the migration through the wiistite layer, and probably also the magnetite layer, is confined to the movement of iron ions. Migration through hematite, however, has been found to take place by oxide ion movement. These behaviors are in direct agreement with the type of vacancies present in the respective oxides. Reduction of Iron Oxides The gaseous reduction of iron oxides, like the oxidation of iron, takes place in a topochemical fashion at distinct interfaces between the appearing phases."-'" Porosities might be expected in these reaction products, since their specific volumes are less than that of the starting material. Frank and van Der Merwe" have shown, however, that it is theoretically possible for nonporous layers to form if the degree of misfit in specific volume is less than about 15 pct. These aspects on porosity formation

Jan 1, 1956

By John Henderson

FOR some time now the most commonly accepted description of liquid silicate structure has been the "discrete ion" theory, proposed originally by Bockris and owe.' This theory is that when certain metal oxides and silica are melted together, the continuous three dimensional silica lattice is broken down into large anionic groups, such as sheets, chains, and rings, to form a liquid containing these complex anions and simple cations. Each composition is characterized by "an equilibrium mixture of two or more of the discrete ions",' and increasing metal oxide content causes a decrease in ion size. The implication is, and this implication has received tacit approval from subsequent workers, that these anions are rigid structures and that once formed they are quite stable. The discrete ion theory has been found to fit the results of the great majority of structural studies, but in a few areas it is not entirely satisfactory. For example it does not explain clearly the effect of temperature on melt structure,3 nor does it allow for free oxygen ions over wide composition ranges, the occurrence of which has been postulated to explain sulfur4 and water5 solubility in liquid silicates. In lime-silica-alumina melts the discrete ion theory is even less satisfactory, and in particular the apparent difference in the mechanism of transport of calcium in electrical conduction8 and self-diffusion,' and the mechanism of the self-diffusion of oxygen8 are very difficult to explain on this basis. By looking at melt structure in a slightly different way, however, a model emerges that does not pose these problems. It has been suggested5" that at each composition in a liquid silicate, there is a distribution of anion sizes; thus the dominant anionic species might be Si3,O9 but as well as these anions the melt may contain say sis0:i anions. Decreasing silica content and increasing temperature are said9 to reduce the size of the dominant species. Taking this concept further, it is now suggested that these complexes are not the rigid, stable entities originally envisaged, but rather that they exist on a time-average basis. In this way large groups are continually decaying to smaller groups and small groups reforming to larger groups. The most complete transport data 8-10 available are for a melt containing 40 wt pct CaO, 40 wt pct SiO2, and 20 wt pct Al2O3. Recalculating this composition in terms of ion fractions and bearing in mind the relative sizes of the constituent ions, Table I, it seems reasonable to regard this liquid as almost close packed oxygens, containing the other ions interstitially, in which regions of local order exist. On this basis, all oxygen positions are equivalent and, since an oxygen is always adjacent to other oxygens, its diffusion occurs by successive small movements, in a cooperative manner, in accord with modern liquid theories." Silicon diffusion is much less favorable, firstly because there are fewer positions into which it can move and secondly, because it has the rather rigid restriction that it always tends to be co-ordinated with four oxygens. Silicon self-diffusion is therefore probably best regarded as being effected by the decay and reformation of anionic groups or, in other words, by the redistribution of regions of local order. Calcium self-diffusion should occur more readily than silicon, because its co-ordination requirements are not as stringent, but not as readily as oxygen, because there are fewer positions into which it can move. There is the further restriction that electrical neutrality must be maintained, hence calcium diffusion should be regarded as the process providing for electrical neutrality in the redistribution of regions of local order. That is, silicon and calcium self-diffusion occur, basically, by the same process. Aluminum self-diffusivity should be somewhere between calcium and silicon because, for reasons discussed elsewhere,' part of the aluminum is equivalent to calcium and part equivalent to silicon. Consider now self-diffusion as a rate process. The simplest equation is: D = Do exp (-E/RT) [I] This equation can be restated in much more explicit forms but neither the accuracy of the available data, nor the present state of knowledge of rate theory as applied to liquids justifies any degree of sophistication. Nevertheless the terms of Eq. [I] do have significance;12 Do is related, however loose this relationship may be, to the frequency with which reacting species are in favorable positions to diffuse, and E is an indication of the energy barrier that must be overcome to allow diffusion to proceed. For the 40 wt pct CaO, 40 wt pct SiO2, 20 wt pct Al2O3, melt, the apparent activation energies for self-diffusion of calcium, silicon, and aluminum are not significantly different from 70 kcal per mole of diffusate,' in agreement with the postulate that these elements diffuse by the same process. For oxygen self-diffusion E is about 85 kcal per mole,' again in agreement with the idea that oxygen is transported,

Jan 1, 1963

By J. M. Dahl, R. J. Warrick, O. K. Riegger, H. Van Vlack

A liquid which is rich in oxygen (and silicon) develops at steel rolling temperatures in resulfurized and plain-carbon steels. This liquid fluxes solid manganese sulfide. The composition of the liquid and the resulting microstructures vary with the manganese/silicon/oxygen ratios in the steel. The observations and conclusions zuhich are presented are used 1) to suggest a mechanism for MnS deformation in resulfurized steels, and 2) to rationulize hot-shortness anomalies between plain-carbon and resulfurized steels. 1 HIS is a summary of the dependence of sulfide inclusion microstructures upon steel composition, and an interpretation of the behavior of sulfide inclusions at steel-rolling temperatures. Sulfide inclusions are known to have several metallurgical effects upon steel. Their effects upon hot-shortness are best appreciated. Although hot-shortness is well known, it is questionable whether its mechanism is fully understood because there are several anomalies in our interpretation. For example, why should not high sulfur contents produce extensive hot-shortness in free-machining steels ? A second metallurgical consequence of sulfide inclusions appears because sulfide inclusions improve machinability. Furthermore, globular inclu- sions are to be desired over elongated inclusions. To date, we know that lower silicon contents (< 0.010 pct) favor the globular sulfides, but we do not know why this interrelationship between silicon content and inclusion shape exists. Finally, the evidence is strong that the surface quality of a hot-rolled steel is particularly sensitive to the sulfur content. SUMMARY OF PREVIOUS WORK The role of sulfur in hot-shortness was first suggested during the past century. It is uncertain who was the first to conclude that sulfur was less deleterious in the presence of manganese, and that manganese altered the sulfide phase from a liquid FeS to a solid MnS at steel-rolling temperatures. However, Urban and chipman&apos; emphasized that liquid FeS may accumulate as an interdendritic film during solidification. More recently, Keh and Van lack&apos; showed that the intergranular liquid does not completely surround the metal grains at all rolling temperatures. Rather, a small finite dihedral angle exists when geometric equilibration is attained at

Jan 1, 1962

By J. Chipman, G. G. Hatch

T. ROSENQVIST*—It is a pleasure to see the excellent way in which the experimental part of this work has been handled. There seems to be little doubt that the distribution data obtained corresponds most closely to thermodynamic equilibrium under the prevailing reducing conditions, namely equilibrium with graphite and one atmosphere CO pressure. The desulphurization curves in Fig 10 show the same general feature as the curves given by Holbrook and Joseph, but the distribution ratios are from 20 to 40 times greater—undoubtedly due to a closer approach to true equilibrium. In the theoretical discussion, the authors calculate a theoretical distribution (S) ration -jg-. which they find to be about 50 times greater than the experimental. The deviation is so great that the basis for their calculation needs a more thorough examination. The authors base their thermodynamic calculation on free energy expressions where diluted solutions of FeS and CaS are used as standard states. (The activity coefficient in diluted solutions is taken to equal unity.) Such a standard state will change when the nature of the solvent is changed. Taking the free energy of the reaction [FeS] ? (FeS), Eq 2, which is derived from the distribution of sulphur between an iron and a FeO-melt, it is very unlikely that the free energy of this reaction will be the same for a distribution between pig iron and a calcium silicate slag. Therefore a more fundamental basis for the thermodyuamic calculations seems needed, where all thermodynamic equations are referred to unambiguously defined standard states. The most natural standard states for CaO and CaS are the pure solid substances at the same temperature. As standard state for sulphur in iron, pure liquid FeS can be used. This rules out Eq 2 [FeS] ;=s (FeS) because ?F° = 0. The standard equation will then be: FeS, + CaO6 + Cgraph ?Fei + CaS8 + CO. vFo1773 = 25,000 cal It would be more universal and also simpler to refer the escaping tendency of sulphur in liquid iron to the corresponding H2S/H2 ratio which can readily be determined experimentally. As standard state a gas mixture H2S/H2 = 1/1 can be used. (This corresponds at the temperature of liquid iron closely to one atmosphere S2 vapor.) Thus the standard equation for the sulphur reaction can be formulated as follows: H2S0 + CaO3 + Cgraph ?H2o + CaS8 + COg The standard free energy of this reaction has been calculated from the best available data to AF°m3 = —35,000 cal. This gives for the equilibrium constant at 1500°C Now, the solubility of CaS in blast furnace slags has been determined by McCafferey and Oesterle* and corresponds at 1500°C to about 10 pet S (varying somewhat with the composition of the slag.) If the activity of CaS is assumed linear between 0-10 pet as curve 1, (see Fig 11), then acaO = 0.1 (S); (S) being wt. pet sulphur in the slag. For a diluted solution of sulphur in an iron melt saturated with carbon, the ratio H2S/H2 is, according to Kitchener, Bockris and Liberman,f about 0.01 [S], [S] being wt. pet sulphur in iron. Substituting these values in the expression for Kp we find The value 2.103 is only 4 times greater than the experimental coefficient found by Hatch and Chipman, but the value is very sensitive to a small error in AF°. A better agreement with the experimental distribution coefficient can be obtained if one assumes the activity of CaS to run like curve 2 (Fig 11). This (S) will give a lower theoretical W, value, a value which varies with (S) exactly as Hatch and Chipman learned. Such a shape of the activity curve, which corresponds to a positive deviation from Raoult&apos;s law, is actually to be expected from the fact that liquid silicate and sulphide phases usually show incomplete miscibility. A closer agreement between experimental and theoretical data can not be expected before we have more complete data for the individual activities of CaS and CaO in the slag. The activities acaS and Ocao referred to the solid phases as standard states, are exact defined quantities contrary to the somewhat undefined expression "free lime," and they are independent of any theory for the constitution of liquid slag. J. CHIPMAN (authors&apos; reply)—The authors wish to thank Mr. Rosenqvist for his very interesting and useful thermodynamic addition. Curve 2 of his figure offers the needed basis for explaining the increase in the ratio (S)/[S] with increasing sulphur content. Attention is called to an error in the printed paper: Fig 2 and 3 are reversed. M. TENENBAUM*—In the figures showing the relationship between excess base and sulphur distribution (Fig 6, 7 and 9) the slope of the curve tapers off in the negative basicity range. Somewhat the same thing is observed with open hearth slags. In that case, the fact that some sulphur distribution between slag and metal is obtained with negative basicity is interpreted as indicating some dissociation of the lime silicate compounds whose existence in oxidizing basic slags has been used to explain various observed phenomena with regard to other slag-metal reactions. In the case of the blast furnace slags, the reduced slope of the sulphur distribution curve with decreasing excess base is attributed to the amphoteric effect of alumina. Has the possibility of other explanations been investigated ?

Jan 1, 1950

By E. N. Silverman

Mixtures of&apos; MnS, ZFeO-SiO,, ZMnO.SiO, FeO, and MnO were prepared synthetically and used to study the planes ZFeO. Si0,-ZMnO.SiO,-MnS, FeO-MnO-MnS, and Fe0-MnO.Si0,-MnS in the quaternary system. Because the oxides and silicates studied flux MnS, inclusions in this system are partly liquid during rolling. Also, oxygen can modify subsurface inclusions through the reaction MnS + FeO—FeS + MnO during the soaking of ingots and the reheating of blooms and billets. NONMETALLIC inclusions in steel have long been a subject of study and conjecture because the properties of many steels are known to be affected by the the kind and amount of inclusions present. In free-machining steels, the type of inclusion present has a direct bearing on their machinability; in other steels, cleanliness is important and metallurgists direct their efforts at reducing the amount of non-metallic inclusions. Thus, the accumulation of knowledge about the formation and behavior of inclusions during all stages of steelmaking and processing should aid in improving the quality of steels and decreasing the production costs resulting from rejections. Previous studies of inclusions have contributed greatly to our understanding. The investigators have generally examined inclusions in finished steel. Boulger, Moorehead, and Garvey1 found that in Bessemer steels, low silicon contents results in globular sulfide inclusions with attendant high machin- ability ratings. Van vlack2 also found that low silicon contents favor inclusion morphology that contributes to high machinability ratings. Boquist and Stoll* found that silicon content affected sulfide inclusions and machinability of blooms, billets and bars more than did the finishing temperature in the range between 1775° and 2075°F. Carney and Rudolphy3 studied the distribution of sulfur, manganese, silicon, and carbon in a commercial ingot of resulfur-ized steel and found that silicon was concentrated near the bottom. This helped to explain why bottom cuts had poorer machinability. Klinger and Koch4 developed techniques for extracting inclusions from steel, and the technique devised by Gurry, Chris-takos, and stricker5 for extracting carbides can also be adapted to the study of inclusions. A different approach to the study of inclusions in steel consists of synthesizing the constituents that form inclusions and determining the phase eauilibria in the system that contains the constituents. The resulting phase diagrams can then be used to determine the effect of the various constituents upon each other. In this way, it becomes possible to under-

Jan 1, 1962

By N. A. Gokcen, E. S. Tankins, G. R. Belton

Equilibrium in the reaction H2(g) + O[in liquid iron, cobalt, or nickel] = H2O(g) has been investigated over wide temperature and composition ranges. Oxygen has been found to obey Henry&apos;s Law in each of the three metals at 1550°C. The standard free-energy changes for the above reaction are given by: (1550° to 1700°C) ?Foco = -45,360(±290) + 16.16(±0.16)T (1490° to 1700°C) ?Fo Ni = -48,170(±220) + 16.12(±0.12)T (1465° to 1700°C) where the standard state for dissolved oxygen is the I wt pet solution. ThE thermodynamic behavior of dilute solutions of oxygen in molten iron has been the subject of many investigations over the last thirty years,&apos;-&apos;4 the majority of the investigators having utilized the reaction:

Jan 1, 1964

By Robert Baschwitz, John Chipman

The distribution of silicon between liquid silver and Fe-Si-C alloys has been studied at 1420oand 1530°C. The data are consistent with earlier studies. New data of Hager on the liquidus lines of the system Ag-Si and the distribution data are used to obtain the activity coefficient of silicon in both liquid phases. Data on the heat of mixing in iron permit accurate extension to 1600°C. Equilibrium data involving SiO2 and silicon in liquid iron together with revised data on the free energy of SiO2 are used to fix the activity of silicon in the infinitely dilute solution. The binary system exhibits strong negative deviation from ideality. At infinite dilution ? Si at 1600" is 1.25 x 10&apos;3, and at concentrations up to NSi = 0.4 the slope d InySi/dNSi has a constant value of r; = 13. It is found that logysi in the ternary solutzon is approximately but not exactly the same function of Nsi + NC as of NSi in the binary. The results are consistent with currently available data on the free energy of Sic and its solubility in molten iron. LIQUID solutions of the system Fe-Si-C have acquired considerable importance as the laboratory prototypes of blast furnace hot metal. Equilibrium studies involving such solutions and slags approximating those of the blast furnace have yielded useful information concerning the thermodynamic properties of blast furnace slags. In studies of this kind great importance attaches to a knowledge of the thermodynamic activity of silicon in the solution as a function of temperature and composition. An attempt was made by Chipman, Fulton, Gokcen, and askey&apos; to evaluate all of the pertinent data on this system and to deduce the desired relation between activity, composition, and temperature. These authors published data on the solubility of graphite and Sic in molten Fe-C-Si solutions and on the distribution of silicon between liquid iron and liquid silver. They showed further how the activity of silicon in very dilute solutions in liquid iron could be calculated from equilibrium data involving the molten alloy and solid SiO,. These calculations rested on the published thermodynamic properties of SiO, in- cluding its heat of formation which at that time was recorded as -209.8 kcal. This value has been under suspicion for some time and has recently been replaced by the concordant results from two independent laboratories2,3 which place the heat of formation of a-quartz at -217.6 kcal. This revision necessitates a re-evaluation not only of the activity of SiO2 in slag but also of silicon in molten iron. It is the purpose of this paper, therefore, to recalculate the activity of silicon, and in furtherance of this objective to present new data on its distribution between liquid Fe-Si-C alloys and liquid silver. HEAT OF SOLUTION OF SILICON IN IRON In order to determine the effect of temperature upon the activity coefficient it is necessary to know the heat of solution of silicon in iron as a function of composition. This is found in the data of Korber and Oelsen4 shown in Fig. 1. The curve corresponds to the following equation, which is of a form suggested by Wagner:5 Here AH is the heat absorbed in kilocalories in forming one gram atom of molten alloy from its molten elements and the N&apos;s are atom fractions. The relative partial molal enthalpies of the components, each referred to its pureliquid state and defined as zFe = aFe - PFe and zsi = HSi — -psi, are shown graphically. At low concentrations zSi = -28.5 kcal, in agreement with Kijrber and Oelsen&apos;s computation. This is in good agreement with the value of -29.3 kcal obtained by Chipman and Grant6 using an entirely different method. ACTIVITY AT INFINITE DILUTION From the known free energy of SiO, it is possible to obtain the activity of silicon in dilute solution in liquid iron from equilibrium studies. The heat of formation of a-quartz is —217.6 kcal and the heat capacity and entropy data are given by Kelley and ~ing.&apos; The free energy of formation of ß-cristo-balite at temperatures above the melting point of silicon is expressed by the following equation: Si(Z) + O2(g) = SiO2 (crist); AF" =-226,500 + 47.50T [I] The value of the deoxidation product for silicon [%Si] x [%O]2 at 1600°C according to Gokcen and chipmans is 2.8 x 10"5, in agreement with results of Hilty and Crafts.9 More recent works of Matoba,

Jan 1, 1963

By Donald B. Evans, Robert D. Pehlke

The solubility of nitrogen in liquid Fe-A1 alloys has been measured up to the solubility limit for formation of aluminum nitride using the Sieverts method. The activity coefficient of nitrogen decreases slightly with increasing aluminum content in the range of 0 to 4 wt pct Al. Based on a nitride composition, AlN, the standard free energy of formation of aluminum nitride from fhe elements dissolved in liquid iron has been determined to be: ?F" = -59,250 + 25.55 T in the range from 1600º to 1750ºC. The solubility of nitrogen in liquid iron alloys and the interaction of nitrogen with dissolved alloying elements in liquid iron have been the subject of a number of research investigations.&apos; Most of this work, however, has been reported for concentrations well below those necessary for the formation of the alloy nitride phase. Data in the concentration region near the solubility limit of the alloy nitride, particularly for systems exhibiting stable nitrides, are important in evaluating the denitrifying power of various alloying elements. They are also useful in determining the stability of a given nitride if it is to be used as a refractory to contain liquid iron alloys. In view of the importance of aluminum as a deoxidizing agent in commercial steelmaking and the fact that its nitride, AIN, is a highly stable compound and has merited some consideration as an industrial refractory, the following investigation was undertaken. The use of the Sieverts technique provided a measurement of the equilibrium nitrogen solubility in liquid Fe-A1 alloys as a function of nitrogen gas pressure up to 3.85 wt pct A1 in the temperature range of 1600º to 1750°C. The values obtained by the Sieverts method were checked by means of a quenching method in which liquid iron was equilibrated with an A1N crucible under a known partial pressure of nitrogen gas, and the solubility of A1N in liquid iron determined by chemical analysis. EXPERIMENTAL PROCEDURE The theoretical considerations involved in determining the solubility product of a solid alloy nitride phase in liquid iron by measuring the point of departure of the nitrogen gas solubility from Sieverts law have been discussed by Rao and par lee.&apos; The principal problem is to determine the variation of nitrogen solubility in an alloy as a function of the pressure of nitrogen gas over it with sufficient precision to establish the break point in the curve at the solubility limit of the alloy nitride phase. A fairly large number of data points are required to do this. A second problem is the determination of the composition of the precipitated solid nitride phase. This is necessary in order to define completely the thermodynamic relationships. The Sieverts apparatus used to make the nitrogen solubility measurements in this investigation is of essentially the same design as that described by Pehlke and E1liott.l The charge materials were Ferrovac-E high purity iron supplied by Crucible Steel Co. and 99.99+ pct pure aluminum. Recrystal-lized alumina crucibles were used, and were not attacked by the liquid alloys. The hot volume of the system which was measured for each melt ranged from 46 to 50 standard cu cm and was found to decrease linearly with decreasing pressure and with increasing temperature. The temperature coefficient of the hot volume at 1 atm pressure of argon gas was essentially constant for all experiments at a value of -6 X 10-3 cu cm per "C. The melt temperature was measured with a Leeds and Northrup disappearing filament type optical pyrometer sighted vertically downward on the center of the melt surface. The temperature scale was calibrated against the observed melting point of pure iron taken as 1536°C. The emissivity of all melts was assumed to be that of pure iron, taken as 0.43. The charge weights ranged from 110 to 140 g and the range of aluminum contents covered was from 0 to 3.85 wt pct. Aluminum additions were made as 12 to 15 wt pct A1-Fe master alloys previously prepared in the system under purified argon. The compositions of the master alloys were checked by chemical analysis and found to be in agreement with the charge analyses. Vertical cross sections of the master-alloy ingots were used as charge material for the equilibrations in order to minimize the effect of any segregation which might have occurred during solidification of the master alloys. Determinations of the solubility product of

Jan 1, 1964

By Donald B. Evans, Robert D. Pehlke

The solubility of nitrogen in liquid Fe-B alloys has been measured up to the solubility limit for the formation of boron nitride. The activity coefficient of nitrogen increases with increasing boron content in the range 0 to 7 wt pct B. From experimen -tal data, values have been calculated for the B-N interaction parameter e3 at temperatures in the range 1550" to 1750°C. A value of 0.038 has been estimated for the boron self-interaction parameter eg at 1550°C. The standard free energy of decomposition of boron nitride into the elements dissolved in liquid iron has been determined to be: ?F° = 45,900 -21.25T in the range from 1550° to 1750°C. The nitride is assumed to be of composition BN. BORON nitride has an unusual combination of properties which make it appear attractive in a wide range of engineering applications. Some of its more important and most recent applications are in the nuclear area, particularly in connection with the liquid-metal cooled reactor concept now receiving considerable emphasis. Boron nitride has a high degree of stability at elevated temperatures. It also has excellent ma-chinability and the ease with which its crystals deform suggests applications as a lubricant. These properties stem from a hexagonal layer-type structure similar to the structure of graphite. One of its primary uses to date has been for seals in liquid-metal pumping systems. It is also used in nuclear reactors as an insulating layer to separate two solid metals which are not themselves compatible under the conditions of temperature and atmosphere in which they are used. Its inertness to liquid metals has also suggested use as a mold-release agent in casting processes. In addition to its excellent machinability and reported inertness to liquid metals such as iron, silicon, aluminum, copper, and zinc, boron nitride has high thermal conductivity and excellent thermal shock resistance. This combination of properties would make it appear ideal as a refractory crucible material for refining of high-purity liquid metals, for example high-quality steels. However, since it is known that concentrations of boron as low as 50 ppm can have a marked effect on the physical properties of certain steels,&apos; in particular on the creep and stress-rupture properties, an investigation was undertaken to define accurately the chemical equilibrium among boron, nitrogen, and liquid iron in the range of steelmaking temperatures. EXPERIMENTAL PROCEDURE Two experimental approaches to this problem were employed: a Sieverts&apos; method and a quenching method. In the first method, the Sieverts&apos; technique was used to measure the equilibrium nitrogen solubility in liquid Fe-B alloys of 0 to 7 pct B as a function of nitrogen gas pressure over the melt. The solubility limit of the boron nitride phase formed was determined by the point of departure of the nitrogen absorption from Sieverts&apos; Law. This technique has been applied to liquid Fe-Ti alloys by Rao and parlee,&apos; to liquid Fe-A1 alloys by Evans and Pehlke,3 and to solid Fe-V alloys by Fountain and Chipman.4 In the second method a melt of liquid iron was held in a crucible of boron nitride under a known partial pressure of nitrogen gas. After thermodynamic equilibrium was attained, the melt was quenched in a stream of helium and then analyzed by wet-chemical methods for boron and nitrogen. The Sieverts&apos; apparatus used in the first method was essentially of the same design as the one described by Pehlke and Elliott.5 The charge materials were vacuum-melted high-purity iron (Ferro-vac E) supplied by the Crucible Steel Co. and -325 mesh boron powder supplied by Cooper Metallurgical Associates of Cleveland, Ohio. The boron contained less than 0.02 wt pct O, according to supplier&apos;s analysis. Recrystallized alumina crucibles were used to contain the melt. Examination of solidified melts showed these crucibles to be satisfactory with no evidence of any reaction or physical penetration of the crucible wall by the melt. The melt temperature was measured by a disappearing filament-type optical pyrometer sighted vertically downward on the melt surface through a 1/4-in.-diam sight hole in the crucible lid. The pyrometer was calibrated against the melting point of pure iron in the same apparatus, taking the emissivity of

Jan 1, 1964

By G. W. Healy, J. C. Schottmiller

Molten silicates were found to exist in the Cr-Si-0 system at temperatures above 1450°C. one atom of oxygen is readily removed from CrzOs in the presence of Si02 at 1700°C, forming a silicate melt. Steam-hydrogen ratio measurements made at this temperature indicate that the free energy of the reaction: Cr20s(c) + Hz = 2CrO (in Si02 -saturated melt) + H20 is approximately 5600 * 1500 cal. A deep-blue chromous silicate, probably Cr2Si04, forms from these melts at about 1400" to 1500°C, and is characterized by.its distinctive X-ray diffraction pattern and optical properties—distinct pleochroism and high refractive indices. THE question of the state of oxidation of chromium in slags has vexed metallurgists for many years The experiments of Koerber and Oelsenl suggested the probable presence of divalent chromium in silica-saturated slags equilibrated with high-chromium melts. Hilty and coworkers2 had also found chromium at lower valency, as Cr304, in inclusions formed from oxygen-saturated Fe-Cr melts with as little as 10 pct Cr. In both cases, disproportiona-tion of the lower oxide to CraO3 and chromium was found to occur to a greater or lesser extent, depending upon the rate of cooling from the high temperature. The work on Cr304 was sufficiently detailed for chipmans to derive its thermodynamic properties, but Oelsen&apos;s experiments in the silicate system were not extensive enough since most of his runs were made with slags containing large amounts of manganese oxide in addition to iron and chrome oxide and silica. The present study describes experiments carried out to map out an approximate liquidus diagram in the CrzOs, "CrO", SiOa portion of the Cr-Si-0 system, and to measure the partial free energy of solution of oxygen in the system. EXPERIMENTAL 1) Approximate Melting-Point Determinations. Preliminary trials showed that when an equimolar mixture of Cr20s and SiOa was treated with hydrogen in the temperature range 1500" to 1600°C, copious amounts of steam were evolved, and a low-melting liquid formed which could be quenched to a blue, crystalline, material having a distinctive X-ray diffraction pattern. It was found that the identical material could be produced by heating intimately mixed, finely divided Cra03, Cr, and SiOa, held in a tungsten crucible, under purified argon. Thus, it was possible to prepare mixtures containing any desired proportions of chromium, silicon, and oxygen, pel-letize them, heat to temperature in the furnace shown in Fig. 1, and observe, after cooling, whether the pellet remained unchanged or showed signs of incipient, partial, or complete melting. 2) Free Energy of Solution of Oxygen. This was studied by measurement of steam-hydrogen ratios for the reaction: H20 (gas) + O (dissolved in charge) + Ha (gas) The experimental method chosen was to prepare a

Jan 1, 1964