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By H. Eyring, J. W. Fredrickson

Many theories have been advanced to explain the phenomena of elastic and plastic deformation. The object of this report is to present a mechanism for deformation, not radically differing from existing theories, but embodying.certain principles and concepts of these theories. This proposed theory cannot be the complete answer to the age-old problem of deformation but it is believed that it will be a step further in the proper direction. The plastic properties of solids have recently been reviewed by Seitz and Read.' ZenerZ has reviewed and correlated the existing work on the anelasticity of metals. With these and others3'4'5 resent theories of deformation are well-known. Of the many theories now extant, those receiving most attention have been proffered by Griffith,6 Becker,l Smekal,8 Polanyilg Orowanlo and Taylor.'' Although differing quite widely in details, these theories are based more or less on a common principle of flaws or dislocations that exist in the crystal lattice. The effect of stress concentrations surrounding a flaw or crack has been treated most recently by Fisher and Hollomon.12 The various authors differ as to the exact nature of the flaws or dislocations, the causes of their appearance, or the manner in which they initiate slip. All of the theories agree that the number of dislocations increases with progressing deformation. It is universally agreed that the shearing process is not a simple gliding of ideally perfect atomic lattice planes. It is also agreed that the "macroscopic" process of gliding, or slip, is brought about by the spreading of a large number of locally initiated "glide steps" through the whole crystal. The occurrence of these "glide steps" is closely related to the presence of some kind of deviations from the "ideal" lattice structure, even in the undeformed crystal. These deviations are overcome ''one at a time," in a manner of speaking, due to some stress-concentrating effect.13 Proposed Theory Many solids show a continuously increasing deformation if subjected to a constant load of suitable magnitude in an extended range of temperature for a sufficient time. This deformation, or flow, may be said to be due, in general, to the tendency of the atoms or molecules in an elastically stressed body to readjust themselves in such a way that if the deformation were kept constant, a release of stress would take place. If rupture does not occur, this process of relaxation causes the rate of flow to obtain a final constant value, as shown in the curve (Fig I) obtained by Taylor and Quinney.14 In this state the release of stress by relaxation exactly counterbalances the increase due to the elastic strain on the test pieces.

Jan 1, 1949

By L. D. Jaffe, D. C. Buffum

The importance of temper brittleness in alloy steels has long been realized in Europe. In the United States recognition of its importance has developed within the last several years. Many brittle failures encountered in parts made of alloy steel are undoubtedly attributable, wholly or in part, to temper brittleness. These failures have occurred particularly at low temperatures, under shock loading, and in the presence of notches, restraints and combined stresses. It has been almost universally accepted that plain carbon steels are not susceptible to temper brittleness. As a result, considerable effort has been expended in attempts to discover the manner in which alloying elements introduce temper brittleness. It is desired to raise the possibility that plain carbon steels are not only susceptible to temper brittleness, but are much more susceptible than alloy steels, in the sense that they embrittle much more rapidly. It is suggested that previous views, (that plain carbon steels are not susceptible), are based upon an erroneous interpretation of the experimental results. Interpretation of Previous Data For many years the standard criterion for susceptibility of a steel to temper brittleness consisted of comparison of the energy absorption, in a notched-bar impact test at room temperature, of specimens that had received two heat treatments.' One specimen (or group of specimens) was quenched, tempered, and quenched from the temper. The other specimen (or group) differed only in having received an embrittlement treatment, which might consist of a slow cool from the tempering temperature, or of an isothermal hold in the range 45w550°C subsequent to quenching from the temper. If the energy absorption of the specimen receiving the embrittlement treatment was lower than that of the specimen not given this treatment, it was said that the steel was susceptible to temper brittleness. If the energy absorptions were the same, it was said that the steel was not susceptible. Several investigators2J found that plain carbon steels showed negligible difference in energy absorption in such tests (at least when their manganese content was below 0.60 pct), and concluded that plain carbon steels are not susceptible to temper brittleness. Recently the deficiencies of the criterion mentioned above have been pointed out.'-= The criterion now generally used for susceptibility to temper brittleness retains the heat treatments used in the earlier scheme, but instead of the notched-bar impact tests being made only at room temperature, they are made over a range of temperatures, covering the transition from ductile to brittle fracture. If the transition temperature of the specimens receiving the embrittlement treatment is higher than those

Jan 1, 1949

By J. B. Bassett, E. S. Rowland

The experimental work reported herein firas inspired by the publication of a paper by Grange and stewart, in which it was suggested that at low chromium contents the effect of this element on the Ma point was greater than at higher chromium contents, It was argued that, if a small increase of chromium in solution in austenite caused a sharp lowering of the M, point, it might lead to an explanation of occasional unsolved stress cracking of SAE 52100 during heat treatment. Commercial hardening is carried out at temperatures of incomplete solution, leaving carbides available to enrich the austenite during any inadvertent increase in temperature, such enrichment of the austenite in chromium content would lower the M, point nearer room temperature where cracking is more likely to occur. According, it was decided to investigate the effect of chromium, in quantities between o and 2 pet, on the M, point of commercial steels. Fig I summarizes the data available on chromium steels. For purposes of comparison, the M, temperature values were corrected by means of the Grange and Stewart formulas to (I) Hypoeutectoid compositions: 0.50 pct carbon and 0.80 pct manganese (z) Hsereutectoid compos-sition': 1." pet carbon and O.35 pct manganese. Only the steels having compositions reasonably close to these corrected values were included, because large adjustments inevitably introduce increased error such as are experienced in using the formula. The straight lines of Fig I were drawn by the authors. These data obviously do not settle the question Of the effect Of chromium On the Md Oint, particularly at low chromium contents. While it is fairly well established that chromium causes the M, to be lowered when present in quantities of I pet Or more, Grange and Stewart1 have suggested that in small amounts (less than 1 pet) it is more effective in lowering the Ma point, while Klier and Troiano2 have suggested that in small amounts it is less effective in lowering the Ma point. Payson and Savage3 concur with Klier and Troiano and further suggest that the effect is not linear at these concentrations. Rose and Fischer4 report that small amounts of chromium actually raise the Ma point of a 0.20 carbon steel. Material Two series of 35 Ib induction heats were melted for this work, the analyses of which are given in Table I. One series was aimed at 0.50 pct carbon and 0.80 pct manganese with chromium additions in 0,25 pet increments from o to 1.5 pct with a single heat at 2.0 pct Cr. This series was considered necessary because the commercial heat treatment of SAE 52100 does not generally introduce more than about 0.60 pct carbon in solution in the austenite. The high

Jan 1, 1949

By W. P. Sykes

In connection with a study of tungsten steels, Honda and Murakamil reported an investigation of the system iron-tungsten. This report included a tentative equilibrium diagram, photomicrographs of various alloys in the series, and qualitative measurements of hardness and magnetic properties. The present investigation originally had as its object the study of the hardness and tensile properties of a series of carbon-free iron-tungsten alloys. But, as the work progressed it became evident that the system possessed other features of interest, so the scope of the investigation was enlarged. Materials The iron powder used in preparing the alloys was obtained as follows : Iron oxalate was precipitated from a ferrous-sulfate solution by the addition of a solution of oxalic acid. This oxalate, after thorough washing, was ignited to iron oxide, which after reduction in hydrogen at about 1000" C. yielded iron powder containing but about 0.2 per cent. iron oxide and about 0.005 per cent. carbon. Tungsten metal in the form of powder, such as is used in the manufacture of filaments for incandescent lamps, was the other constituent. This material analyzed 99.8 per cent. tungsten and contained no carbon. The metal powders, in the desired proportions, were mixed by tumbling and formed, under pressure of 20 tons per sq. in., into rods 3/8 by 3/8 by 10 in. To form the alloys without contamination, sections of the pressed rods were placed on alundum slabs and heated in an alundum tube wound with a tungsten resistor and packed in heat-insulating material. While the resistor was heated, a stream of hydrogen was continually passed through both the tube and the case containing it. The tungsten winding on the tube, as well as the material heated within the tube, was thus protected from oxidation. Such a furnace will operate constantly for several days at temperatures up to 1500' C. and for shorter periods

Jan 1, 1926

By Robert S. Williams

NO new ferrous alloys have been produced in the last five or six years that are as outstanding contributions to civilization as were the high-speed steels of the early part of the century or the stainless steels announced at the close of the World War. There have been, however, many developments of major importance having to do either with modifications of existing alloys or with new methods of treatment of older iron or steel alloys. Only a limited number of these new developments can be discussed and their selection does not necessarily imply that these are the most important but merely that they are of general interest rather than perhaps of greater importance in highly specialized fields.

Jan 1, 1939

CHARLES KIRCHHOFF, Chairman. ALBERT SAUVEUR, Vice-Chairman. A. A. STEVENSON, Vice- Chairman. HERBERT M. BOYLSTON, Secretary, Abbot Bldg., Cambridge, Mass. John Birkinbine, J. Esrey Johnson, Jr., Felix A: Vogel, William H. Blauvelt, William Kelly, Leonard Waldo, James Gayley, Richard Moldenke, William R. Walker, Henry D. Hibbard, Joseph W. Richards, William R. Webster. Henry M. Howe, E. Gybbon Spilsbury, Frederick W. Wood. Robert W. Hunt, A. A. Stevenson is appointed an additional Vice-Chairman of the Committee. He will serve as Acting Chairman during the absence abroad of Professor Sauveur. For the meeting of the Institute which is to be held in New York City, Oct. 16 and 17, 1913, the following papers are now assured

Jan 7, 1913

By Wm. A. Haven

As soon as the likelihood of American participation in the war was established, and in spite of the fact that we can produce almost as much as all other countries combined, the demand for prompt deliveries of iron and steel products from American factories began to exceed their utmost capacity, and, after fifteen months of practically full operation, has continued to do so. Therefore, it would seem worth while to review some of the aspects of the supply and demand for steel as well as the efforts that are being made to bring them into balance. It is still being debated as to whether the situation should be defined as a real shortage that will persist through the emergency period or as a temporary one than can be relieved by less drastic measures than a Government-sponsored program of expanding existing plants and building new ones.

Jan 1, 1942

Correlation between Metallography and Mechanical Testing By H F MOORE (Henry Marion Howe Memorial Lecture Trans, vol 120 11,000 words ) The lack of effective correlation between metallography and mechanical testing from the days of Reaumur to the present is noted, and two general problems of such cor¬relation are stated (1) the correlation of qualitative test results with quantitative, and (2) the correlation of average results with maximum or minimum results Some correlation has been achieved, notably in the field of heat-treatment and in the study of elastic strength of metals Correlation is in the process of being developed in several other fields, including the quantitative study of the effect of grain size of metal on mechanical properties, of grain size and type of crystalline structure on the sensitivity of metals to weakening at holes, notches and other "stress raisers," and in the study of the starting phases of spreading, "fatigue" cracks Correlation of mechanical test results and metallographic study is needed in the study of such problems as the various theories of failure of metals, the meaning of and tests for service ductility of metals, and the relative damaging effect of defects inherent in metals, such as minute cracks and segregation, as compared with the damaging effect of "imposed" stress raisers, such as screw threads, sharp shoulders, etc The neces¬sity for more correlation of metallography and mechanical testing during student days is emphasized, and also the need of a better understanding on the part of engi¬neers of the significance and limitations of metallographic apparatus and results and of X-ray diffraction This should be matched by a better understanding on the part of metallurgists of the significance and limitations of mechanical tests of metals, of testing machines and strain-measuring apparatus

Jan 1, 1937

By Ming-Kao Yen

A considerable amount of work has been reported in the literature in regard to the texture and earing behavior of copper strip. The rolling texture of copper has been confirmed as (110) [112] and (112) [111], which yields ears of a drawn cup at the position 45" from the rolling direction.1-3 The recrystallization texture has been established as the cubic or (100) [001] texture, where the earing positions are at 0" and 90" to the rolling direction.4-8 It has also been reported that in the development of cubically aligned grains of copper strips, the percentage of this cubic texture increased with an increase of final reduction and final annealing temperature.8,9 A comprehensive study on H.C. copper (British commercial copper of high-conductivity quality = Cu 99.95 pct, O2 0.03 pct, Ag 0.003 pct, Fe 0.005 pct and Pb < 0.001 pct) was made by Cook and Richards.6 They concluded that the recrystallization textures could be described as one or more of the following textures: (1) a single texture (100) 10011, (2) a twin texture (110) [112] and (3) a random orientation, depending upon the previous history of the specimen concerned. The effect of various alloying additions in copper was reported by Dahl and Pawlek.10 They found that certain alloying additions, such as 5 pct Zn, 1 pct Sn, 4 pct Al, 0.5 pct Be, 0.5 pct Cd, or 0.05 pct P suppressed the formation of cubic texture. Brick, Martin and Angierll reported that the cold rolled textures due to various additions fitted a rather simple pattern. However, the recrystallization textures were subject to very considerable variations. In the discussion of this paper, Baldwin stated that deoxidized copper containing 0.02 pct P gave a complicated recrystallization texture at lower temperature. When this copper was annealed at high temperature, a single texture appeared which was described as (110) [ill] but. according to a pri- vate communication from Baldwin, this orientation reported was in error and should have been reported as (110)[112]. He also reported that the earing positions of drawn cups were at 60" to the rolling direction.12 Recently, Howald, in his discussion on the paper by Hibbard and Yen,13 reported that the rolling texture of phosphorus deoxidized copper, containing from 0.006 to 0.020 pct phosphorus, was of the pure copper type. When these coppers were annealed at lower temperatures, they exhibited a random orientation, and when they were annealed at higher temperatures they had a mixed (111)[110] and (100)[001] texture, depending on the severity of the final reduction and annealing temperature. However, the specific influence of phosphorus and other impurities on the recrystallization textures and the deep drawing properties of copper strip has not been thoroughly reported. Therefore, an attempt has been made in the present work to determine the rolling and recrystallization textures and also the earing behavior of five types of commercial copper and thereby to evaluate the effect of phosphorus and some other significant impurities on the development of texture for cold reductions of about 87 to 89 pct. Materials Used The five types of copper employed in the present investigation were two phosphorus deoxidized coppers of different phosphorus content (0.007 and 0.013 pct P), an oxygen-free copper (OFHC), an electrolytic tough-pitch copper, and a fire-refined tough-pitch copper. These materials were subjected to a thorough spectroscopic and chemical analysis. The designations and the chemical compositions were as shown in Table 1. The coppers, FA1, FA2 and FA3. were hot-forged from 3-in. billets into a ½ X 6-in. plate and cold rolled to the ready-to-finish gauge indicated below. FA4 and FA5 were hot rolled and scalped to ready-to-finish gauge. The grain size of all the materials in the ready-to-finish condition was about 0.030 to 0.045 mm. Table 2 shows the last stage of the production schedule for each copper strip used. Experimental Procedure ANNEALING, GRAIN SIZE AND HARDNESS DETERMINATIONS Specimens of each type of copper were finally annealed in air for periods of one hour at temperatures ranging from 300 to 1600°F and were subsequently cooled in air. The average grain diameter of the annealed specimen was estimated by comparing with a standard grain size chart. Hardness was determined on the Rockwell 15 T scale. CUPPING TESTS Cups were made in a blanking and drawing set, in which blanks of 2-in. diam were drawn to a cup of 1.25-in. diam with an average depth of about 0.75 in. The clearance between the punch and die was about 0.032 in. The ears of the cup were measured with a special fixture which read the height of ears to one-thousandth of an inch on every ten-degree interval along the circumference of the cup. POLE FIGURES The usual transmission diffraction method with unfiltered copper radiation was employed to determine the pole-figures of the specimens cold-rolled or annealed at 900°F. All the pole-figures were derived from the positions of intensity maxima on 111 diffraction rings of the X ray photo-grams taken at 10 rotation of a

Jan 1, 1950

By B. M. Larsen, T. E. Brower

A perspective is presented of the steel plant sulphur distribution and elimination problem from coal to liquid steel ready for teeming, giving distributions of sulphur over a range of coke sulphur content, and some methods of sulphur control, in the blast furnace, external desulphuriza-tion between blast furnace and open hearth, distribution between fuel, slag, and metal, and methods and limitations of control of sulphur in the open hearth furnace. AS a part of the 1951 AIME symposium on sulphur in steelmaking, it was thought that a discussion of the distribution of sulphur throughout the whole series of operations, from coal and ore to finished steel ingots, might have some value in giving a perspective on the whole problem. The following discussion is an attempt to present such an overall picture. The order is that of the actual plant operations, beginning with a very brief consideration of the coking process. Sulphur in Coal and Coke Since by far the largest source of sulphur entering the steelmaking cycle is in the coal used to make coke for the blast furnace, it would seem reasonable to eliminate some of it, either from the coal, or the coke, or during the coking process. This has appeared impracticable up to the present, at least, for two main reasons: the low activity of the organic sulphur in either coal or coke, and because of price limitations involved in treating a low cost material such as coke. A variable portion, usually M or less, of the sulphur is present in coal in the form of pyrites or similar compounds, and a large part of this sulphur may be removed in the coal washery. Most of the sulphur, however, is normally present as "organic" sulphur, intimately associated with the coal structure. Its distribution prevents any separation by mechanical means. Its low activity makes improbable &apos; any rapid chemical removal, although hydrogen will remove sulphur from both coal and coke. Thus, prolonged recirculation of coke oven gas in the coking process would tend to leave a smaller percentage of the total sulphur in the coke residue. Table I shows a typical distribution of sulphur from coal into products in the coking process. As the sulphur in the coal increases, the sulphur in the coke tends to increase in about the same proportion. Sulphur in the Blast Furnace The best picture of the situation in the blast furnace is provided by a sulphur balance of raw materials entering, and of products leaving, the furnace. The difficulties in accurate weighing and sampling of the variable solid materials entering this process, and the number of hours required for the raw materials to descend through the furnace under variable operating conditions, make it difficult to obtain an accurate balance. However, balances made over periods of weeks or months tend to average out some of these uncertainties. Table I1 presents three typical sulphur balances similar to a number that the writers have calculated. In most of these the slag volume calculated from the sulphur balance is, in some instances more, and in other instances less, than the value corresponding to the best input and output balances of the other slag constituents (lime, silica, alumina, etc.). Probably the greatest source of error in these calculations is the sulphur content of the slag. Despite some possible inaccuracies the balances of Table II show rather definitely the following points: 1—That 87 to 95 pct of the total sulphur input is in the coke and 95 to 97 pct of the total sulphur output is in the slag. Also, that if any sulphur leaves the furnace with the gas it is relatively small, amounting to a possible 1 pct or less. 2—At the lower sulphur coke level of 0.86 pct the total amount of sulphur charged is 15 Ib of sulphur per ton iron increasing to 26 lb per ton at the higher sulphur, intimately associated with the coal struc-rare burdens containing sulphur-rich ores will the total sulphur burden fail to be nearly proportional to the content in the coke used. 3—The 7 to 9 pct of the total sulphur input from the limestone of furnaces B and C is due to the relatively high sulphur content of the stone, 0.226 and 0.265 pct, respectively. In the case of furnace A, the sulphur content of the limestone was only 0.06 pct which resulted in only 3 pct of the total sulphur input coming from this source. It is rather interesting to compare the sulphur balances of a typical ferromanganese furnace with

Jan 1, 1952

By T. C. M. Pillay, John Chipman

The Si-0 equilibrium in liquid iron was investigated in some detail by Gokcen and Chipman1 who reported equilibrium constants of the following reactions: In each case the thermodynamic equilibrium constant K was obtained by extrapolating K&apos; to infinite dilution. Their results on reaction [I] were consistent with those of earlier observers, and there seems to be little doubt as to the approximate validity of the equilibrium product determined. On the other hand, the results on reactions [2] and [3] indicated unusual curvilinear relationships between the activity coefficients of silicon and of oxygen and the silicon concentration. Since a linear relation is observed in all other similar instances, it was concluded that some source of error had affected the gas-metal equilibrium relationships. Recently the authors undertook a restudy of these equilibria with the view especially to establishing the equilibrium constant of reaction [2]. Their experiments, however, were discontinued before a tbroughly satisfactory answer had been obtained and consequently have remained unpublished until the present time. Quite recently Matoba, Gunji, and Kuwana2 published the results of a very thorough study of these equilibria. It is the purpose of the present note to Compare these three sets of experimental data, a comparison which leads to Confirmation of the results of Matoba and coworkers. The experimental arrangements described by Dutilloy and Chipman were employed. A controlled mixture of Hz, H,0, and argon was led into the induction furnace over the surface of the Fe-Si melt contained in a silica crucible. On account of the slow approach to equilibrium, runs were continued for at least 8 hr and, in some cases, as much as 12 hr. Samples were taken by means of silica tubes at the start and during the final 4-hr period of each run. Out of twelve runs only seven appeared to have approached equilibrium with sufficient certainty to be reported here. The results of these experiments, all at 1600°c, are shown by the triangles in Fig. 1. Equilibrium was approached from both oxidizing and reducing conditions and the direction of approach is indicated by the shape of the triangle. Attempts to extend the study to higher silicon concentrations were unsuccessful on account of the transfer of silica from the metal bath to the upper portions of the crucible by means of a silicon-monoxide mechanism. Matoba, Gunji, and Kuwana used similar methods and approached equilibrium from both sides. Their results at three temperatures are shown by the solid lines of Fig. 1. A line corresponding to 1600 deg is interpolated from their equations and is shown by the dotted line. The results of Gokcen and Chipman at 1600 deg in the range 0.02 to 2.0 pct Si are also shown. At concentrations up to about 0.6 pct Si, these investigators approached equilibrium from both directions, and their results in this range appear to be trustworthy. Above 1 pct Si, however, all of their experiments involved oxidation only, as shown by a decrease in the silicon content of the bath. It now seems certain that they failed to reach equilibrium at these higher silicon concentrations and that the curvilinear relation shown by their results is incorrect. It is evident that their results in the lower silicon range are in good agreement with those of Matoba and coworkers. In fact, a straight line drawn through their results would parallel the dotted line and would differ from it by only 0.05 in log K. The new data also lie close to that line. It may thus be concluded that the three sets of data are in good agreement and that the equations proposed by Matoba and coworkers for reaction [2] are satisfactory . They take the limiting value of log K&apos;, at 0 pct Si as log K ,. The slope of each line gives the value

Jan 1, 1962

By Arnulf Muan, Egil Aukrust

Activities of COO in the three solid solution series COO-MgO, COO-MnO, and COO-"FeO" have been determined at 1200°C by equilibrating oxide samples with a metal phase (cobalt or a Co-Fe alloy) in atmospheres of known oxygen activities. The systems COO-MgO and COO-MnO are ideal, within limits of experimental error, whereas the system COO-"FeO" shows a slight positive deviation from ideality at the temperature of this investigation. KNOWLEDGE of activities of oxide components in solid solutions is necessary for the calculation of the cation distribution among coexisting oxide phases in many important metallurgical, ceramic, and geochemical systems. Very few activity-composition curves are available for such systems at the present time. A start on the prodigious job of obtaining such data has been made by studies of solid solutions with simple sodium chloride type structure containing NiO or "FeO" as a component.1-3 The present paper describes similar studies for solid solutions with sodium chloride type structure containing COO as a component. For the systems COO-MgO and COO-MnO, the activity solid solution is determined by the equation Here po and are the oxygen activities of the gas phase in equilibrium with the phase assemblages metallic cobalt plus oxide solid solution, and metallic cobalt plus pure COO, respectively. In the system CoO-"Fe07&apos;, the metal phase is a Co-Fe alloy rather than pure cobalt. In this case the activity of COO in the oxide solid solution is calculated from the equation where a& is the activity of cobalt in the metal phase. Because the system Co-Fe is close to ideal4 and the metal phase appearing in the present work contains less iron than cobalt even at the lowest COO concentrations of the oxide phases used may be set equal to N . Hence, Eq. [2] becomes The value of Nc, can be estimated from the relative stabilities of the oxides COO (see below) and I) EXPERIMENTAL METHOD Two different but closely related procedures were used in the present investigation to determine the activity-composition curves. In one method, the oxide solid solutions were equilibrated against

Jan 1, 1963

By J. Chipman, H. R. Larson

The data previously reported for the quantity as a function of oxygen pressure at 1550°C have been used to compute the activities of Fe, FeO, Fe2O3, and COO in slags of the ternary system. Activities of the first three have been obtained also for two quasi-ternaries involving fixed CaO:SiO2 ratios. IN a previous paper1&apos; the authors reported the results of an investigation into the effects of oxygen pressure on the composition of various simple slags analogous to some of those which are found in steel-making practice. The ratio of ferric iron to total iron was studied at 1550°C in iron oxide melts to which lime, magnesia, lime-plus-silica, and other oxides were added. The oxygen pressures involved those of air, carbon dioxide, and carbon dioxide plus carbon monoxide in several proportions. Although very low oxygen pressures could not be used, the slag-metal equilibrium studies of Fetters and Chip-man&apos; permitted extending the results to slags in equilibrium with iron. For the ternary system CaO-FeO-Fe,O, the oxygen pressure-composition relationship has been determined from zero percent lime to lime saturation over an oxygen pressure range from air to that represented by equilibrium with liquid iron. Lime and silica were added to iron oxide in the ratios 0.54, 1.306, and 2.235 to form three quasi-ternary systems which were also studied over the entire region of liquid melts at 1550°C. Ternary Gibbs-Duhem Equation Wagner" has developed a method by which the activities of two components of a ternary system can be calculated if&apos; the activity of the third component is known throughout the composition range being considered. The fundamental form of the Gibbs-Duhem equation for ternary systems is N, d In a, + N2 d In a, + N3 d In a3 — 0. Wagner has developed a usable form of this equation by introducing the term y = N3/(N3 + N1) and rearranging the equation to give (? In a1/?N2)x = -y/(1-N2)2 (? In a2/?y) x2 N2/1-N2 (? In a2/?N2) y To apply this equation to the slag system CaO-FeO-Fe,O,, the activity of one of these components must be known. However, the only activity which is known from the experimental data is the activity or pressure of oxygen in the gas phase with which the slag is in equilibrium. The activity of oxygen in the slag can be defined as the square root of the oxygen pressure. In order to use oxygen as one component, the composition of the slags must be converted to the basis Fe-O-CaO. Oxygen, exclu-sive of that contained in CaO, then becomes com-ponent 2 in Eq. 1, Fe is selected as 1 and CaO as 3. Then y = Ncao/(Ncao + NFe). Eq. 1 then becomes (? In a fe/? Nx) y = -y/(1-No)2 (? In ao/?y) xp - No/ 1- No (? In ao/ O No) y. In order to evaluate Eq. 2, the boundary conditions must be known. The obvious choice of a standard state for iron is to assign slags in equilibrium with iron an activity of one. Then In a,., or log aFe, which is substituted for convenience, is determined by integrating along a line of constant y from the slag in equilibrium with iron to the composition at which log a,, is to be determined. Mathematically this can be expressed as log aFe (Nlog y) = log a&apos;Fe + (? log a Fe/? No) dNo where the primes indicate equilibrium with liquid iron and a&apos;Fe = 1. When Eq. 3 is integrated along a line of constant y, the following is obtained: log a Fe (No, y) = - y No ?/ No ? y ( log ao/(1- No)2) No d No - No/1-No d log ao. Lime-Iron Oxide Slags Iron Activity: The oxygen pressure of lime-iron oxide slags is shown in Fig. 1 as a function of j, de-fined as j = Fe+++/(Fe++ + Fe +++) for various constant mol percentages of lime. The j values at 0 to 60 pct lime were then determined for oxygen pressures of 1, 10-1, etc. to 10-" atm. For purposes of calculation, the line for zero percent lime was extrapolated to

Jan 1, 1955

By E. T. Turkdogan

Jan 1, 1962

By John Chipman, Peter J. Koros

The distribution of copper between liquid silver and liquid iron, Fe-C, and Fe-C-Si alloys was studied at 1600°C. From the data and the activity of copper in silver obtained from the phase diagrams, the activity coefficients of copper in iron up to 3 atomic pct are log y, = 0.90 — 2.4 N, and y": = 8.0. The effect of carbon up to saturation is expressed as log y = 1.8 Ni. The effect of silicon up to 11 atomic pct is insignificant. SLOWLY rising copper content of steel being produced is focusing attention on the lack of information on the behavior of copper in liquid steel. Chou&apos; measured the activity coefficient of copper in liquid iron through its distribution between liquid silver and iron. He obtained a value of 9.12 at infinite dilution. Chipman&apos; calculated the activity coefficient of copper in liquid iron at 1600°C from the phase diagram and obtained a value of 12. Ameiln and Pfeiffer- investigated the possibility of removing copper from liquid steel through the use of lead, silver, and bismuth as solvents. They also calculated that the activity coefficient of copper in liquid iron would be increased five or sixfold by addition of 4 pct C. Langenberg and Lindsay&apos; are investigating the effectiveness of lead and sodium sulfide in removing copper from liquid iron. The experimental method was similar to that used by Chou&apos; and Chipman and Floridis." Silver, copper, and iron were melted together, the copper distributing itself between the two layers so that at equilibrium a,.,, (in Ag) - a,.,, (inFe) and N,.,, (in Ag) yc,, (inFe) (in Ag). [21 N,.,, (inFe) Thus, from the concentration of copper in each layer and the activity coefficient of copper in liquid silver, the activity coefficient of copper in liquid iron can be calculated. The effect of an alloying element such as carbon on the activity coefficient of copper can also be determined from its effect on the distribution, provided the new element is essentially insoluble in silver. The one serious limitation of this method is the increasing mutual solubility of silver and iron in the presence of a third component. The mutual solubilities of pure iron and silver are negligible, but increase with increasing copper concentration beyond the range investigated. Ag-Cu System—The activity coefficient of copper in silver at 1600°C had to be estimated, since direct experimental data are not available. Kawakami" measured the heat of mixing of Ag-Cu alloys at 1200°C by a calorimetric technique. Scheil&apos; calculated the heat of mixing in this system by use of the phase diagram, assuming that the solution is random and that the heat of mixing is independent of temperature. The terminal solid solution may be assumed to obey Raoult&apos;s law within the requirements

Jan 1, 1957

By C. E. Lesher

Two processes for agglomerating fine sized ores with low temperature coke are described. One process (Orcarb) agglomerates ores with limited amounts of carbon; the other (ore-carbon pellets) pelletizes fine sized ores, using low temperature cokes as the binder. Data are presented on the products obtained when taconite, magnetite, and hematite concentrates and several titanium oxide ores were used. REDUCTION of finely divided oxide ores has long been a major metallurgical problem. Various methods of agglomerating, notably briquetting, sintering, nodulizing, and pelletizing, have been developed and are in industrial use. Carbon is the usual reducing agent for oxide ores and attempts have been made to agglomerate fine ores with coking coal in order to get the metallurgical advantages of intimate contact and increased particle sizes suitable for subsequent smelting or reduction.1-3 For the past three years, research has been in progress on two processes for combining fine sized ores with reactive carbon in the form of low temperature coke. By one process, designated Orcarb, the ore is coated with coke carbon in amounts limited to that required for reduction and the result is an appreciable but limited increase in product particle size over that of the original ore. The second process, called ore-carbon pellets, produces pellets of fine ore bound by low temperature coke. The first objective was to develop a process for combining an oxide ore with only enough carbon for its reduction, the amount being that derived from the simple reduction equations. The carbon (that amount theoretically required for the reduction of the oxides in an agglomerate with ore), calculated to CO, ranges from 10 pct for zinc calcines Or phosphate Ores to 21 pct for rutile and 26 pct for silica. The carbon required for reducing iron oxides falls between the extremes—10 and 26 pet&apos; When the carbon in the agglomerate is low temperature coke, reduction by hydrogen may be expected to take place, thus modifying the calculated carbon require- ment. On the other hand, in practice, an excess of reducing agent over that required in theory is normally used. Screen sizes of three iron ores on which the greater part of the experimental work was performed are given in Table I. Most of the testing was done with Freeport (Renton Mine) coal, but Pittsburgh and Elkhorn bed coals were also used. The first objective was to agglomerate the ore with sufficient low temperature coke to provide carbon for reduction, i.e., in the range from 10 to 25 pct. This was accomplished by preheating the ore, mixing with pulverized coking coal, and completing the coal carbonization in a plain steel rotating retort. The pilot retort, as it is now constituted, is shown by diagram in Fig. 1 and photograph, Fig. 2. Ore is fed into the upper kiln by a screw; it is heated to 900° to 1100°F by direct flame and it is then discharged into an insulating hopper from which it goes

Jan 1, 1956

By E. C. Nelson

An equation is derived for calculating approximately the solubility of nitrogen in an alloy steel over a temperature range from 1200" to 1900°C using data on the effects of alloys on the activity coefficient of nitrogen in iron at 1600°C. Calculated values are compared with experimental data. STUDIES involving the solubility of nitrogen in complex iron alloys have been aided by a system devised by Lanenberg for calculating the solubility of nitrogen in ternary and higher alloys of iron at 1600C using available data on the activity coefficient of nitrogen in binary alloys. Calculated nitrogen solubilities agree well with experimental determinations on complex alloys,&apos; and the calculation procedure is useful when experimental data on specific alloys are not required. A method for extending calculations of nitrogen solubilities in alloy steels to temperatures other than 1600°C would be useful because of the paucity of experimental data in this area. Assumptions which must be made to permit this calculation introduce an element of approximation; however, the calculation is useful in cases where a large number of alloys and temperatures are under consideration, and the precision of experimental measurements is not required. The activity coefficient can be freed from the isothermal restriction of Langenberg&apos;s calculation if it is assumed that the solution is "regular". This is a reasonably good assumption for solutions of nitrogen in iron alloys, because usually they are dilute and the heat of solution is relatively small. The presence of large amounts of certain alloying elements would make this assumption less tenable.3 One of the properties of a regular solution is: when "f" is the activity coefficient of a solute and "T" is the temperature in OK. If the activity coefficients at temperatures T and T2 are denoted by subscripts 1 and 2 respectively: The activity coefficient of nitrogen in a complex alloy then can be calculated at any temperature from the sum of the effects of the other solutes on the activity coefficient of nitrogen at 1600°C. It will be convenient to incorporate this expression into an equation which will give directly the logarithm of the solubility of nitrogen at any temperature in the presence of alloys. Two constants, the solubility of nitrogen in pure iron at 1600 "C and the partial molar heat of solution of nitrogen in iron are required. The heat of solution value of 1400 cal at 1600 °C quoted by Kubaschewski and vans will be used. The accuracy of the calculation will suffer if the heat of solution should change materially with temperature. In the absence of experimental data on the temperature dependence of this quantity, it will be expedient to restrict the application of the calcula-

Jan 1, 1963

By G. W. Healy, W. D. Forgeng, N. R. Griffing

The liquidus surface of the C-Cr-Fe system to 1900°C has been mapped from carbon solubility and freezing point measurements, metallographic observations, and published data. In the graphite field, the surface drops precipitously to meet a slope descending from the Cr3C2-C join through the mixed carbide fields to a meandering eutectic valley; this flows from the Cr-C to the Fe-C edge. On the low-cavbon side of the valley, the surface rises to the chromium corner. 1 HE purpose of this inquiry was to gain a more complete knowledge of the C-Cr-Fe liquidus surface. This required an extensive study of carbon and chromium-rich alloys which, in the past, have received the least attention. Sanbongi el al.1 measured the solubility of carbon in liquid alloys containing up to 25 wt pct Cr at temperatures of 1460" and 1545°C and found that it increased with chromium content and temperature. Lucas and wentrup2 also made solubility experiments in graphite crucibles, but at higher chromium contents in the range of 38 to 80 pct and at temperatures of 1550" to 1700°C; however, their experiments were relatively few, and the effect of temperature was not consistent. The present investigation included compsitions ranging from 0 to 94 pct Cr and 1.4 to 11.7 pct C at temperatures up to 1800°C. Saturation experiments in graphite crucibles, thermal analyses during freezing, and metallographic and X-ray diffraction studies were employed to obtain desired information which, together with published data from other investigations about the binary edges and the ternary system at lower carbon levels, made possible the construction of a comprehensive C-Cr-Fe liquidus surface. C-CR-FE LIQUIDUS SURFACE The derived ternary liquidus surface is shown in Fig. 1 where the solid curves are liquidus isotherms. The dashed curves and the binary edges are the liquidus field boundaries of the primary crystals, the arrows along these boundaries indicating their direction with decreasing temperature. Each intersection of three internal boundaries is the composition of liquid involved in four-phase isothermal equilibrium with the three primary crystals whose liquidus fields touch there. It may be observed that a total of six four-phase equilibria occur at melt temperatures. The general topography of the liquidus surface consists of a eutectic valley sloping down from the Cr-C edge at 3.2 pct C to the Fe-C edge at 4.25 pct C. The liquidus surface rises in two directions from the valley, one hillside terminating at the Cr-Fe edge while the opposing hillside rises toward the carbon corner out beyond the top edge of Fig. 1. Both hillsides are interrupted by intersections of surfaces or joints representing the liquid colmpositions of various three-phase equilibria. Seven primary crystals occur, namely a(bcc), ?(fcc), graphite, (Fe, Cr)3C capable of dissolving as much as 15 pct Cr, and three different chromium carbides in which iron is partially soluble by substitution for chromium. The composition ranges of the primary l~hases at melt temperatures are shown

Jan 1, 1962

By G. W. Healy

Activities of oxides in the ternary FeO-MnO-SiO system are calculated from data on the binaries, using the Gibbs -Schuhmann method. These activity data are used, together with thermodynamic relations, to calculate the contents of oxygen, silicon, and manganese in liquid iron, and the composition of the corresponding oxide phases in equilibrium with it. Excellent agreement was found between calculated and experimental metal compositions, indicating that thermodynamic methods, such as those presented, can supply useful data on slag-metal equilibria, and are therefore capable of wider application. When molten steel containing oxygen is treated with a deoxidizer, the product of the reaction is composed of the oxides of the deoxidizing elements together with more or less iron oxide. This assemblage of oxides is presumably very close to equilibrium with the melt since it is formed directly from elements dissolved in it. Therefore, it is likely that metal analyses corresponding to a given oxide phase composition are calculable with considerable precision where sufficient information is available on activities in both phases. The present study was made to test this hypothesis in the case of deoxidation by silicon and manganese. Here experimental data on the relationship among oxygen, silicon, and manganese in the metal phase are available, as are activity data for the MnO-SiO2, FeO-SiO2, and FeO-MnO binary systems. A further objective was to demonstrate in a specific example how gaps in the data may be filled by judicious interpolation, and how the Gibbs-Schuhmann method is applied to obtain activities in ternary systems. The rigorous derivation of the latter, requiring considerable facility in handling partial differential equations, tends to mask the real simplicity of the method, and repel the very engineers and experimenters who would benefit the most from its use. I) GENERAL PROCEDURE Deoxidation of steel with silicon and manganese results in the formation of two phases, metal and oxide. The number of components is four—Fe, Mn, Si, and 0. The phase rule then gives the number of degrees of freedom: If temperature and pressure are fixed, 2 deg of freedom remain. Consequently, it suffices to fix the concentration of two of the components of the system for all the concentrations to be determined and, at least in principle, calculable. The practical problem is then to discover a straightforward method of calculation with a minimum of trial and error solutions. The type of reactions to be dealt with are like the following: Si (in steel) + 2 O(in steel) = SiO2(in oxide phase) The standard free energy of this reaction is readily obtained from the literature, so that the equilibrium constant can be calculated by the expression: Since the metallurgist is concerned with percentages rather than activities, a real solution requires information that will permit activities to be converted to concentrations. This information is fairly adequate for the elements dissolved in molten steel, but is not available for the FeO-MnO-SiO2 oxide system. Therefore, the first step in this study was to find the relation between activity and concentration for the three oxides making up the system. The next step was the calculation of slag and metal compositions that are in equilibrium with one another. As shown by the phase rule, it is sufficient to assume the values of two composition variables in order to calculate the rest. But which two are picked makes a vast difference in how difficult the calculation turns out to be, so that it is necessary to make preliminary trial of several routes to find the easiest one. A convenient method is to begin with the slag, and to pick FeO concentration and that of one of the other oxides. Starting with the relation between oxygen dissolved in the liquid steel and FeO in the slag permits a good first estimate to be made of the oxygen content of the steel in equilibrium with FeO, because the iron content being fairly close to 100 pct, Fe has an activity close to unity. Using this first estimate of oxygen activity, approximate values of the manganese and silicon contents of the melt are readily calculated from the appropriate equilibrium constants. A minor correction for the depar-

Jan 1, 1963

By J. F. Peters

The term solution loss is discussed and defined. Examples are given showing that solution loss may either have a favorable or unfavorable effect on blast furnace performance. A theory is advanced explaining the contradictions encountered during earlier studies of the problem. MANY papers have been written and numerous theories advanced concerning the chemical reactions in the iron blast furnace. Richards, discussing the utilization of fuel in the blast furnace, said: "All the carbon burnt in the furnace should be first oxidized at the tuyeres to CO and all the reduction of oxides above the tuyeres should be caused by CO, which thus becomes CO,. This dictum is not Gruner&apos;s own words, but expresses their sense, and from the point of view of the present discussion, it is the correct principle upon which to obtain the maximum generation of heat in the furnace from a given weight of fuel." Richards also said the cubic feet of wind per pound of coke does not express how efficiently a furnace is running, but it will be shown later where Howland paid much attention to this figure. Richards pointed out a shortcoming of Grun-er&apos;s discussion in that Richards claimed there is not enough CO generated in a good working blast furnace to possibly combine with the oxygen of the ore, so some of it must be reduced by solid carbon. Mathesius2 made two important contributions. First, he wrote the direct and indirect reduction equations showing the heats of formation in each case and stressed that the indirect reduction reaction produces heat while direct reduction absorbs heat. Secondly, he substantiated Richards&apos; conclusion that there may be a deficiency of CO gas and that some direct reduction is required, but he pointed out that "direct reduction in the hearth has often been confused with direct reduction in the stack, which more correctly should be termed &apos;premature combustion&apos;." Then he said that "as this carbon is consumed partly by economical (direct reduction in the hearth) and partly by wasteful reactions in ever varying proportions, it is evident that the relation of this total carbon gasified above the tuyeres to the fuel consumption of a furnace cannot possibly have a direct bearing on the economy of furnace operation. The premature combustion of carbon (CO2+C?2CO) must therefore in all cases be considered a detrimental reaction." Johnson3 discussed this whole problem. Through experience Johnson found that when departure from Gruner&apos;s ideal working is considerable, the fuel economy is poor, and when the quantity of blast goes up, the fuel economy goes up in the same proportion. Johnson said that the wind per pound of coke is a measure of fuel economy, which will be shown to be wrong. Howland4 established the fact from operating data that there was no relationship between the coke consumption of a furnace and the percentage of coke burned at the tuyeres. But some of his calculations have led to questionable conclusions, such as "it is practically impossible to obtain low coke consumption unless we keep our wind low." Howland made an important contribution when, in his concluding paragraphs, he reached certain negative decisions as to why one coke works better than another. He said the reason one coke works better than another in a blast furnace is not because: 1—there is any difference in the percentage gasified at the tuyeres, or 2—there is any difference of wind required per pound of coke. Korevaar5 propounded a new theory on combustion which disagrees with Gruner because "as far as we can see, this is sufficient proof for the invalidity of Gruner&apos;s ideal, for if it was valid, the Low coke consumption should always be accompanied by a higher percentage of carbon burned at the tuyeres. This not being the case, we must deliberately give up our belief in Gruner&apos;s ideal." The part in italics will be proved to be wrong. Martin" contended that there was not enough reducing capacity in the blast furnace unless solution loss occurs. He came to a series of conclusions, such as: 1—The greatest efficiency of the blast furnace may not be attained when the reduction is performed entirely by carbon monoxide as demanded by Gruner&apos;s definition of the ideally perfect working of the blast furnace. 2—The so-called solution loss reactions, which should more properly be termed direct reduction reactions, promote furnace efficiency. 3—In modern blast furnace practice, the carbon consumption of the process is determined primarily by the carbon needed for reduction purposes, any thermal deficiency created by the reduction process being balanced in practice by the use of blast heat. From a practical standpoint further discussion of the theory of chemical reactions is of little moment. There is however one phase of the general theory of chemical reactions which is very important and that is the combustion of coke in the blast furnace. The purpose of this paper is to show that the reaction

Jan 1, 1955