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By J. H. Wernick

IN a study of the rate of decomposition of %-in. cubes of limestone, dolomitic limestone, and dolomite in a sweeping nitrogen atmosphere, Joseph, Beatty, and Bitsianes' found that the zone of calcination advanced at substantially a constant rate at a given temperature. The purpose of this paper is to show that their data indicate at least two processes occurring when the above carbonates decompose. It is suggested that there are two processes occurring consecutively, one the limiting process in a given temperature range and the other the limiting process over another temperature range. The data of Joseph and co-workers have been extended by use of the Arrhenius equation, Eq. 1, a relation beween the velocity constant k of a reaction and the absolute temperature T. d In k A ' dT Rr2 A is the activation energy in calories per mol (additional energy that must be added for an average molecule to react), and R is equal to 1.986 cal

Jan 1, 1955

By H. H. Podgurski, J. C. M. Li, Y. T. Chou, F. N. Davis, R. A. Oriani

When an Fe-2 pct A1 alloy is nitrogemted at 500ºC with a gus tnixture (NH3-H2) in which the nitrogen activity has been kept Lou] enough to avoid the formation of iron nitride, a two-phase alloy is generuled which consists of AlN particles and a ferrite phase cotaining a heavy network of dislocations. The amount of nitrogen contained in such an alloy, when equilibrated with the nitrogenating atmosphere, far exceeds both that needed to satisfy the normal solution requirements of a Fe and that needed to convert all of the aluninuwi to AlN. This excess nitrogen is accounted for as being trapped on dislocations, adsorbed at the ferrite -AlN interface, and as an en-Iuznced lattice solubility in strained ferrite. This excessive uptake of nitrogen had previously been attribuled by other investigators to the formation of a nonstoichiornetric aluminum nitride. Isotope exchange experiments revealed various amomts of exchargeable N14 present in the originally nitrided samples that could not be removed by reduction with HS at 500ºC. This exchangeable nitrogen has been identified us that bound to the AlN -ferrite interface. Estimates of inter facial areas in alloys containing -3 pct by weight of A1N are as high as 10 sq m per g of alloy. ThE first1 of this series of papers described the experiments forming the basis for the elucidation of the mechanism of the formation of aluminum nitride particles within an Fe-A1 alloy. It was found that not only are dislocations necessary for the nucleation of the AlN particles but also the nitriding reaction in turn produces a dense network of dislocations in the ferrite matrix. It was also observed that nitrogen in excess of that needed for the formation of stoichiometric AIN is taken up by the alloy without the formation of an iron nitride. The present paper is an analysis of the excess nitrogen sorbed by the nitrided Fe-A1 alloy which considers the structure generated by the formation of the aluminum nitride particles. Because of the high density of dislocations, this system has proved to be quite useful in studying nitrogen-dislocation interactions. Wriedt and Darken2 have already reported such studies in a cold-worked ferritic steel. EXPERIMENTAL Nitrogen Sorption. The specimens of this alloy were cold worked (50 pct reduction) to 0.011 in. thickness and chemically cleaned in a 2:1 concentrated phosphoric acid-50 pct hydrogen peroxide solution before nitrogenation. The flowing nitrogenating atmosphere (11 pct NH3-89 pct H2) was established before the ni- trogenating temperature was reached. The same gravimetric and gas-flow equipment described in an earlier paper1 was also used in this investigation. Changes in nitrogen concentration were followed gravimetrically when establishing the isotherms. In the isotope exchange experiments, concentration changes were followed volumetrically. In some instances chemical analyses (Kjeldahl method) were used to check for material balances in both the gravimetric and volumetric procedures. Our objective was to obtain reversible nitrogen sorption isotherms for alloys equilibrated with NH3-H2 gas mixtures over a large range of nitrogen activity* and temperature. The *Defined as equal to PN H /P3/2 ,where P corresponds to partial pressure in atmospheres. Actually the nitrogen activity, aN, in the alloy equals K PNH3 /p3/2H2, where K is the equilibrium constant for the reaction NH3 = N + 3/2H2. upper limit for nitrogen activity was below that which would produce iron nitride; for reasons which will become apparent later, most of the sorption studies were made in the temperature range between 400" and 500'C. The alloy sample studied most extensively in this investigation was given a series of successive reductions (100 pct H2) and nitrogenating treatments at 500°C to attain a stabilized structure. Presumably some dislocations were lost during these treatments. Throughout most of the sorption studies, temperature was held at ± 1°C and the gas-phase composition was held to k0.2 pet NH3. Based upon the results of numerous diffusion experiments with this alloy* the 'Results to be published in a third paper of this series. times chosen for equilibration were considered more than adequate, Nitrogen Isotope Exchange. The first exchange experiments were carried out by circulating a measured quantity of H2 and NH3* (with a predetermined isotopic *The NH3 was synthesized over a synthetic ammonia iron catalyst using H2, and N2, enriched with N15. AS the NH3 was formed it was removed continuously from the gas phase by circulating the mixture through a refrigerated (78ºK) trap. When most of the nitrogen had been converted to NH3, it was distilled from the trap into a glass storage vessel. composition, N15/N14) over a nitrided specimen at 450°C in a closed glass system. Attempts to reach the exchange limit isothermally, i.e., an isotopic ratio (N15/N14) in the gas phase identical with the extractable nitrogen in the nitrided alloy, were futile. The exchange rate between the gas phase and the alloy was too slow, involving many days. To circumvent the slow exchange with the gas phase, measured amounts of N15 and N14 were introduced into the alloy by nitrogenating with an NH3-H2, mixture at 500°C containing an N15/N14 ratio of 7.76; the charged specimens were then sealed by allowing an oxide film to form over them in air, and in this final condition the specimens were subjected to 500°C in vacuum for various times during which isotope exchange was allowed to proceed within the specimen. It was estab-

Jan 1, 1970

By J. B. Conway

RECENT studies of first-stage creep data have led to a special least-squares procedure for use in calculating the Andrade creep-equation constants. This procedure is easy to apply, uses only experimental data, and yields results which do not vary with the individual making the analysis. Taking logarithms of the Andrade equation I = lo(1 +ßl1/3)ekt leads to In l = In lo + In(1 + ßl1/3) + kt A residual can then be defined as the difference between the logarithm of an experimental length value and the logarithm of a calculated length value. Following the usual least-squares approach differentiation with respect to 10 and k of the expression for the sum of the squares of the residuals leads to two equations which cannot be solved directly for the three constants involved. However, values of ß can be assumed and corresponding values of lo and k can then be calculated. Then the sum of the squares of the residuals can be calculated. For a certain value of ß this sum will exhibit a minimum value to identify the least-squares values for ß, lo, and k. An illustration of the application of the above least-squares approach is obtained by employing the creep data for lead originally presented by Andrade' in the development of the one-third creep law. A plot of the sum of the squares of the residuals as a function of ß is shown in Fig. 1 along with the corresponding lo and k values. The minimum is seen to occur at ß = 0.0143 from which lo = 39.149 and k = 7.27 x 10-5. Based on these results the Andrade equation for these lead data is 1 = 39.149(1 + 0.0143 t1/3)e7.27 x 10-5t compared to the equation I = 39.13(1 + 0.01457/l/3)e6.95 X 10-5t developed by Andrade (calculation procedure not given but was probably a graphical solution). Calculating length values using the above equations and comparing with the experimental length values, residuals corresponding to the difference between experimental and calculated length values can be obtained. The sum of the squares of the residuals can then be calculated. For the case of the least-squares approach a value for this sum was found to be 2.37 x 10-3 compared to a value of 3.47 x 10" for the equation originally proposed by Andrade. In other words, while the constants proposed by Andrade provide an excellent fit for the lead data, a still better representation is afforded by applying the least-squares approach discussed above. A similar result was obtained for the lead data at 162oc.l The least-squares approach led to 1 = 9.866(1 + 0.0497 t1/3)el.011 x 10-2t compared to the equation I = 9.94(1 + 0.0433 l1/3)e1.06 X 10-2t given by Andrade. In this case, too, the sum of the squares of the residuals is slightly smaller for the equation based on the least-squares approach. It should be pointed out that a double trial and er-

Jan 1, 1965

By D. R. de Fontaine

By computing the values of the resolved shear stress for a great many orientations of the tensile axis on all 12 (111) <110> slip systems, Taylor and Elam&apos; were able to map out a stereogram of slip systems, given here, Fig. 1, in Diehl&apos;s notation (see, for example Hauser and Jackson&apos;). The purpose of the present note is to derive such a stereogram by mathematical reasoning and to give a simple method of defining the system with highest resolved shear stress for any direction of the tensile axis, thereby eliminating tedious computation of the Schmid factor. Let the tensile axis, of direction cosines [xyz], make the angles 0 and relative to the slip plane normal [HKL] and the slip direction [hkl] respectively. Fig. 2 shows a simple cubic cell whose edges define a Cartesian coordinate system. Let A be a unit vector of origin 0 supported by the tensile axis. Its end point [xyz] then lies on a sphere of unit radius, on the positive hemisphere of which the planes z = 0, x = y, y = z limit the so called primitive triangle of a standard [loo] stereographic projection. For any point of this triangle, the following hold: 0«2 «y *sx « 1 [l] Now, for given [xpz], the most highly stressed system will be the one for which the following expression of the Schmid factor has the largest possible value: (HKL) belonging to the form (111) and [hkl] to the form <110>, the denominator on the right hand side of Eq. [2] is equal to 1.fi.a. The unknowns H, ... h, . . . can now be determined by correctly distributing +1 and -1 to H, K, L, and +1, -1,O to h,k, 1 giving the numerator of Eq. [2] its greatest positive value (the angles 0 and cpare chosen acute) subject to the condition: Hh+Kk+Ll = 0 [3] since the slip direction must lie in the slip plane. Condition [3] can be satisfied only by having one of its terms vanish and the other two equal to +1 and -1. Taking the case of the tensile axis lying in the primitive triangle, x2y, z and therefore Hh will be chosen equal to +1 with H = h = +1 in order to retain a maximum of + signs. There remain two alternatives: either 1 = 0 or k = 0. 1) take 2 = 0 Then Kk = -1 and the numerator of [2] becomes: Retaining a maximum of + signs gives L = k = +1 and becomes: 2) take k = 0 Then L1 = -1and, likewise, the numerator of [2] becomes: xz - z2 + xy + yz [5] But the difference [5] - [4] can be written: (y - z)(x+y +z) [6] positive or zero by virtue of [I]. If [6] is positive (y > z), then the largest possible value for the numerator of [2] for any xyz inside the primitive triangle is given by [5], i.e., for: That is to say, the active slip system corresponding to an orientat:~on of the tensile axis inside the primitive triangle is:

Jan 1, 1962

By Heinrich Siemes, Christa Hennig-Michaeli

The effect of temperature and orientation on the plastic behaviour of an impure titaniferous magnetite from Palabora (South Africa) has been investigated. Prismatic specimens were axially compressed along [001], [110] and [111]. The strain rate was about 3.10-6 s-1. Trace analyses of glide lines revealed a plastic anisotropy due to glide on several non-equivalent crystallographic planes. At 200 °C {111}<112> twinning, {111}<110> slip and {100}<011> slip have been operative. At 400 °C {111}<110> slip and {100}<011> slip are primary glide modes. {110}<110> slip seems to be a secondary glide mode in the [001] specimens at 400 °C. {III}<112> twin glide generates isolated sharp lamellae. {111}<110> slip produces smooth surface striae. {100}<011> slip is characterized by coarse and partially wavy slip bands. {I10}<110> slip displays faint remarkably equidistant glide lines. The critical resolved shear stresses (CRSS) decrease with rise of temperature. The CRSS ratio of slip on [100) to slip on {lll} increases from 1.1 at 200 °C to 1.4 at 400 °C giving rise to a distinct flow anisotropy of the magnetite at 400 °C.

Jan 1, 1982

By G. L. F. Powell

TURNBULL and Cech1 reported undercooling small droplets (approximately 50 mp in diam) of a large number of metals by approximately 0.18 TE (TEoK = the equilibrium freezing point). This was considered to be the undercooling for homogeneous nucleation in

Jan 1, 1968

By John Chipman, Wayne L. Worrell

Using recent thermodynamic data for the carbides and oxides of tantalum, columbium, and vanadium, the stable solid phases aboue 1300°K and at 1 atm CO(g) pressure in each M-C-O system have been determined. Pourbaix-Ellingham diagrams have been constructed and are used to estimate the minimum temperatures necessary to obtain these metals by carbothermic reduction under reduced pressures. TANTALUM, columbium, and vanadium react with carbon to form the M,C and MC carbide phases. The M,C carbides have narrow homogeneity ranges, but the MC phases possess considerable ranges of homogeneity.&apos; The free energies of formation of the MC carbides were recently measured, and those of MC and M2Ccarbides have been estimated., (The notations MC and MC denote, respectively, the carbide compositions at the carbon-rich and at the metal-rich boundaries of the homogeneous MC phase.) In this paper, these carbide data are combined with thermodynamic data for the oxides to predict the stable phases in the Ta-C-O, Cb-C-O, and V-C-O systems at specified temperatures and oxygen pressures. M-C-O PHASE DIAGRAMS The solid phases in equilibrium above 1300°K for each M-C-0 system are shown in Figs. 1, 2, and 3. Above 1300°K, the vapor would consist entirely of carbon monoxide; and these phase diagrams were obtained by specifying the vapor pressure of CO(g) to be 1 atm. The temperature at which any two solid phases become unstable is specified along the two-phase equilibrium lines in each diagram. To simplify the diagrams, all solid phases in the Ta-C-O and Cb-C-O systems are represented as stoichiometric compounds. However, the homogeneity ranges of VO and VC are indicated, because they form a complete series of solid solutions.3 It also should be noted that the homogeneity range of

Jan 1, 1964

By G. Horvay, J. G. Henzel

Nomograms and charts are provided which permit rapid determination of the mold-casting interFace temperature and the speed of solidification when a semiinfinite ingot is cast into a semiinfinite mold. The physical properties of the mold, the frozen metal, and the liquid metal are not assumed to be equal, but they are assumed to be temperature independent. The results are correlated with charts prepared by Paschkis and Hlinka—using Columbia University&apos;s Analog Computer—for the freezing of finite ingot slabs when a constant surface temperature is impressed, and with experimental cooling curves obtained by Pellini in sand and iron molds for 7 in.-sq steel castings. x = coordinate directed to right of zero (freezing direction) y = coordinate directed to left of zero (melting direction) e = position of freezing front (in x direction) ? = position of melting front (in y direction) $ = temperature (function of position and time) ? = fixed temperature 2L = width of finite slab Y = a fixed position in the semiinfinite mold, or the width of a finite mold (see Section 7) Properties y = density, lb per cu ft c = specific heat, Btu per lb deg F k = conductivity, Btu per ft hr deg F X = heat of fusion, Btu per lb k = k/Yc = temperature diffusivity, sq ft per hr* b - = = - heat diffusivity, Btu per sq ft deg F Vhr Subscripts 2 = pertaining to the still-liquid metal (case of freezing), or the medium which furnishes the heat (case of melting) 1 = pertaining to solidified metal (case of freezing), or to molten metal (case of melting)* 0 = pertaining to mold (case of freezing), or to the still-solid metal (case of melting) f = pertaining to fusion condition (eg.,?f - fusion temperature) i = pertaining to the location x = 0 (the mold-ingot "interface")

Jan 1, 1960

By J. Chipman, T. P. Floridis

The equilibrium of gaseous H2O-H2, mixtures with liquid iron is used to establish the activity coefficient of dissolved oxygen as log f o = — 0.20 [%O]. Discrepancies in earlier work are explained or eliminated. Additions of cobalt, molybdenum, tungsten, and platinum increase the activity coefficient of oxygen whereas copper and gold decrease it. THE importance of dissolved oxygen as a refining agent in liquid steel has led to a number of studies of its thermodynamic properties. It is the purpose of this paper to examine the effects of a number of alloying elements on the activity of oxygen in the steel bath. As a preliminary to the study of the effects of alloying elements, it seemed desirable to examine certain minor discrepancies in the published data on the properties of oxygen in simple iron-oxygen solutions. The activity and free energy of oxygen in liquid iron were studied by chipman1 in 1933 by means of the reaction h2(g)+-o = H2O (g);K; = PH2O/PH2[% O] [1] The data were in fairly good agreement with calculations from results on the similar reaction CO (g) + O- = CO2 (g); K2 =P CO2/P CO [% O] [2] which had been obtained by Vacher and Hamilton2 as an incidental part of their study of the reaction of carbon with oxygen dissolved in the melt. It was found that the equilibrium ratio k1 was not a true constant but varied with the oxygen concentration in the manner shown in Fig. 1, which is copied from the original paper. This early observation has been refuted in subsequent work but is confirmed in the present paper, and an explanation of the discrepancy can now be presented. In Fig. 1, the two points on the extreme right were disregarded in drawing the

Jan 1, 1959

By T. J. Koppenaal, M. E. Fine

The critical resolved shear stress of aCu-A1 single crystals has been investigated as a function of composition, testing temperature, and strain rate. The strength, and its temperature and strain rate dependency increase markedly with aluminum content. The solid solution strengthening at 4.2°K is much larger than at room temperature. The activation energy for the low temperature dislocation cutting process increases from about 0.17 ev for a 0.5 at. pct A1 alloy to about 0.50 ev for a 14 at. pct A1 alloy. This result is attributed to an increase in the width of extended dislocations due to the known decrease in the stacking fault energy. In addition, the activation volumes at 4.2 and 77°K decrease with aluminum additions. Concerning the solid solution strengthening in the 296° to 373°K range, no definite conclusions are reached, but it is suggested that subgrains and cutting of subgrains are important. In a previous paper,&apos; the authors reported that adding aluminum to copper produces two types of strengthening: a yield point effect and an increase in the lower yield stress. This behavior is illustrated in Fig. 1 by shear stress-shear strain curves of furnace-cooled single crystals of pure copper and Cu-14 at. pct A1. The yield point effect, At, was thoroughly studied previously1 and shown to be associated with short-range order. The present paper reports an investigation of the lower-yield stress in aCu-A1 single crystals as a function of composition, testing temperature, and strain rate. EXPERIMENTAL PROCEDURE Single cryski1 tensile specimens were prepared in the same manner as previously described.l Briefly,

Jan 1, 1962

By S. R. Maloof

In 1945, Hollomon&apos; showed that after plastic yielding and prior to necking under simple tension, both ferrous and nonferrous materials are approximated by an equation of the following form: where, a = true stress K = strength coefficient 5 = true strain n = strain hardening exponent If a material is stress free, a plot of log o vs log 6 would be expected to be a straight line, the slope of which is the strain-hardening exponent (n) and the intercept on the ordinate axis for 6 = 1. is the logarithm of the strength coefficient K. It can be easily shown that the strain-hardening exponent in the above equation is equal to the total uniform elongation prior to necking. Thus, one can determine the total uniform elongation prior to necking from such a plot and, as pointed out by C. W. MacGregor, 2 it is the amount of strain prior to necking which determines the extent to which a material may be deep drawn. In the view of the current interest in the fabrication of beryllium shapes, it was thought that a knowledge of the strain-hardening exponent of cross-rolled beryllium sheet would be of general interest. Load-strain data after yielding and prior to necking was taken from the work of Greenspan,3 converted to true stress and true-plastic strain, and plotted on log-log paper as shown in Fig. 1. The value of Young&apos;s modulus E for beryllium was taken as 43 . l06 psi. The slope of this plot gave a value of 0.212 for the strain-hardening exponent (n) and the intercept on the ordinate axis a value of 105,000 for K. The total uniform elongation is thus 21.2 pct, in excellent agreement with the value estimated by Greenspan (about 20 pct). Furthermore,

Jan 1, 1960

By Masayuje Otagawa

The mining methods adopted in Japanese mines are less known to the mining world than those of other countries, owing to the geographical remoteness, but they present many features of interest to mining engineers. The Ashio copper mine is at the town of Ashio, in Tochigiken, about 109 mi. (175 km.) from Tokyo by the Imperial Government railroad. The area of the concession owned by the Furukawa Mining Co. is 5000 acres (5023 ha.). The mine was formerly divided into three, viz. Hon-ean, Kotaki, and Tsudo, which are distant 3 mi. (4.8 km.) from each other as shown in Fig. 1. However, in 1914, when the writer was general manager, the three mining plants were amalgamated into one with the result of simplifying the working projects and increasing the operating efficiency. The mine was visited by the Institute party in November, 1911, when K. Inouye was the general manager. Shortly afterward, the writer was appointed manager, and, in 1915, he was succeeded by Dr. K. Asano. At the end of 1917, K. Kibe became general manager and, after two years, was succeeded by the present manager, I. Sugimoto. For some time Dr. J. Kojima was acting general manager. It is necdless to say that the gentlemen mentioned are all members of the Institute, as is also Baron T. Furukawa, president of the company. Since the discovery of the enormous masses, called "kajika" or "bullhead" deposits, which the writer has described elsewhere,&apos; the mining methods have been varied to meet the different conditions presented by these deposits. The Ashio Ore The country rock of this mine consists of the middle Chichibian slate and quartzite into which has been intruded a post-Tertiary plagio-liparite neck more than 2 mi. in diameter, in the shape of an inverted cone which narrows downward to a relatively small diameter in depth.

Jan 1, 1923

By Masayuje Otagawa

The mining methods adopted in Japanese mines are less known to the mining world than those of other countries, owing to the geographical remoteness, but they present many features of interest to mining engineers. The Ashio copper mine is at the town of Ashio, in Tochigiken, about 109 mi. (175 km.) from Tokyo by the Imperial Government railroad. The area of the concession owned by the Furukawa Mining Co. is 5000 acres (5023 ha.). The mine was formerly divided into three, viz. Hon-ean, Kotaki, and Tsudo, which are distant 3 mi. (4.8 km.) from each other as shown in Fig. 1. However, in 1914, when the writer was general manager, the three mining plants were amalgamated into one with the result of simplifying the working projects and increasing the operating efficiency. The mine was visited by the Institute party in November, 1911, when K. Inouye was the general manager. Shortly afterward, the writer was appointed manager, and, in 1915, he was succeeded by Dr. K. Asano. At the end of 1917, K. Kibe became general manager and, after two years, was succeeded by the present manager, I. Sugimoto. For some time Dr. J. Kojima was acting general manager. It is necdless to say that the gentlemen mentioned are all members of the Institute, as is also Baron T. Furukawa, president of the company. Since the discovery of the enormous masses, called "kajika" or "bullhead" deposits, which the writer has described elsewhere,&apos; the mining methods have been varied to meet the different conditions presented by these deposits. The Ashio Ore The country rock of this mine consists of the middle Chichibian slate and quartzite into which has been intruded a post-Tertiary plagio-liparite neck more than 2 mi. in diameter, in the shape of an inverted cone which narrows downward to a relatively small diameter in depth.

Jan 1, 1923

By Martin F. Littmann

Inclusions of MnS and Fe3C have been introduced into single crystals of 3 pct Si-Fe to study their effects on recrystallization behavior and textures after cold rolling and annealing. The presence of MnS in (110) [001] and (111)[112] crystals inhibited primary grain growth and promoted secondary recrystallization but did not alter the texture significantly after annealing at 1200°C. The presence of Fe3C in (llO)[OOl] and (100)[001] crystals caused a refinement of the primary re crystallized grain size but did not promote secondary recrystallization. THE texture behavior of single crystals of 3 pct Si-Fe during deformation and recrystallization has been studied by numerous investigators. The early work of Dunn&apos; followed by Decker and Harker2 involved relatively small cold reductions. More detailed studies of Dunn3&apos;4 and of Dunn and Koh5&apos;6 involved a reduction of 70 pct and recrystallization at 980°C for several crystals. Walter and Hibbard7 studied a greater variety of initial orientations and sought to relate the textures to those of polycrystalline material. Attention was focused on the nucleation process during early stages of annealing and on surface energy effects in studies by Walter and Dunn8 and by HU.9&apos;10 One of the most extensive investigations has been reported by T. Taoka, E. Furubayashi, and S. Takeuchi.11 Most of this work has been conducted using relatively pure crystals with minimal amounts of precipi-tate-forming elements such as carbon, oxygen, sulfur, and nitrogen. Recently, however, S. Taguchi and A. Sakakura have observed that AIN precipitates can alter the recrystallization textures of rolled (100)[001] crystals.12 The present studies were initiated to determine effects of MnS and Fe3C precipitates on recrystalli-zation and grain growth behavior of rolled single-crystals of 3 pct Si-Fe. Both of these types of inclusions play significant roles in the recrystallization behavior leading to the formation of the (110)[001] or cube-on-edge texture in commercial grain-oriented silicon iron. It is well known that (110)[001] primary grains are formed by recrystallization of (110)[001] or (11 l)[ 112] crystals after cold reduction of about 60 pct or more. Crystals of these orientations, therefore, were selected for study of the effect of MnS in-clusions on grain growth. On the other hand, a major component of the texture of cold-rolled, polycrystal-line 3 pct Si-Fe is the (100)[011] orientation. The function of Fe3,C inclusions is of interest for this orientation. EXPERIMENTAL PROCEDURE The single crystals used are listed in Table I and were obtained from commercial Si-Fe alloy processed to produce (110)[001] and (100)[001] texture by secondary growth. The cube-on-edge material was 0.59 mm thick. Suitably large (110)[001] crystals 25 mm wide were selected and their orientations were determined using an optical goniometer. Etch pits for texture determination were formed by a ferric sulfate solution. The other crystals used in the study with (100)[001], (100)[011], and (111)[112] orientations were obtained from sheet which contained large grains developed from secondary recrystallization by a surface-energy driving force.13 Most crystals had a (100) plane very nearly parallel to the sheet surface and the rolling direction could be selected readily. The same sheet also contained a few crystals with (111) planes parallel to the sheet surface, these also being a result of growth by surface energy. The crystals selected from the sheet were about 25 mm wide and 0.25 to 0.28 mm thick. As shown in Table 11, the crystals already contained about 0.070 to 0.10 pct Mn. Inclusions of MnS were incorporated into crystal 36 in the following manner. The crystals were first sulfurized by holding them Table I. Initial Orientations of Crystals Crystal No. Initial Orientation Thickness, mm Special Treatment 34 (I10) [00l]* 0.59 None 36s (110) [001] 0.59 Sulfide precipitates added 30,40 (111)[Ti21 0.28 None 43s (III) [Ti21 0.28 Sulfide precipitates added 37 (100) [Oll] 0.30 None 37C (100) [01I] 0.27 Carbon added 41 (100) (01I] 0.25 None 41C (100) [OI11 025 Carbide precipitates added 42 (100) [OOl] 0.25 None 42C (100) [001] 0.25 Carbide precipitates added *Tilted 4 deg to r~ght about R.D. Table II. Compositions of Crystals Special Treatments Base Analysis ~ ______________________£________________Crys- Crystals Pct Si Pct C Pct Mn Pct S Pct N Pct Al tal Pct C Pct S 34.36 2.93 • 0.099 <0.005 - 0.0014 36S 0.011 30.37 to 42 2.78 0.0057 0.070 0.001 0.0008 0.0011 43S 0.022 37C 0.029 -41C 0.028 -42C 0.026 *Estimate 0.004 pct. Oxygen estimated <0.003 pct on all samples

Jan 1, 1970

Executive Committee. CHESTER F. LEE, Chairman. C. R. CLAGHORN, Vice-Chairman. WC. BUTLER. M. K. LODGERS. Joseph DANIELS, Secretary-Treasurer. Meeting, Jai. 25, 1913.-Dean Milnor Roberts read a paper on The Equipment and Curricula of the College of Mines, University of Washington.

Jan 4, 1913

United Engineering Trustees, Inc. H. G. MOULTON, 40 A. L. QUENEAU, &apos;38 ALBERT ROBERTS, &apos;39 The Engineering Foundation GEORGE D. BARRON, &apos;40 A. L. QUENEAU, &apos;38 Engineering Societies Library Board H. K. MASTERS, &apos;38&apos; A. B. PARSONS, Ex-officio F. F. SHARPLESS, &apos;37

Jan 1, 1936

By John Chipman, Carl F. Floe

The system molten copper-oxygen-sulphur is of interest from both the practical and theoretical standpoints; practically, because oxygen and sulphur play an important role in the commercial production of copper, and theoretically because it offers a chance to study the reaction between an oxide gas, SO2, and a molten metal. Reactions of this type, such as that between CO or CO2 and molten iron, are important in many metallurgical operations. Again, the determination of any gas-metal equilibrium is of theoretical interest and value insofar as it contributes to the general development of the laws for solution of gases in metals. This problem is being given an increasing amount of attention in modern metallurgy. This paper is the first of a series and presents the results of a determination of the solubility of sulphur dioxide in molten copper at temperatures from 1100" to 1500ºC., and pressures from about 20 mm. of mercury to somewhat above atmospheric. Review of Previous Work The solubility of gases in most liquids has been found to obey Henry&apos;s law, which states that for a fixed temperature the amount of gas dissolved is directly proportional to the pressure. However this law does not hold for liquid metals. Instead, it appears that for many gases, particularly the simple diatomic ones such as O2, N2 and H2, the amount dissolved in liquid metals is directly proportional to the square root of the pressure. This is known as Sieverts&apos; law. It may be recognized as a special case of Henry&apos;s law, which applies when the solute gas is dissociated into atoms or when it reacts with the solvent metal to form solute molecules containing single atoms of the gaseous elements, as, for example: 2M+G2=2MG K=(MG)2 /M2 X PG2 [I] which, if the solubility is not too high and the concentration of metal therefore nearly constant, can be simplified to: MG = K1 vPG2 [2 Following the same reasoning, it would be expected that the solubility of SO2 in copper would be proportional to the cube root of the pressure, the reaction being as follows: 6Cu + SO2 = Cu2S + 2Cu2O [3] K= Cu2S X (Cu2O)2 i - (Cu)6(Pso2) [4] or, neglecting the change in metal concentration and considering that the concentrations of sulphur and oxygen are both proportional to the volume V of gas absorbed: K1 =V3SO2 /Pso2 or

Jan 1, 1941

By Carl F. Floe, John Chipman

The system molten copper-oxygen-sulphur is of interest from both the practical and theoretical standpoints; practically, because oxygen and sulphur play an important role in the commercial production of copper, and theoretically because it offers a chance to study the reaction between an oxide gas, SO2, and a molten metal. Reactions of this type, such as that between CO or CO2 and molten iron, are important in many metallurgical operations. Again, the determination of any gas-metal equilibrium is of theoretical interest and value insofar as it contributes to the general development of the laws for solution of gases in metals. This problem is being given an increasing amount of attention in modern metallurgy. This paper is the first of a series and presents the results of a determination of the solubility of sulphur dioxide in molten copper at temperatures from 1100" to 1500ºC., and pressures from about 20 mm. of mercury to somewhat above atmospheric. Review of Previous Work The solubility of gases in most liquids has been found to obey Henry&apos;s law, which states that for a fixed temperature the amount of gas dissolved is directly proportional to the pressure. However this law does not hold for liquid metals. Instead, it appears that for many gases, particularly the simple diatomic ones such as O2, N2 and H2, the amount dissolved in liquid metals is directly proportional to the square root of the pressure. This is known as Sieverts&apos; law. It may be recognized as a special case of Henry&apos;s law, which applies when the solute gas is dissociated into atoms or when it reacts with the solvent metal to form solute molecules containing single atoms of the gaseous elements, as, for example: 2M+G2=2MG K=(MG)2 /M2 X PG2 [I] which, if the solubility is not too high and the concentration of metal therefore nearly constant, can be simplified to: MG = K1 vPG2 [2 Following the same reasoning, it would be expected that the solubility of SO2 in copper would be proportional to the cube root of the pressure, the reaction being as follows: 6Cu + SO2 = Cu2S + 2Cu2O [3] K= Cu2S X (Cu2O)2 i - (Cu)6(Pso2) [4] or, neglecting the change in metal concentration and considering that the concentrations of sulphur and oxygen are both proportional to the volume V of gas absorbed: K1 =V3SO2 /Pso2 or

Jan 1, 1941

By M. J. Pool, P. J. Spencer, J. R. Guadagno

THE present work was carried out as part of a general program to study the magnitude and concentration dependence of exothermic heats of solution of various metals in liquid tin. Thermodynamic data for the Te-Sn system are very limited, particularly for dilute alloys. Indeed, apart from the prior heat of solution determinations of Pool,1 the only other published values are those of McAteer and seltz2 who carried out electromotive force measurements on solid alloys of compositions from 75 to 95 at. pct Te at temperatures between 543° and 668° K, of Colin and Dro-wart3 who made mass-spectrometric studies of the vapor in equilibrium with solid SnTe, and of Robinson and Bever4 who determined the heat of formation of SnTe by solution calorimetry. The present work thus provides further thermodynamic information for this system, specifically in a temperature range where the solubility of tellurium is very limited. The liquid metal solution calorimeter used in this investigation has been described previously.&apos; No further discussion of its design or operation will therefore be given here. For each series of measurements a tin bath consisting of approximately 70 g of 99.99 pct pure Sn was carefully weighed out and inserted in the calorimeter which was maintained at the desired temperature to within ± 1/4°c. Tellurium, of 99.999 pct purity, was supplied by Atomergic Chemetals Co., a division of Gallard-Schlesinger Chemical Manufacturing Corp. Approximately 0.0006 to 0.001 g-atom specimens of tellurium were used for each experimental run and these were added to the calorimeter from a temperature of 273" K. Heat of solution values were calculated by analysis of a strip-chart recorder plot of the heat effect accompanying the solution of each sample in the tin bath. The heat capacity of the latter was determined at regular intervals during a series of tellurium drops by using pure tin or tungsten calibration samples, the heat contents of which were obtained from the standard compilation of Hultgren et al5 The composition of the bath was calculated after each drop from a knowledge of the weights of tin and tellurium added. Values of the heat of solution of tellurium at 725" and 775° K are listed in Table I together with the corresponding tellurium contents of the alloys. The data are also plotted as a function of tellurium concentration in Fig. 1, which includes the equivalent data obtained by pool&apos; at 750° K. It can be seen that the heat of solution remains

Jan 1, 1969

United Engineering Trustees, Inc. WILLIAM H. BASSETT H. G. MOULTON R. M. ROOSEVELT The Engineering Foundation GEORGE D. BARRON E. DEGOLYER R. M. ROOSEVELT Engineering Societies Library Board ROBERT PEELE H. K MASTERS A. B. PARSONS, Ex-officio F. F. SHARPLESS R. H. SWEETEST

Jan 1, 1934