The diffusion of antimony in alpha iron

Diffusion coefficients of antimony in α-iron were determined in the temperature range 700 to 900°C using the residual activity method. Specimens were large-grained polycrystals for the higher temperature measurements and single crystals for the low temperature measurements. Above 800°C the data may be represented by the equationD _sb(cm²/s) = (440 ± 200) exp [- (270,000 ± 7000)/RT]. The activation energy (reported in J/mole) is approximately equal to that measured for iron self-diffusion in this same temperature range, although the antimony diffusion coefficients are a factor of ten larger than the iron self diffusion coefficients. The potential for strongly coupled vacancy-antimony motions is demonstrated, based on the observed enhancement of iron self diffusion in dilute iron-antimony alloys. Finally molybdenum is shown to have a negligible effect on the diffusion of antimony in α-iron. These results are discussed in relation to the phenomenon of temper brittleness in steels. Embrittlement kinetics in iron-antimony alloys are shown to be consistent with an antimony diffusion controlled segregation mechanism.     

The self diffusion of iron in Fe2SiO4 and CaFeSiO4 melts

The self diffusion of iron in Fe₂SiO₄ and CaFeSiO₄ melts has been measured in the temperature range 1250° to 1540°C using Fe⁵⁹ as the radio tracer and the capillary-liquid reservoir method of diffusion measurement. The results obtained are represented by $$log D_{Fe} = - \frac{{3800 \pm 500}}{T} - 2.74 \pm 0.29$$ for Fe₂SiO₄, and $$log D_{Fe} = - \frac{{5450 \pm 620}}{T} - 1.93 \pm 0.37$$ for CaFeSiO₄. Excellent agreement is obtained with the self-diffusivity of iron calculated from the measured interdiffusivity of iron and oxygen in iron oxide melts.     

Radiochemical measurement of diffusion in the cesium chloride-tin and barium-sodium systems

Diffusion coefficients for cesium chloride in molten tin from 308 to 396°C and for barium in liquid sodium from 290 to 396°C were measured radiochemically. Cesium chloride coefficients were determined by measurement of nonsteady state surface concentrations of cesium chloride deposited on the surface of cylindrical specimens. Cesium chloride traced with cesium-137 was deposited on the specimens from a supersaturated solution dispersed in the molten tin and measured by gamma spectroscopy after removal of the specimen from the solution. Diffusion coefficients for barium in sodium were measured by a unique method in which radioactive cesium-137 (half-life of 33 yrs) dissolved in the sodium served as a volumetric source for its short-lived barium-137 daughter (half-life of 2.6 min). Since barium deposited on the surface of the specimen, measurement of the equilibrium barium-137m gamma activity on the specimen surfaces after removal from the solution yielded barium diffusion coefficients. The equations for these diffusion coefficients with 95 pct confidence limits are: Cesium chloride in tin (308 to 396°C):D=[(4.84±0.14)×10^?4] exp [?6720±6400)/RT] Barium in sodium (290 to 496°C):D=[(4.27±0.05)×10^?4] exp [?4600±990)/RT]     

Dislocation creep controlled superplasticity in the high carbon steel 140NiCr16-6

Superplasticity in the alloyed high carbon-steel 140NiCr16-6 with phosphorus additions and a fine grained microdupiex structure – containing cementite in volume fractions of 22 % (Fe,Cr,Ni)₃C, particle size of about 1 μm and with a medium ferrite grain size of about 2 μm – has been investigated in the temperature regime of 550 to 675°C and in the strain rate range of 10^?5 to 5 · 10^?2 s^?1. Maximum strain rate exponents of m = 0,45 at 675°C with strain rates of the order of 10^?4 s^?1 have been determined. Maximum superplastic elongations of about 700 % were detected. At higher strain rates of 10^?3 s^?1 superplastic elongations of about 570 % were achieved. At relatively low test temperatures of 550°C elongations up to 230 % were recorded. The activation analysis in the temperature regime of 550 to 650°C show an activation energy for superplastic flow of 250 ± 20 kJ/mol. This is in agreement with the activation energy for lattice self diffusion of iron in α-iron. Above 650°C the activation energy decreases to 70 kJ/mol. This is due to a stress induced decrease in the eutectoid α-γ-transformation temperature from 685°C to somewhat lower temperatures during superplastic deformation. The superplastic deformability (m > 0.3) of this steel in a wide strain rate range at relatively low temperatures above 550°C allows near net shape forming of complex parts applying low flow stresses.     

Determination of diffusion coefficient of oxygen in γ-iron from measurements of internal oxidation in Fe-Al alloys

The internal oxidation of iron alloys containing between 0.069 and 0.274 wt pct aluminum was investigated in the temperature range from 1223 to 1373 K for the purpose of determining the diffusion coefficients in γ-iron as well as in the internal oxidation layer. A parabolic rate law is obeyed in the internal oxidation of the present alloys. The rate constant for penetration of the oxidation front, the oxide formed, and the concentration of aluminum in the oxidation layer were determined. Pronounced enrichment of aluminum in the oxidation layer was observed, resulting from the counterdiffusion of aluminum. The oxygen concentration at the specimen surface was determined by combining the thermodynamic data on the dissociation of FeO and the solution of oxygen in y-iron. The diffusion coefficient of oxygen in the internal oxidation layer,D _o ¹⁰ , was evaluated on the basis of the rate equation for internal oxidation.D _o ¹⁰ increases at a given temperature as the volume fraction of oxide,f ¹⁰, in the oxidation layer increases. The diffusion coefficient of oxygen in γ-iron,D _o, was determined by extrapolation ofD _o ¹⁰ = 0.D _o may be expressed as $$D_o = \left( {1.30\begin{array}{*{20}c} { + 0.80} \\ { - 0.50} \\ \end{array} } \right) \times 10^{ - 4} \exp \left[ { - \frac{{166 \pm 5(kJ \cdot mol^{ - 1} )}}{{RT}}} \right]m^2 \cdot s^{ - 1} .$$ D _o is close to the diffusion coefficients of carbon and nitrogen in γ-iron.     

Desulfurization of liquid steel containing aluminum or silicon with lime

The reaction mechanisms for the desulfurization of iron containing from 0.04 to 1.5 pct aluminum or 1.1 to 3.7 pct silicon by CaO at 1600° have been examined. The desulfurization of Fe-Al by CaO is considerably faster than that of Fe-Si. The basic difference between the two processes is that Fe-Al alloys can be desulfurized by the formation of AI₂O₃, whereas for Fe-Si melts it is necessary to form Ca₂SiO₄. The rate of desulfurization of Fe-Si alloys by CaO is controlled by the formation of the gaseous intermediates, SiS and S, and is the same as that for desulfurization in vacuum. The rate of desulfurization of Fe-Al melts is fast, and is apparently controlled by the diffusion of sulfur to the liquid metal—CaO interface. Experiments were also conducted to demonstrate that sulfur could be transferred to CaO by the gaseous intermediates SiS and S.     

The Diffusion Coefficient of Scandium in Dilute Aluminum-Scandium Alloys

The diffusion coefficient of Sc in dilute Al-Sc alloys has been determined at 748 K, 823 K, and 898 K (475 °C, 550 °C, and 625 °C, respectively) using semi-infinite diffusion couples. Good agreement was found between the results of the present study and both the higher temperature, direct measurements and lower temperature, indirect measurements of these coefficients reported previously in the literature. The temperature-dependent diffusion coefficient equation derived from the data obtained in the present investigation was found to be \( D \left( {{\text{m}}^{2} /{\text{s}}} \right) = \left( {2.34 \pm 2.16} \right) \times 10^{ - 4} \left( {{\text{m}}^{2} /{\text{s}}} \right) { \exp }\left( {\frac{{ - \left( {167 \pm 6} \right) \left( {{\text{kJ}}/{\text{mol}}} \right)}}{RT}} \right). \) Combining these results with data from the literature and fitting all data simultaneously to an Arrhenius relationship yielded the expression \( D \left( {{\text{m}}^{2} /{\text{s}}} \right) = \left( {2.65 \pm 0.84} \right) \times 10^{ - 4} \left( {{\text{m}}^{2} /{\text{s}}} \right) { \exp }\left( {\frac{{ - \left( {168 \pm 2} \right) \left( {{\text{kJ}}/{\text{mol}}} \right)}}{RT}} \right). \) In each equation given above, R is 0.0083144 kJ/mol K, T is in Kelvin, and the uncertainties are ±1 standard error.     

Structure and properties of rapidly solidified dispersion-strengthened titanium alloys: Part I. Characterization of dispersoid distribution,structure, and chemistry

The feasibility of developing dispersion-strengthened powder metallurgy Ti alloys was determined in Ti-RE (RE = Ce, Dy, Er, Gd, La, Nd, or Y) alloys prepared by rapid solidification processing. The alloys were produced by electron-beam melting and splat quenching. Dispersoid precipitation and growth were studied as functions of annealing temperature, 700 to 1000 °C, for annealing times between 5 and 50,000 minutes. Dispersoid diameters, spacings, compositions, and crystal structures were characterized by transmission and scanning electron microscopy, X-ray and electron diffraction, energy-dispersive X-ray analysis, and scanning Auger microscopy. Two classes of dispersoid coarsening behavior at temperatures below theβ-transus were identified. In Ti-Ce, Ti-Gd, and Ti-Nd alloys, equilibrium rare earth sesquioxide (RE₂O₃) dispersoids form early in the annealing process and coarsen rapidly to > 1 μm diameter. The Ti-Nd alloys additionally contain large volume fractions of small (< 100 nm diameter) dispersoids. In the other Ti-RE alloys, dispersoids identified as Ti-RE-O-C compounds coarsen relatively slowly. Ti-Er is the most promising of the investigated systems for application in a multicomponent dispersion-strengthened alloy because long-time annealing at 700 to 800 °C produces stable dispersoids of 50 to 150 nm average diameter and 300 to 600 nm inter-particle spacing.     

High Temperature Vaporization Chemistry in the Gold-Chlorine System Including Formation of Vapor Complex Species of Gold and Silver with Copper and I ron

The thermodynamic properties of the vaporization reactions in the gold-chlorine system have been investigated in the temperature ranges of 580 to 649 K and 882 to 1107 K. The experimental technique consisted of classical transpiration vapor pressure measurement and analysis with a newly developed Flow Reactor-Mass Spectrometer system, which connects a high temperature, ambient pressure reactor to a TOF mass spectrometer. Metallic gold reacts with chlorine gas at temperatures above 750 K, forming the vapor species Au₂Cl₂(g). The resulting reaction 2Au(c) + Cl₂(g) = Au₂Cl₂(g) was found to have a standard Gibbs free energy change of ΔG° = 89,000 (±2400) ? 27.8 (±1) T (J/mole) ΔG° = 21,300 (±570) ? 6.6 (±0.2) T (cal/ mole) in the temperature range 882 to 1107 K. At lower temperatures, the vaporization reaction is 2Au(c) + 3Cl₂(g) = Au₂Cl₆(g) with a standard Gibbs free energy change of ΔG° = ?102,000 (±20,000) + 235 (±29) T (J/mole) ΔG° = ?24,400 (±4800) +56 (±7) T (cal/mole) in the temperature range 580 to 649 K. The above results are combined with equilibrium dissociation data obtained in the literature for condensed gold chloride phases to construct a phase stability-vapor pressure diagram for the gold-chlorine system. Consideration is given to some possible operating conditions for a gold chlorination-vaporization process to treat low grade and refractory gold ores. Transpiration vapor pressure measurements in the Au-Cu-Cl, Ag-Cu-Cl, and Ag-Fe-Cl systems showed significant enhancement of the vapor pressure of Au₂Cl₂(g) or AgCl(g) in the presence of Cu₃Cl₃(g) or FeCl₃(g). This indicates the formation of binary vapor phase complexes in these systems. The species AuCu₂Cl₃(g), AgCu₂Cl₃(g), and AgFeC₄ have been proposed.     

A mass-spectrometric study of the thermodynamics of the Fe−Cu and Fe−Cu−C(sat) systems at 1600°C

The Knudsen cell-mass spectrometer combination has been used to study the Fe?Cu and Fe?Cu?C_(sat) alloys at 1600°C. Activity coefficients in the Fe?Cu system are closely represented by the equations $$\begin{gathered} \ln \gamma _{Fe} = 1.86N_{Cu}^2 + 0.03, (0< N_{Fe}< 0.7) \hfill \\ \ln \gamma _{Cu} = 2.25N_{Fe}^2 - 0.19, (0.7< N_{Fe}< 1.0) \hfill \\ \end{gathered} $$ with an uncertainty in the quadratic terms of about 5 pct. For the iron-rich carbon-saturated alloys, the activity coefficient of copper is given by the equation $$\ln \gamma _{Cu} = 2.45(N'_{Fe} )^2 + 0.3N'_{Fe} + 0.03, (0< N'$$ to within an uncertainty of about 10 pct. N _Fe ^′ represents the fraction N_Fe/(N_Fe+N_Cu), etc. The activity coefficient of iron in this region is found to be essentially constant at 0.69±0.05.     

Rates and mechanisms of FeO reduction from slags

The rate of FeO reduction from CaO?SiO₂ slags has been determined in a stationary magnesia crucible with a graphite rod as the reductant. At 1650°±35°C, with a slag basicity CaO/SiO₂=1.2±0.3, and in the total iron concentration,C, range of 10 pct<C<40 pct, the specific rate,Q _E , of reduction in g-moles per min per sq cm was found to beQ _E =8.25×10^?6 C ^1.77 According to this and previous literature results, the reduction rates span three possible stages. The order of reaction in regard to concentration varies from first order in the low concentrations (0.1 pct<C<1.5 pct) to second order in the intermediate concentrations (1.5 pct<C<15 pct) and finally to 1.77th order in the high concentrations (15 pct<C<40 pct). The reduction reaction can best be interpreted by convective mechanisms because the observed rates fall between the rates of the chemically limited graphite gasification and the rates predicted from molecular diffusion.     

Influence of Phase Transitions on the Macrokinetics of the Gaseous Reduction of Iron Oxide

The reduction of magnetite pellets is studied by thermogravimetry in a hydrogen flow upon linear heating. A multiple decrease in the reduction rate is observed at the degrees of metallization higher than 70% in a temperature range of 920–950°C. The observed effect of a decrease in the reduction rate by four to five times on heating is related to the α → γ phase transition of iron. Temperature cycling (heating–holding–heating–holding…) in the phase transition range shows that this effect is reversible; i.e., the reduction rate increases on cooling, unlike that on heating. The average temperature of the beginning of the effect in heating–cooling cycles (913°C) turns out to be close to the thermodynamic temperature (911°C) of the α → γ phase transition of iron. Isothermal studies of the reduction rate of pellets in a temperature range of 900–1000°C also confirm this phenomenon. The effect of a decrease in the reduction rate is assumed to appear due to a decrease in the effective coefficient of gas diffusion in γ-iron as compared to α-iron.     

Grain boundary segregation in Ni and binary Ni alloys doped with sulfur

The kinetics and temperature dependence of sulfur segregation in Ni and binary alloys of Ni with Cu, Al, Cr, Mo, W, and Hf, all containing between 70 and 100 atomic ppm sulfur, have been measured using Auger spectroscopy over the temperature range 500 to 1000 °C. No evidence for cosegregation of these alloying elements with sulfur is found. The alloying elements do influence the precipitation of sulfides, however, and this influences the amount of segregation which occurs. Comparison with theoretical models of grain boundary segregation allows the sulfur solubility in the various alloys to be determined.     

Dynamic strain aging in carbide-strengthened molybdenum alloys

A study was made of the dynamic strain aging processes in carbide strengthened molybdenum alloys. Strengthening due to dynamic strain aging was observed in the temperature range 1500° to 2400°F. The strengthening is believed to arise from the pinning of mobile dislocations by metal-carbon pairs or clusters and/or carbide precipitates. This immobilization of the dislocations led to a high work hardening rate and the observed strengthening.     

The undercooling of aluminum

An important parameter affecting microstructure development during solidification is the amount of undercooling prior to nucleation. The undercooling potential of aluminum has been assessed by thermal analysis measurements on powder dispersions of the liquid metal. A number of variables have been identified which influence the undercooling of powder Al samples including powder coating, powder size, melt cooling rate, and melt superheat. Surface analysis by Auger electron spectroscopy indicates that changing the medium in which the powders are produced is an effective method of altering the coating chemistry. Factorial design analysis has been employed to quantify the potential of processing variables to increase the undercooling level obtainable in aluminum. The factorial analysis indicates that control of the powder coating through changing the medium in which the powders are produced is most effective in decreasing the nucleation temperature. Additionally, the finest powders produced in the medium which induces the least catalytic coating, when cooled at high rates,T = 500 °C/s, from low superheats,T _s = 710 °C, are found to achieve the deepest undercooling, ΔT = 175 °C. These studies provide the basis for further increases in undercooling and for future investigations into the solidification reactions which produce both stable and metastable structures in aluminum alloys.     

Activity coefficient of antimony and arsenic in molten iron and carbon saturated iron

In connection with removal of tramp elements from molten iron and steel, activity coefficients of arsenic and antimony in carbon saturated iron and in pure iron were determined between 1473 to 1923 K and between 1823 to 1923 K, respectively, by measuring the distribution of those elements between the metals and silver. The following activity coefficients are observed: for carbon saturated iron: log γ_As = 3560/T + 1.14, log γ_Sb = 6890/T + 5.06; and for molten iron: log γ°_As = - 198/T - 1.41, log γ°_Sb = - 7030/T + 4.38, respectively. The possibility of the removal of arsenic and antimony from molten iron by using a BaO-BaF₂ flux is discussed.     

Rate of change of oxygen chemical potential in liquid PbO−SiO2 binary solution

The rate of the chemical potential change of oxygen in a liquid PbO?SiO₂ binary solution, with SiO₂ contents of 10, 20, and 30 mol pct, and in pure PbO, has been measured at temperatures of 900°, 950°, 1000°, 1050°, and 1100°C. The rate increased with temperature according to the Arrhenius type relation and decreased with the increase of the silica content. It is suggested that the rate-controlling step is the counter diffusion rate of Pb²⁺ and Pb⁴⁺ ions, which are considered to be the most easily movable ions in the PbO?SiO₂ solution. The relation between the rate of oxygen chemical potential change and the electrical conductivity is also discussed for the liquid PbO?SiO₂ system.     

Thermodynamic properties of molten PbO−SiO2 systems

Thermodynamic properties of the molten PbO?SiO₂ systems were determined by electromotive force measurements employing solid electrolytes. The experiments were carried out in the temperature range 900° to 1000°C on the galvanic cells \(Pb_{(l)} + PbO_{(l)} \left| \begin{gathered} Electrolyte \\ ZrO_2 \times CaO \\ \end{gathered} \right|Pb_{(l)} + PbO - SiO_{2(l)} \) From the results the activities of PbO and SiO₂, the partial and integral free energies, entropies, and heats of mixing were determined. These experimental results are interpreted and compared with previous experimental data. Based on the present data and according to the model of Toop and Samis,¹the activity coefficients of the free oxygen ion O^2? concentration is reduced.     

Effect of Sulfur and Antimony on the Intergranular Fracture of Iron at Cathodic Potentials

Antimony, segregated to grain boundaries of iron, was found to be five times more effective than sulfur in promoting intergranular fracture of iron when tested in IN H₂SO₄ at cathodic potentials. A decrease in the ductility of iron accompanied the fracture mode change at increasing cathodic potentials. The effectiveness of antimony relative to sulfur was determined from straining electrode tests on iron and iron + 250 appm antimony alloys heat treated at 800 °C and 600 °C to produce different grain boundary chemical compositions. Grain boundary compositions were determined by Auger Electron Spectroscopy (AES). Similar grain boundary sulfur concentrations of 0.2 monolayers were observed by AES for the iron and iron + 250 appm antimony alloy after an anneal of 240 hours at 600 °C, while 0.08 monolayers of antimony was observed for the iron + 250 appm antimony alloy. These results suggest that sulfur and antimony do not compete for grain boundary sites.     

Thermodynamic Relation between Silicon and Aluminum in Liquid Iron

The thermodynamic relation between silicon and aluminum in liquid iron was studied by measuring the effect of silicon on the solubility product of AlN in liquid Fe-Si-Al-N alloys containing silicon up to 1.5?mass pct in the temperature range from 1823?K to 1923?K (1550 °C to 1650 °C). The effects of aluminum and silicon on nitrogen solubility in liquid iron were separately determined in the same temperature range. The experimental results were thermodynamically analyzed using Wagner??s interaction parameter formalism to determine the first-order interaction parameters of silicon on nitrogen and aluminum in liquid iron as follows: $ e_{\text{N}}^{\text{Si}} = 0. 0 6 7 3, \;e_{\text{Al}}^{\text{Si}} = 0.009 $ (1823?K to 1923?K (1550 °C to 1650 °C), Si ?? 1.5?mass pct)