Electrical conductivity of molten cryolite-based mixtures obtained with a tube-type cell made of pyrolytic boron nitride

A pyrolytic boron nitride tube-type cell was used to measure the electrical conductivity for molten cryolite, for binary mixtures of cryolite with Al₂O₃, AlF₃, CaF₂, KF, Li₃AlF₆, and MgF2, and for ternary mixtures Na₃AlF₆-Al₂O₃-CaF₂ (MgF₂) and Na₃AlF₆-AlF₃-KF (Li₃AlF₆). The cell constant was about 40 cm^?t. The temperature and concentration dependence of the conductivity in the investigated concentration range was described by the equation $$\begin{gathered} \kappa /S cm^{ - 1} = 7.22 exp\left( { - 1204.3/T} \right) - 2.53\left[ {Al_2 O_3 } \right] - 1.66\left[ {AlF_3 } \right] \hfill \\ - 0.76\left[ {CaF_2 } \right] - 0.206\left[ {KF} \right] + 0.97\left[ {Li_3 AlF_6 } \right] - 1.07\left[ {MgF_2 } \right] \hfill \\ - 1.80\left[ {Al_2 O_3 } \right]\left[ {CaF_2 } \right] - 2.59\left[ {Al_2 O_3 } \right]\left[ {MgF_2 } \right] \hfill \\ - 0.942\left[ {AlF_3 } \right]\left[ {Li_3 AlF_6 } \right] \hfill \\ \end{gathered} $$ whereT represents the temperature in Kelvin and the brackets represent the mole fractions of the additions. The standard deviation was found to be 0.026 S cm^?1 (～1 pct). For practical reasons, it is often desired to express composition in weight percent. In that case, it holds that $$\begin{gathered} \ln \kappa = 1.977 - 0.0200\left[ {Al_2 O_3 } \right] - 0.0131\left[ {AlF_3 } \right] - 0.0060\left[ {CaF_2 } \right] \hfill \\ - 0.0106\left[ {MgF_2 } \right] - 0.0019\left[ {KF} \right] + 0.0121\left[ {LiF} \right] - 1204.3/T \hfill \\ \end{gathered} $$ whereT represents the temperature in Kelvin and the brackets denote the concentration of the additives in weight percent. However, in this case, the maximum relative error of the conductivity equation can reach up to 2.5 pct.     

The precipitation of epsilon-carbide in twinned martensite

Tempering of martensite has been investigated by means of thin foil electron microscopy in a high carbon steel, a high nickel steel, and a silicon steel. ε carbide has been unambiguously identified in each steel. It was found that the carbide was precipitated with the Jack orientation relationship: $$\begin{gathered} \left( {0001} \right)_\varepsilon \parallel \left( {011} \right)_{\alpha '} \hfill \\ \left( {10\bar 10} \right)_\varepsilon \parallel \left( {2\bar 11} \right)_{\alpha '} \hfill \\ \end{gathered} $$ In the silicon steel the ε carbide precipitated in the form of needles which grew with a \(\left[ {01\bar 10} \right]_\varepsilon \) close to \(\left[ {21\bar 1} \right]_{\alpha '} \) . This growth direction minimizes the surface energy of the needles, yet allows growth in a direction of low mismatch.     

Self-diffusion of copper in Cu−Al solid solutions

Self-diffusion coefficients of copper in Cu?Al solid solutions in the concentration interval 0 to 19 at. pct Al and in the temperature range 800° to 1040°C have been determined by the residual activity method using the isotope Cu⁶⁴. The values of the self-diffusion coefficients in the concentration interval 0 to 14.5 at. pct Al satisfy the Arrhenius relation and their temperature dependence can be expressed by the following equations $$\eqalign{ & D_{Cu}^{Cu} = \left( {0.43_{ - 0.11}^{ + 0.15} } \right) exp \left( { - {{48,500 \pm 700} \over {RT}}} \right) cm^2 /\sec \cr & D_{Cu - 2.80 at. pct Al}^{Cu} = \left( {0.46_{ - 0.16}^{ + 0.23} } \right) exp \left( { - {{48,000 \pm 900} \over {RT}}} \right) cm^2 /\sec \cr & D_{Cu - 5.50 at. pct Al}^{Cu} = \left( {0.30_{ - 0.07}^{ + 0.09} } \right) exp \left( { - {{47,000 \pm 600} \over {RT}}} \right) cm^2 /\sec \cr & D_{Cu - 8.83 at. pct Al}^{Cu} = \left( {0.46_{ - 0.09}^{ + 0.11} } \right) exp \left( { - {{47,100 \pm 500} \over {RT}}} \right) cm^2 /\sec \cr & D_{Cu - 11.7 at. pct Al}^{Cu} = \left( {0.61_{ - 0.13}^{ + 0.17} } \right) exp \left( { - {{47,200 \pm 600} \over {RT}}} \right) cm^2 /\sec \cr & D_{Cu - 14.5 at. pct Al}^{Cu} = \left( {4.2_{ - 1.5}^{ + 2.2} } \right) exp \left( { - {{51,110 \pm 1000} \over {RT}}} \right) cm^2 /\sec \cr} $$ An analysis of the results leads to the conclusion that, in the concentration interval 0 to 11.7 at. pct Al, the frequency factor and activation enthalpy concentration dependences can be described by the following equations whereD _0Cu ^Cu and ΔH _Cu ^Cu are diffusion characteristics for self-diffusion in pure copper,X _Al is the atomic percent of aluminum, andK andB are experimental constants.     

Reaction Mechanism and Kinetics of Enargite Oxidation at Roasting Temperatures

Roasting of enargite (Cu₃AsS₄) in the temperature range of 648?K to 898?K (375?°C to 625?°C) in atmospheres containing variable amounts of oxygen has been studied by thermogravimetric methods. From the experimental results of weight loss/gain data and X-ray diffraction (XRD) analysis of partially reacted samples, the reaction mechanism of the enargite oxidation was determined, which occurred in three sequential stages:

1. $4{\text{Cu}}_{ 3} {\text{AsS}}_{ 4} \left( {\text{s}} \right){\text{ + 13O}}_{ 2} \left( {\text{g}} \right){\text{ = As}}_{ 4} {\text{O}}_{ 6} \left( {\text{g}} \right){\text{ + 6Cu}}_{ 2} {\text{S}}\left( {\text{s}} \right){\text{ + 10SO}}_{ 2} \left( {\text{g}} \right) $
2. $ 6{\text{Cu}}_{ 2} {\text{S}}\left( {\text{s}} \right){\text{ + 9O}}_{ 2} \left( {\text{g}} \right){\text{ = 6Cu}}_{ 2} {\text{O}}\left( {\text{s}} \right){\text{ + 6SO}}_{ 2} \left( {\text{g}} \right) $
3. $ 6{\text{Cu}}_{ 2} {\text{O}}\left( {\text{s}} \right){\text{ + 3O}}_{ 2} \left( {\text{g}} \right){\text{ = 12CuO}}\left( {\text{s}} \right) $

The three reactions occurred sequentially, each with constant rate, and they were affected significantly by temperature and partial pressure of oxygen. The kinetics of the first stage were analyzed by using the model X?=?k ₁ t. The first stage reaction was on the order of 0.9 with respect to oxygen partial pressure and the activation energy was 44?kJ/mol for the temperature range of 648?K to 898?K (375?°C to 625?°C).     

Thermodynamic behavior in binary metallic solutions

The applicability of Krupkowski’s formalism $$\begin{gathered} ln \gamma _1 = \omega \left( T \right)\left( {1 - X_1 } \right)^m \hfill \\ ln \gamma _2 = \omega \left( T \right)\left[ {\left( {1 - X_1 } \right)^m - \frac{m}{{m - 1}}\left( {1 - X_1 } \right)^{m - 1} + \frac{1}{{m - 1}}} \right] \hfill \\ \end{gathered} $$ in interpreting experimental data is shown for several binary systems. Both dilute and concentrated solutions are considered. In dilute solutions (Henry’s law region) these equations exclude constant values of the activity coefficients. These formulae withm>1 satisfy Raoults law and Henry’s law as limiting cases. However, experimental data indicate that only in two systems, namely Zn-Sn and Zn-Bi,γ _Zn ⁰ =γ _Zn over a finite composition range. Whenm is close to unity, as is the case for the Zn-Sn and Zn-Bi systems Raoult’s law is not satisfied untilX _Zn is infinitesimally close to unity. Data for concentrated zinc solutions for both systems support this conclusion. A comparison of Krupkowski’s method with Darken’s quadratic formalism was also carried out, and it was shown that both methods give similar results whenm=2.     

Lattice and grain boundary diffusion of cerium and neodymium in nickel

Diffusion of cerium and neodymium in nickel has been studied by the serial sectioning technique using radioactive tracers¹⁴¹Ce and¹⁴⁷Nd, in the temperature ranges 700° to 1100°C for volume and 500° to 875°C for grain boundary diffusion respectively. Volume diffusivities can be expressed as: $$\begin{gathered} D_{Ce/Ni} = (0.66 \pm 0.18)\exp \left( { - \frac{{60,800 \pm 810}}{{RT}}} \right)cm^2 /\sec \hfill \\ D_{Nd/Ni} = (0.44 \pm 0.13)\exp \left( { - \frac{{59,820 \pm 830}}{{RT}}} \right)cm^2 /\sec \hfill \\ \end{gathered} $$ and grain boundary diffusivities by: $$\begin{gathered} Dg_{Ce/Ni} = 0.11\exp \left( { - \frac{{29,550}}{{RT}}} \right)cm^2 /\sec \hfill \\ Dg_{Nd/Ni} = 0.07\exp \left( { - \frac{{28,580}}{{RT}}} \right)cm^2 /\sec \hfill \\ \end{gathered} $$ Results of volume diffusion have been compared with those calculated from the theories of diffusion based on size and charge difference between the solute and the solvent atoms. Whipple and Suzuoka methods have been used to evaluate the grain boundary diffusion coefficients. Both the methods give similar results.     

The thermodynamic properties of liquid Fe−Si alloys

The thermodynamic properties of liquid Fe?Si alloys have been determined electrochemically by use of the following galvanic cells: $$\begin{gathered} Cr - Cr_2 O_3 (s)|ZrO_2 (CaO)|Fe - Si(l), SiO_2 (s) \hfill \\ Cr - Cr_2 O_3 (s)|ThO_2 (Y_2 O_3 )|Fe - Si(l), SiO_2 (s) \hfill \\ \end{gathered} $$ The free energy of formation of SiO₂ was measured and is ?139.0 and ?134.3 kcals per mole at 1500° and 1600°C, respectively. The activity coefficients of iron and silicon for the atom fraction of siliconN _Si<0.35 at 1600° and 1500°C can be represented by the quadratic formalism. $$\begin{gathered} \left. {\begin{array}{*{20}c} {log \gamma _{Fe} = - 2.12 N_{Si}^2 } \\ {log \gamma _{Si} = - 2.12 N_{Fe}^2 - 0.22} \\ \end{array} } \right\}1600^ \circ C (2912^ \circ F) \hfill \\ \left. {\begin{array}{*{20}c} {log \gamma _{Fe} = - 2.50 N_{Si}^2 } \\ {log \gamma _{Si} = - 2.50 N_{Fe}^2 - 0.13} \\ \end{array} } \right\}1500^ \circ C (2732^ \circ F) \hfill \\ \end{gathered} $$ The results indicate that an excess stability peak occurs at about the equimolar composition. Combining the heats of solution determined in this study with previous data indicates that the heats also follow the quadratic formalism. The partial molar heats, \(\bar L_{Si} \) and \(\bar L_{Fe} \) , are represented by $$\begin{gathered} \bar L_{Si} = - 31 N_{Fe}^2 - 4 kcals per mole \hfill \\ \bar L_{Fe} = - 31 N_{Si}^2 kcals per mole \hfill \\ \end{gathered} $$ ForN _Si less than 0.35 and by $$\begin{gathered} \bar L_{Si} = - 22 N_{Fe}^2 \hfill \\ \bar L_{Fe} = - 22 N_{Fe}^2 - 7.0 \hfill \\ \end{gathered} $$ forN _Fe less than 0.35. There is an inflection point in the transition region similar to an excess stability peak for the excess free energies. At 1600°C the ThO₂(Y₂O₃) electrolyte exhibited insignificant electronic conductivity at oxygen partial pressures as low as that in equilibrium with Si?SiO₂ (2×10^?16 atm).     

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.     

Effect of Slag Composition on the Concentration of Al2O3 in the Inclusions in Si-Mn-killed Steel

The thermodynamic equilibria between CaO-Al₂O₃-SiO₂-CaF₂-MgO(-MnO) slag and Fe-1.5 mass pct Mn-0.5 mass pct Si-0.5 mass pct Cr melt was investigated at 1873 K (1600 °C) in order to understand the effect of slag composition on the concentration of Al₂O₃ in the inclusions in Si-Mn-killed steels. The composition of the inclusions were mainly equal to (mol pct MnO)/(mol pct SiO₂) = 0.8(±0.06) with Al₂O₃ content that was increased from about 10 to 40 mol pct by increasing the basicity of slag (CaO/SiO₂ ratio) from about 0.7 to 2.1. The concentration ratio of the inclusion components, \( {{X_{{{\text{Al}}_{2} {\text{O}}_{3} }} \cdot X_{\text{MnO}} } \mathord{\left/ {\vphantom {{X_{{{\text{Al}}_{2} {\text{O}}_{3} }} \cdot X_{\text{MnO}} } {X_{{{\text{SiO}}_{2} }} }}} \right. \kern-0pt} {X_{{{\text{SiO}}_{2} }} }} \) , and the activity ratio of the steel components, \( {{a_{\text{Al}}^{2} \cdot a_{\text{Mn}} \cdot a_{\text{O}}^{2} } \mathord{\left/ {\vphantom {{a_{\text{Al}}^{2} \cdot a_{\text{Mn}} \cdot a_{\text{O}}^{2} } {a_{\text{Si}} }}} \right. \kern-0pt} {a_{\text{Si}} }} \) , showed a good linear relationship on a logarithmic scale, indicating that the activity coefficient ratio of the inclusion components, \( {{\gamma_{{{\text{SiO}}_{2} }}^{i} } \mathord{\left/ {\vphantom {{\gamma_{{{\text{SiO}}_{2} }}^{i} } {\left( {\gamma_{{{\text{Al}}_{2} {\text{O}}_{3} }}^{i} \cdot \gamma_{\text{MnO}}^{i} } \right)}}} \right. \kern-0pt} {\left( {\gamma_{{{\text{Al}}_{2} {\text{O}}_{3} }}^{i} \cdot \gamma_{\text{MnO}}^{i} } \right)}} \) , was not significantly changed. From the slag-steel-inclusion multiphase equilibria, the concentration of Al₂O₃ in the inclusions was expressed as a linear function of the activity ratio of the slag components, \( {{a_{{{\text{Al}}_{2} {\text{O}}_{3} }}^{s} \cdot a_{\text{MnO}}^{s} } \mathord{\left/ {\vphantom {{a_{{{\text{Al}}_{2} {\text{O}}_{3} }}^{s} \cdot a_{\text{MnO}}^{s} } {a_{{{\text{SiO}}_{2} }}^{s} }}} \right. \kern-0pt} {a_{{{\text{SiO}}_{2} }}^{s} }} \) on a logarithmic scale. Consequently, a compositional window of the slag for obtaining inclusions with a low liquidus temperature in the Si-Mn-killed steel treated in an alumina ladle is recommended.     

Ti3Ga and αTi interdiffusion between 660° and 860°C

Diffusion experiments were conducted in vacuum with bimetallic couples of the Ti₃Ga (α₂) composition and unalloyed α titanium. The gallium-composition profiles after various timetemperature exposures were determined by microprobe analyzer transverses and evaluated by established techniques. Results from this evaluation include the definition of the α to α +α₂ and the α + α₂ to α₂ phase boundaries for the Ti?Ga system and the determination of the interdiffusion coefficients for gallium in the α Ti and Ti₃Ga (α₂) phases. The interdiffusion coefficients were found to conform to the relationships: $$\tilde D_{\alpha Ti} = 4.4 \times 10^{ - 4} \exp [ - (43.4 \pm 4.7)10^3 /RT]cm^2 /\sec $$ $$\tilde D_{\alpha Ti_3 Ga} = 7.4 \times 10^{ - 5} \exp [ - (43.8 \pm 10.7)10^3 /RT]cm^2 /\sec $$      

Crystallographic Reconstruction Study of the Effects of Finish Rolling Temperature on the Variant Selection During Bainite Transformation in C-Mn High-Strength Steels

The effect of finish rolling temperature on the austenite-(γ) to-bainite (α) phase transformation is quantitatively investigated in high-strength C-Mn steels using an alternative crystallographic γ reconstruction procedure, which can be directly applied to experimental electron backscatter diffraction mappings. In particular, the current study aims to clarify the respective contributions of the γ conditioning during the hot rolling and the variant selection during the phase transformation to the inherited texture. The results confirm that the sample finish rolled at the lowest temperature [1102 K (829 °C)] exhibits the sharpest transformation texture. It is shown that this sharp texture is exclusively due to a strong variant selection from parent brass {110} \( \left\langle {1\bar{1}2} \right\rangle \) , S {213} \( \left\langle {\bar{3}\bar{6}4} \right\rangle \) and Goss {110}〈001〉 grains, whereas the variant selection from the copper {112} \( \left\langle {\bar{1}\bar{1}1} \right\rangle \) grains is insensitive to the finish rolling temperature. In addition, a statistical variant selection analysis proves that the habit planes of the selected variants do not systematically correspond to the predicted active γ slip planes using the Taylor model. In contrast, a correlation between the Bain group to which the selected variants belong and the finish rolling temperature is clearly revealed, regardless of the parent orientation. These results are discussed in terms of polygranular accommodation mechanisms, especially in view of the observed development in the hot-rolled samples of high-angle grain boundaries with misorientation axes between 〈111〉γ and 〈110〉γ.     

The Solubility and Diffusivity of Fluorine in Solid Copper from Electrochemical Measurements

The solubility and diffusivity of fluorine in solid copper were determined electrochemically using the double solid-state cell $$Ni + NiF_2 \left| {CaF_2 } \right|Cu\left| {CaF_2 } \right|Ni + NiF_2 .$$ In the temperature range 757 to 920°C, the diffusivity of fluorine in solid copper was found to be $$D_F \left( {{{cm^2 } \mathord{\left/ {\vphantom {{cm^2 } s}} \right. \kern-\nulldelimiterspace} s}} \right) = 9.32 \times 10^{ - 2} \exp \left( {\frac{{ - 98,910 {J \mathord{\left/ {\vphantom {J {mole}}} \right. \kern-\nulldelimiterspace} {mole}}}} {{RT}}} \right).$$ . The results obtained for the dissolution of fluorine as atoms in solid copper showed large scatter. However, the equilibrium dissolution of fluorine follows Sieverts’ law. Above the melting point (770°C) of CuF₂, the mean solubility of fluorine in solid copper, for the equilibrium Cu(s)+ CuF ₂(l), follows the relationship $$N_F^s (atom fraction) = 0.98 \exp \left( {\frac{{ - 79,500 {J \mathord{\left/ {\vphantom {J {mole}}} \right. \kern-\nulldelimiterspace} {mole}}}} {{RT}}} \right).$$      

Mass-spectrometric determination of activities in Fe−Cr and Fe−Cr−Ni alloys

The Knudsen cell-mass spectrometer combination has been used to study the Fe?Cr system and some Fe?Cr?Ni liquid alloys. The Fe?Cr liquid alloys at 1600°C are found to be essentially ideal when referred to pure liquids as standard states. Phase equilibria over a limited composition range for this system are derived from the behavior of the ion-current ratios. The necessary equations are derived to apply the integration technique to the measured ion current ratios in a ternary system and the method is applied to the Fe?Cr?Ni system at 1600°C. The results are represented, within experimental error, by the following equations: forN _Fe≥0.6, $$\begin{gathered} ln \gamma _{Fe} = - 0.08 N_{Ni}^2 \hfill \\ \ln \gamma _{Cr} = 0.09 - 0.08 N_{Ni}^2 \hfill \\ \ln \gamma _{Ni} = - 0.26 - 0.08(1 - N_{Ni} )^2 \hfill \\ \end{gathered} $$ forN _Fe=0.45, $$\begin{gathered} \ln \gamma _{Fe} = - 0.20 N_{Ni}^2 \hfill \\ \ln \gamma _{Cr} = 0.09 - 0.20 N_{Ni}^2 \hfill \\ \ln \gamma _{Ni} = - 0.19 - 0.20(1 - N_{Ni} )^2 \hfill \\ \end{gathered} $$      

Thermodynamic properties of the Cu−Sn and Cu−Au systems by mass spectrometry

The activities and partial molar heats of mixing have been determined in the liquid Cu?Sn system at 1320°C and the liquid Cu?Au system at 1460°C. The experimental technique consisted of the analysis of Knudsen cell effusates with a T.O.F. mass spectrometer. The ion current ratio for the alloy components was measured for each system over a range of temperature and composition and the thermodynamic values calculated by a modified Gibbs-Duhem equation. Both systems exhibited negative deviations from ideal behavior. The results can be partially represented by the equations $$\begin{gathered} \log \gamma _{Cu} = - 0.0175x^2 _{Sn} - 0.302 (0 \leqslant x_{Cu} \leqslant 0.20) \hfill \\ log \gamma _{Sn} = - 0.342x^2 _{Cu} + 1.084(0 \leqslant x_{Sn} \leqslant 0.20) \hfill \\ \end{gathered} $$ for the Cu?Sn system at 1320°C and by $$\begin{gathered} \log \gamma _{Cu} = - 0.703x^2 _{Au} - 0.083(0 \leqslant x_{Cu} \leqslant 0.52) \hfill \\ \log \gamma _{Au} = - 1.057x^2 _{Cu} + 0.098(0 \leqslant x_{Au} \leqslant 0.47) \hfill \\ \end{gathered} $$ for the Cu?Au system at 1460°C.     

An investigation of the thermodynamics and kinetics of the ferric chloride brine leaching of galena concentrate

The solution thermodynamics of acidified ferric chloride brine lixiviants and the dissolution kinetics of a galena concentrate in such solutions have been investigated. The distribution of the various metal chloro complexes calculated from available thermodynamic data shows that the distribution is shifted to the higher complexes, predominantly FeCl ₃ ^o , FeCl ₂ ^o , and PbCl ₄ ⁼ , as the total Cl^? concentration increases, and that the distribution is unaffected by the extent of reaction. The dissolution of PbS concentrate is presented as a competition between a nonoxidative reaction with H⁺ and the oxidative reaction with ferric ion. Acid dissolution of PbS predominates when the activity ratio of hydrogen ion to ferric ion is high. Under these conditions H₂S is produced. When the activity ratio of hydrogen ion to ferric ion is low, and especially when the concentration of Fe³⁺ is greater than 0.15 M, oxidative dissolution of PbS becomes the controlling reaction. The dissolution can be represented by a shrinking core model with a surface chemical reaction as the rate controlling step. This is supported by the activation energy of 72.1 kJ/mole and the dependence of the rate on the inverse of the particle radius. The following rate equation was found to be in excellent agreement with the experimentally observed leaching behavior for 0.15 to 0.6 M [Fe⁺³] _T up to approximately 90 to 95 pet extraction: $$1 - \left( {1 - \alpha } \right)^{1/3} = \left[ {\frac{{2.3 x 10^{12} }}{{r_0 }}\left[ {{\text{Fe}}^{{\text{ + 3}}} } \right]_T^{0.21} \exp \left( {\frac{{ - 72100}}{{{\text{R}}T}}} \right)} \right]t$$ The rate deviates from the 0.21 order for Fe⁺³ concentrations greater than 0.6 M. The deviation from the surface model at higher values of PbS conversion is due to the presence of solid PbCl₂ in the pores of the reacting particles.     

Nonmetallic inclusions in 13XΦA steel for reliable oil-line pipe

Nonmetallic inclusions in low-alloy 13XΦA steel mass-produced at OAO Severskii Trubnyi Zavod are studied. Corrosive nonmetallic inclusions of two types, identified by etching, are found to consist of two phases: MgO · Al₂O₃; and CaS with some quantity of Mn. The orientations identified are \(\{ 111\} _{CaS} \left\| {\{ 110\} _{MgO \cdot Al_2 O_3 } } \right.\) and \(\left\langle {1\bar 10} \right\rangle _{CaS} \left\| {\left\langle {1\bar 11} \right\rangle _{MgO \cdot Al_2 O_3 } } \right.\) .     

Nitrogen diffusion and solubility in tungsten

The diffusion and solubility of nitrogen in tungsten were determed using an ultrahigh vacuum-and mass-spectrometric technique capable of measuring concentrations of 10^?2 ppm and degassing rates of 10^?3 ppm N per hr. The technique is based on measuring the degassing rate of nitrogen as a function of time from a resistivity heated tungsten wire previously engassed with nitrogen between 1 and 25 torr. The diffusion and solubility constants between 1000° and 1800°C may be summarized by $$D = (2.37 \pm 0.43) \times 10^{ - 3} \exp [( - 35,800 \pm 3900)/RT] cm^2 /\sec ,$$ , and $$S = (0.21 \pm 0.06) \exp [( - 17,600 \pm 5900)/RT] torr \cdot liter cm^{ - 3} torr^{ - 1/2} .$$ . The concentration of nitrogen in tungsten at 760 torr according to these results are 0.4 and 9.2 ppm at 1000° and 2000°C, respectively. The expression for the permeation constants calculated fromD andS is $$K = 5 \times 10^{ - 4} \exp ( - 53,400/RT) torr \cdot liter cm^{ - 1} sec^{ - 1} torr^{ - 1/2} .$$ .     

Plastic deformation of titanium at elevated temperatures

The deformation of iodide titanium single crystals containing 200 to 250 ppm O, was studied in compression at temperatures from 25° to 800°C. Reduction of about 5 pct along thec axis was accommodated almost entirely by \(\left\{ {11\bar 22} \right\}\) twinning from 25° to 300°C, and above 400°C by \(\left\{ {10\bar 11} \right\}\) twinning in combination with c+a slip. The stress for \(\left\{ {11\bar 22} \right\}\) twinning increased with increasing temperature, and twin formation was accompanied by a load drop, while the stress for \(\left\{ {10\bar 11} \right\}\) twinning decreased with increasing temperature and twinning was not accompanied by a load drop. Crystals reduced normal to thec axis deformed by a combination of prism slip and \(\left\{ {10\bar 12} \right\}\) twinning at 25°C and by prism slip alone above 500°C.     

Oxygen pressure dependence of tracer diffusivities of Ca and Fe in liquid CaO-SiO2-FexO system

The tracer diffusivities of calcium and iron in a steel-making slag of 33 pct CaO-27 pct SiO₂-40 pct Fe₂O₃ by charge composition have been measured at 1360 to 1460°C as a function of temperature and oxygen pressure in the gas phase. The results expressed in cm²/s (in SI unit of m²/s, the following equation should be divided by 10,000) are given by $$D^{tr} = D_0 \left[ {P_{O_2 } } \right]^{1/\chi } \exp \left[ { - \frac{E}{{RT}}} \right](at 1360 to 1460^\circ C)$$ where for tracer diffusion of iron, D_o is 0.2,x is 8.5, andE is 26 kcal/mol (1.09 x 10⁴ J/ mol) and for tracer diffusion of calcium, D_o is 0.1,x is 12.5, andE is 28 kcal/mol (1.17 × 10⁴ J/mol). Prior to diffusion runs, the slag was equilibrated with the gas mixture of carbon monoxide and dioxide with an oxygen pressure of 10^?11 to 10^?8 atm. The diffusivity was measured by the instantaneous plane source method, using radioactive tracers of calcium and iron. The increase of the tracer diffusivities with the oxygen pressure was interpreted in relation to a probable increase of the divalent cation vacancies in the slag.