Development of Low-Fluoride Slag for Electroslag Remelting: Role of Li<Subscript>2</Subscript>O on the Viscosity and Structure of the Slag

The effect of the substitution of CaF₂ with Li₂O on the viscosity and structure of low-fluoride CaF₂-CaO-Al₂O₃-MgO slag was studied with an aim to develop low-fluoride slag for electroslag remelting. Increasing Li₂O addition up to 4.5 mass pct was observed to significantly reduce the slag viscosity monotonically. Increasing temperature significantly lowered the viscosity of slag, whereas this influence is less effective with increasing Li₂O content especially above 3.5 mass pct. The activation energy for viscous flow decreases with increasing Li₂O content. The polymerization degree of aluminate networks decreased with increasing Li₂O content, as demonstrated by Raman analysis. The dominant structural unit in [AlO₄]^5?-tetrahedral network is \( {\text{Q}}_{\text{Al}}^{4} \). The amount of symmetric Al-O⁰ stretching vibrations significantly decreased with increasing Li₂O content. The relative fraction of \( {\text{Q}}_{\text{Al}}^{4} \) in the [AlO₄]^5?-tetrahedral units shows a decreasing trend, whereas \( {\text{Q}}_{\text{Al}}^{2} \) increases with the increase in Li₂O content accordingly. The change in slag viscosity with chemistry variation agrees well with the changes in slag structural units.     

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.     

Effects of Interfacial Lattice Mismatching on Wetting of Ni-Plated Steel by Magnesium

In this study, wetting has been characterized by measuring the contact angles of AZ92 Mg alloy on Ni-electroplated steel as a function of temperature. Reactions between molten Mg and Ni led to a contact angle of about 86 deg in the temperature range of 891 K to 1023 K (618 °C to 750 °C) (denoted as Mode I) and a dramatic decrease to about 46 deg in the temperature range of 1097 K to 1293 K (824 °C to 1020 °C) (denoted as Mode II). Scanning and transmission electron microscopy (SEM and TEM) indicated that AlNi + Mg₂Ni reaction products were produced between Mg and steel (Mg-AlNi-Mg₂Ni-Ni-Fe) in Mode I, and just AlNi between Mg and steel (Mg-AlNi-Fe) in Mode II. From high resolution TEM analysis, the measured interplanar mismatches for different formed interfaces in Modes I and II were \( 17{\kern 1pt} \;{\text{pct}}_{{\{ 10\overline 11\}_{\text{Mg}} //\{ 110\}_{\text{AlNi}} }} \)-\( 104.3\;{\text{pct}}_{{\{ 110\}_{\text{AlNi}} //\left\{ {10\overline{1}0} \right\}_{{{\text{Mg}}_{ 2} {\text{Ni}}}} }} \)-\( 114\,{\text{pct}}_{{\left\{ {0003} \right\}_{{{\text{Mg}}_{ 2} {\text{Ni}}}} //\{ 111\}_{\text{Ni}} }} \) and \( 18\,{\text{pct}}_{{\{ 10\overline 11\}_{\text{Mg}} //\{ 110\}_{\text{AlNi}} }} \)-\( 5\,{\text{pct}}_{{\left\{ {110} \right\}_{\text{AlNi}} //\{ 110\}_{\text{Fe}} }} \), respectively. An edge-to-edge crystallographic model analysis confirmed that Mg₂Ni produced larger lattice mismatching between interfaces with calculated minimum interplanar mismatches of \( 16.4\,{\text{pct}}_{{{\text{\{ 10}}\overline 1 1 {\text{\} }}_{\text{Mg}} / / {\text{\{ 110\} }}_{\text{AlNi}} }} \)-\( 108.3\,{\text{pct}}_{{{\text{\{ 110\} }}_{\text{AlNi}} / / {\text{\{ 10}}\overline 1 1 {\text{\} }}_{{{\text{Mg}}_{ 2} {\text{Ni}}}} }} \)-\( 17.2\,{\text{pct}}_{{{\text{\{ 10}}\overline 1 1 {\text{\} }}_{{{\text{Mg}}_{ 2} {\text{Ni}}}} / / {\text{\{ 100\} }}_{\text{Ni}} }} \) for Mode I and \( 16.4\,{\text{pct}}_{{{\text{\{ 10}}\overline1 1 {\text{\} }}_{\text{Mg}} / / {\text{\{ 110\} }}_{\text{AlNi}} }} \)-\( 0.6\,{\text{pct}}_{{{\text{\{ 111\} }}_{\text{AlNi}} / / {\text{\{ 111\} }}_{\text{Fe}} }} \) for Mode II. Therefore, it is suggested that the poor wettability in Mode I was caused by the existence of Mg₂Ni since AlNi was the immediate layer contacting molten Mg in both Modes I and II, and the presence of Mg₂Ni increases the interfacial strain energy of the system. This study has clearly demonstrated that the lattice mismatching at the interfaces between reaction product(s) and substrate, which are not in direct contact with the liquid, can greatly influence the wetting of the liquid.     

The Thermodynamic Behavior of Copper in the CaO-B<Subscript>2</Subscript>O<Subscript>3</Subscript> and BaO-B<Subscript>2</Subscript>O<Subscript>3</Subscript> Slag System

The Cu solubility was measured in the CaO-B₂O₃ and BaO-B₂O₃ slag systems to understand the dissolution mechanism of Cu in the slags. The Cu solubility had a linear relationship with oxygen partial pressure in the CaO-B₂O₃ slag system, which corresponds with previous studies. Also, the Cu solubilities in slag decreased with increasing the slag basicity, which value of slope was close to –0.5 in logarithmic form. From the results of experiment, the Cu dissolution mechanism established as follows:

Density Measurements of Low Silica CaO-SiO<Subscript>2</Subscript>-Al<Subscript>2</Subscript>O<Subscript>3</Subscript> Slags

Density measurements of a low-silica CaO-SiO₂-Al₂O₃ system were carried out using the Archimedes principle. A Pt 30 pct Rh bob and wire arrangement was used for this purpose. The results obtained were in good agreement with those obtained from the model developed in the current group as well as with other results reported earlier. The density for the CaO-SiO₂ and the CaO-Al₂O₃ binary slag systems also was estimated from the ternary values. The extrapolation of density values for high-silica systems also showed good agreement with previous works. An estimation for the density value of CaO was made from the current experimental data. The density decrease at high temperatures was interpreted based on the silicate structure. As the mole percent of SiO₂ was below the 33 pct required for the orthosilicate composition, discrete \textSiO₄^{4 -} {\text{SiO}}_{4}^{4 - } tetrahedral units in the silicate melt would exist along with O^2– ions. The change in melt expansivity may be attributed to the ionic expansions in the order of

On the Prediction of <Emphasis Type="Italic">α</Emphasis>-Martensite Temperatures in Medium Manganese Steels

A new composition-based method for calculating the α-martensite start temperature in medium manganese steel is presented and uses a regular solution model to accurately calculate the chemical driving force for α-martensite formation, \( \Delta G_{\text{Chem}}^{\gamma \to \alpha } \). In addition, a compositional relationship for the strain energy contribution during martensitic transformation was developed using measured Young’s moduli (E) reported in literature and measured values for steels produced during this investigation. An empirical relationship was developed to calculate Young’s modulus using alloy composition and was used where dilatometry literature did not report Young’s moduli. A comparison of the \( \Delta G_{\text{Chem}}^{\gamma \to \alpha } \) normalized by dividing by the product of Young’s modulus, unconstrained lattice misfit squared (δ ²), and molar volume (Ω) with respect to the measured α-martensite start temperatures, \( M_{\text{S}}^{\alpha } \), produced a single linear relationship for 42 alloys exhibiting either lath or plate martensite. A temperature-dependent strain energy term was then formulated as \( \Delta G_{\text{str}}^{\gamma \to \alpha } \left( {{\text{J}}/{\text{mol}}} \right) = E\varOmega \delta^{2} (14.8 - 0.013T) \), which opposed the chemical driving force for α-martensite formation. \( M_{\text{S}}^{\alpha } \) was determined at a temperature where \( \Delta G_{\text{Chem}}^{\gamma \to \alpha } + \Delta G_{\text{str}}^{\gamma \to \alpha } = 0 \). The proposed \( M_{\text{S}}^{\alpha } \) model shows an extended temperature range of prediction from 170 K to 820 K (?103 °C to 547 °C). The model is then shown to corroborate alloy chemistries that exhibit two-stage athermal martensitic transformations and two-stage TRIP behavior in three previously reported medium manganese steels. In addition, the model can be used to predict the retained γ-austenite in twelve alloys, containing ε-martensite, using the difference between the calculated \( M_{\text{S}}^{\varepsilon } \) and \( M_{\text{S}}^{\alpha } \).     

Oxidation and Condensation of Zinc Fume From Zn-CO<Subscript>2</Subscript>-CO-H<Subscript>2</Subscript>O Streams Relevant to Steelmaking Off-Gas Systems

The objective of this research was to study the condensation of zinc vapor to metallic zinc and zinc oxide solid under varying environments to investigate the feasibility of in-process separation of zinc from steelmaking off-gas dusts. Water vapor content, temperature, degree of cooling, gas composition, and initial zinc partial pressure were varied to simulate the possible conditions that can occur within steelmaking off-gas systems, limited to Zn-CO₂-CO-H₂O gas compositions. The temperature of deposition and the effect of rapidly quenching the gas were specifically studied. A homogeneous nucleation model for applicable experiments was applied to the analysis of the experimental data. It was determined that under the experimental conditions, oxidation of zinc vapor by H₂O or CO₂ does not occur above 1108 K (835 °C) even for highly oxidizing streams (CO₂/CO = 40/7). Rate expressions that correlate CO₂ and H₂O oxidation rates to gas composition, partial pressure of water vapor, temperature, and zinc partial pressure were determined to be as follows:

$$ {\text{Rate}}\left( {\frac{\text{mol}}{{{\text{m}}^{2} {\text{s}}}}} \right) = 406 \exp \left( {\frac{{ - 50.2 \,{\text{kJ}}/{\text{mol}}}}{RT}} \right)\left( {p_{\text{Zn}} p_{{{\text{CO}}_{2} }} - p_{\text{CO}} /K_{{{\text{eq}},{\text{CO}}_{2} }} } \right)\,\frac{\text{mol}}{{{\text{m}}^{2} \times {\text{s}}}} $$

$$ {\text{Rate}}\left( {\frac{\text{mol}}{{{\text{m}}^{2} {\text{s}}}}} \right) = 32.9 \exp \left( {\frac{{ - 13.7\, {\text{kJ}}/{\text{mol}}}}{RT}} \right)\left( {p_{\text{Zn}} p_{{{\text{H}}_{2} {\text{O}}}} - p_{{{\text{H}}_{2} }} /K_{{{\text{eq}},{\text{H}}_{2} {\text{O}}}} } \right)\,\frac{\text{mol}}{{{\text{m}}^{2} \times {\text{s}}}} $$

It was proven that a rapid cooling rate (500 K/s) significantly increases the ratio of metallic zinc to zinc oxide as opposed to a slow cooling rate (250 K/s). SEM analysis found evidence of heterogeneous growth of ZnO as well as of homogeneous formation of metallic zinc. The homogeneous nucleation model fit well with experiments where only metallic zinc deposited. An expanded model with rates of oxidation by CO₂ and H₂O as shown was combined with the homogenous nucleation model and then compared with experimental data. The calculated results based on the model gave a reasonable fit to the measured data. For the conditions used in this study, the rate equations for the oxidation of zinc by carbon dioxide and water vapor as well as the homogeneous nucleation model of metallic zinc were applicable for various temperatures, zinc partial pressures, CO₂:CO ratios, and H₂O partial pressures.

     

Effect of Tungsten on Primary Creep Deformation and Minimum Creep Rate of Reduced Activation Ferritic-Martensitic Steel

Effect of tungsten on transient creep deformation and minimum creep rate of reduced activation ferritic-martensitic (RAFM) steel has been assessed. Tungsten content in the 9Cr-RAFM steel has been varied between 1 and 2 wt pct, and creep tests were carried out over the stress range of 180 and 260 MPa at 823 K (550 °C). The tempered martensitic steel exhibited primary creep followed by tertiary stage of creep deformation with a minimum in creep deformation rate. The primary creep behavior has been assessed based on the Garofalo relationship, \( \varepsilon = \varepsilon_{\text{o}} + \varepsilon_{\text{T}} [1-\exp (-r^{\prime} \cdot t)] + \dot{\varepsilon }_{\text{m}} \cdot t \) , considering minimum creep rate \( \dot{\varepsilon }_{\text{m}} \) instead of steady-state creep rate \( \dot{\varepsilon }_{\text{s}} \) . The relationships between (i) rate of exhaustion of transient creep r′ with minimum creep rate, (ii) rate of exhaustion of transient creep r′ with time to reach minimum creep rate, and (iii) initial creep rate \( \dot{\varepsilon }_{\text{i}} \) with minimum creep rate revealed that the first-order reaction-rate theory has prevailed throughout the transient region of the RAFM steel having different tungsten contents. The rate of exhaustion of transient creep r′ and minimum creep rate \( \dot{\varepsilon }_{\text{m}} \) decreased, whereas the transient strain ? _T increased with increase in tungsten content. A master transient creep curve of the steels has been developed considering the variation of \( \frac{{\left( {\varepsilon - \varepsilon_{\text{o}} } \right)}}{{\varepsilon_{\text{T}} }} \) with \( \frac{{\dot{\varepsilon }_{\text{m}} \cdot t}}{{\varepsilon_{\text{T}} }} \) . The effect of tungsten on the variation of minimum creep rate with applied stress has been rationalized by invoking the back-stress concept.     

Sulfate Formation and Decomposition of Nickel Concentrates

Nickel sulfide concentrates from two Canadian nickel concentrators were investigated to improve the understanding of SO₂ formation and release during processing. The concentrates were heated in gases of various oxygen concentrations up to 1573 K (1300 °C) in a thermal gravimetric analysis unit to simulate what may take place during calcine collection and processing. The resulting SO₂ gases were also measured. It was determined that during oxidation, there are competing reactions, such as \( 3{\text{FeS}} + 5{\text{O}}_{2} = {\text{Fe}}_{3} {\text{O}}_{4} + 3{\text{SO}}_{2} \) leading to mass loss, or \( 2{\text{FeS}} + 5{\text{O}}_{2} + {\text{SO}}_{2} = {\text{Fe}}_{2} \left( {{\text{SO}}_{4} } \right)_{3} \) causing mass gain. At temperatures up to approximately 973 K (700 °C), sulfates were formed readily, whereas at higher temperatures, they would decompose, evolving SO₂. By lowering the oxygen content in the surrounding gas, the sulfates decomposed more readily. In an argon or hydrogen atmosphere or in vacuum, it is possible to enhance the sulfate decomposition greatly, possibly allowing for reduced SO₂ emissions from the electric furnaces.     

Isothermal Reaction of NiO Powder with Undiluted CH<Subscript>4</Subscript> at 1000 K to 1300 K (727 °C to 1027 °C)

In this study, isothermal reaction behavior of loose NiO powder in a flowing undiluted CH₄ atmosphere at the temperature range 1000 K to 1300 K (727 °C to 1027 °C) is investigated. Thermodynamic analyses at this temperature range revealed that single phase Ni forms at the input \( {{n_{{{\text{CH}}_{ 4} }}^{\text{o}} } \mathord{\left/ {\vphantom {{n_{{{\text{CH}}_{ 4} }}^{\text{o}} } {\left( {n_{{{\text{CH}}_{ 4} }}^{\text{o}} + n_{\text{NiO}}^{\text{o}} } \right)}}} \right. \kern-0pt} {\left( {n_{{{\text{CH}}_{ 4} }}^{\text{o}} + n_{\text{NiO}}^{\text{o}} } \right)}} \) mole fractions (\( X_{{{\text{CH}}_{ 4} }} \)) between ~0.2 and 0.5. It was also predicted that free C co-exists with Ni at \( X_{{{\text{CH}}_{ 4} }} \) values higher than ~0.5. The experiments were carried out as a function of temperature, time, and CH₄ flow rate. Mass measurement, XRD and SEM-EDX were used to characterize the products at various stages of the reaction. At 1200 K and 1300 K (927 °C and 1027 °C), the reaction of NiO with undiluted CH₄ essentially consisted of two successive distinct stages: NiO reduction and pyrolytic C deposition on pre-reduced Ni particles. At 1200 K (927 °C), 1100 K (827 °C), and 1000 K (727 °C), complete oxide reduction was observed within ~7.5, ~17.5, and ~45 minutes, respectively. It was suggested that NiO was essentially reduced to Ni by a CH₄ decomposition product, H₂. Possible reactions leading to NiO reduction were suggested. An attempt was made to describe the NiO reduction kinetics using nucleation-growth and geometrical contraction models. It was observed that the extent of NiO reduction and free C deposition increased with the square root of CH₄ flow rate as predicted by a mass transport theory. A mixed controlling mechanism, partly chemical kinetics and partly external gaseous mass transfer, was responsible for the overall reaction rate. The present study demonstrated that the extent of the reduction can be determined quantitatively using the XRD patterns and also using a formula theoretically derived from the basic XRD data.     

Diffusion of Au in the Intermetallic Compound Ti<Subscript>3</Subscript>Al

The Au diffusion in the Ti₃Al compound was investigated at six compositions from 25 to 35 at. pct Al by using the diffusion couples (Ti-X at. pct Al/Ti-X at. pct Al-2 at. pct Au; X = 25, 27, 29, 31, 32, and 35) at 1273 to 1423 K. The diffusion coefficients of Au in Ti₃Al ( D_\textAu^{\textTi₃ \textAl} ) \left( {D_{\text{Au}}^{{{\text{Ti}}_{3} {\text{Al}}}} } \right) are relatively close to those of Ti. The D_\textAu^{\textTi₃ \textAl} \texts {D}_{\text{Au}}^{{{\text{Ti}}_{3} {\text{Al}}}} {\text{s}} slightly increase with Al concentration within the same order of magnitude. The activation energies of Au diffusion, Q_\textAu^{\textTi₃ \textAl} \texts, Q_{\text{Au}}^{{{\text{Ti}}_{3} {\text{Al}}}} {\text{s}}, evaluated from the Arrhenius plots were relatively close to those of Ti diffusion, Q_\textTi^{\textTi₃ \textAl} \texts, Q_{\text{Ti}}^{{{\text{Ti}}_{3} {\text{Al}}}} {\text{s}}, rather than those of Al diffusion, Q_\textAl^{\textTi₃ \textAl} \texts; {Q}_{\text{Al}}^{{{\text{Ti}}_{3} {\text{Al}}}} {\text{s}}; therefore, it was suggested that Au atoms diffuse by the sublattice diffusion mechanism in which Au atoms substitute for Ti sites preferentially in Ti₃Al and diffuse by vacancy mechanism on Ti sublattice. The influence of the D0₁₉ ordered structure (hcp base) of Ti₃Al on diffusion of Au and other elements is discussed by comparing the diffusivities in Ti₃Al and α-Ti.     

Enrichment Mechanism of Phosphate in CaO-SiO2-FeO-Fe2O3-P2O5 Steelmaking Slags

The phosphate-enrichment behavior has experimentally been investigated in CaO-SiO₂-FeO-Fe₂O₃-P₂O₅ steelmaking slags. The reaction ability of structural units in the slags has been represented the mass action concentration \( N_{i} \) from the developed ion and molecule coexistence theory (IMCT)- \( N_{i} \) model based on the IMCT. The defined enrichment possibility \( N_{{{\text{c}}i{\text{ {-}c}}j}} \) and enrichment degree \( R_{{{\text{c}}i{\text{{-}c}}j}} \) of solid solutions containing P₂O₅ from the developed IMCT- \( N_{i} \) model have been verified from the experimental results. The effects of binary basicity, the mass percentage ratio \( {{ ( {\text{pct Fe}}_{t} {\text{O)}}} \mathord{\left/ {\vphantom {{ ( {\text{pct Fe}}_{t} {\text{O)}}} { ( {\text{pct CaO)}}}}} \right. \kern-0pt} { ( {\text{pct CaO)}}}} \) , and mass percentage of P₂O₅ in the initial slags on phosphate-enrichment behavior in the slags has also been discussed. The results show that the P₂O₅ component can easily be bonded by CaO to form tricalcium phosphate 3 CaO·P₂O₅, and the formed 3CaO·P₂O₅ can react with the produced dicalcium silicate 2CaO·SiO₂ to generate solid-solution 2CaO·SiO₂-3CaO·P₂O₅ under fixed cooling conditions. The maximum value of the defined enrichment degree \( R_{{{\text{C}}_{ 2} {\text{S{-}}} {\text{C}}_{ 3} {\text{P}}}} \) of solid-solution 2CaO·SiO₂-3CaO·P₂O₅ is obtained as 0.844 under conditions of binary basicity as 2.5 and the mass percentage ratio \( {{ ( {\text{pct Fe}}_{t} {\text{O)}}} \mathord{\left/ {\vphantom {{ ( {\text{pct Fe}}_{t} {\text{O)}}} { ( {\text{pct CaO)}}}}} \right. \kern-0pt} { ( {\text{pct CaO)}}}} \) as 0.955 at fixed cooling conditions.     

Sulfide capacities of CaO-CaF2-CaCl2 melts

The sulfide capacity of CaO-CaF₂-CaCl₂ slag was determined at temperatures from 1000 °C to 1300 °C by equilibrating molten slag, molten silver, and CO-CO₂-Ar gas mixture. The sulfide capacity increases with replacing CaCl₂ by CaF₂ in slags of constant CaO contents. The sulfide capacity also increases with increasing temperature as well as with increasing CaO content at a constant ratio of CaF₂/CaCl₂ of unity. A linear relationship between the sulfide capacity and carbonate capacity in literature was observed on a logarithmic scale. SIMEON SIMEONOV, formerly Visiting Research Fellow, Institute of Industrial Science, University of Tokyo. TOSHIHIKO SAKAI, formerly Research Fellow, Institute of Industrial Science, University of Tokyo.     

Mass spectrometric determination of activities for alloys with complex vapor species: Bi-Pb and Bi-Tl

For solutions from which complex species vaporize (Bi₂, Si₂, Al₂O, Sb₄, and so forth) new methods of determining the thermodynamic properties from mass spectrometric data are demonstrated. In order to test the feasibility of these new techniques, experiments have been carried out on the liquid Bi-Pb and Bi-Tl systems for which adequate thermodynamic data are available. In evaluating the thermodynamic properties, the ion current ratiosI _Pb ⁺/IBi₂/+ andI _Tl ⁺/IBi₂/+ were employed,e.g. $$\log {\text{ }}\gamma _{{\text{Bi}}} {\text{ = - }}\mathop {\int {\frac{{N_{Pb} }}{{1{\text{ + }}N_{Pb} }}d} }\limits_{N_{Bi} = 1}^{N_{{\text{Bi}}} = N_{Bi} } {\text{ }}\left\{ {{\text{log}}\frac{{{\text{1}}_{{\text{Pb}}}^{\text{ + }} {\text{ }}N_{Bi}^2 }}{{I_{Bi2}^ + {\text{ }}N_{Pb} }}} \right\}$$ Measuring these particular ion current ratios eliminates errors resulting from the fragmentation of the complex vapor species in evaluating the thermodynamic properties. A dimer-monomer technique, which corrects for fragmentation, was also demonstrated. The results using these two independent approaches are in good agreement with each other as well as with previous investigations. The activity coefficients in both systems adhere to the quadratic formalism over large composition ranges,e.g. $$\begin{gathered} \log {\text{ }}\gamma _{{\text{Pb}}} {\text{ = - 0}}{\text{.255 }}N_{Bi}^2 {\text{ }}N_{{\text{Bi}}} {\text{< 0}}{\text{.8}} \hfill \\ \log {\text{ }}\gamma _{{\text{Tl}}} {\text{ = - 0}}{\text{.805 }}N_{Bi}^2 {\text{ }}N_{{\text{Bi}}} {\text{< 0}}{\text{.7}} \hfill \\ \end{gathered} $$      

Determination of Gibbs Energy of Formation of Molybdenum-Boron Binary System by Electromotive Force Measurement Using Solid Electrolyte

The standard Gibbs energies of formation of Mo₂B, ??MoB, Mo₂B₅, and MoB₄ in the molybdenum-boron binary system were determined by measuring electromotive forces of galvanic cells using an Y₂O₃-stabilized ZrO₂ solid oxide electrolyte. The results are as follows: $$ \begin{aligned} \Updelta_{\text{f}} {\text{G}}^\circ \left( {{\text{Mo}}_{2} {\text{B}}} \right)/{\text{J}}\,{\text{mol}}^{ - 1} & = - 193100 + 44.10T \pm 700\left( {1198{\text{ K to }}1323{\text{ K}}\left( {925^\circ {\text{C to }}1050^\circ {\text{C}}} \right)} \right) \\ \Updelta_{\text{f}} {\text{G}}^\circ (\alpha {\text{MoB}})/{\text{J}}\,{\text{mol}}^{ - 1} & = - 164000 + 26.45T \pm 700\left( {1213{\text{ K to }}1328{\text{ K}}\left( {940^\circ {\text{C to }}1055^\circ {\text{C}}} \right)} \right) \\ \Updelta_{\text{f}} {\text{G}}^\circ \left( {{\text{Mo}}_{2} {\text{B}}_{5} } \right)/{\text{J}}\,{\text{mol}}^{ - 1} & = - 622500 + 117.0T \pm 3000\left( {1205{\text{ K to }}1294{\text{ K}}\left( {932^\circ {\text{C to }}1021^\circ {\text{C}}} \right)} \right) \\ \Updelta_{\text{f}} {\text{G}}^\circ \left( {{\text{MoB}}_{4} } \right)/{\text{J}}\,{\text{mol}}^{ - 1} & = - 387300 + 93.53T \pm 3000\left( {959{\text{ K to }}1153{\text{ K}}\left( {686^\circ {\text{C to }}880^\circ {\text{C}}} \right)} \right) \\ \end{aligned} $$ where the standard pressure is 1 bar (100 kPa).     

Thermal Conductivity Measurements of Some Synthetic Al2O3-CaO-SiO2 Slags by Means of a Front-Heating and Front-Detection Laser-Flash Method

The thermal conductivities of some synthetic slags containing Al₂O₃, CaO, and SiO₂ have been determined in the temperature range between 1623?K (1350?°C) and 1823?K (1550?°C) by applying a front-heating front-detection laser-flash method. In this method, the temperature response curve is measured in the short initial time period immediately after irradiating a laser pulse. The resultant values obtained by this method are unaffected by the radiative heat transfer. The temperature dependence of the thermal conductivity values were found not to be significant for all slag samples currently investigated. The thermal conductivity (??) of samples is represented with standard deviation less than 2?pct of its value as a function of compositions as follows: $$ \lambda = - 0.48\left[ {{\text{Al}}_{ 2} {\text{O}}_{ 3} } \right] - 0.57\left[ {\text{CaO}} \right] - 0.55\left[ {{\text{SiO}}_{2} } \right] + 57.1{\text{ W m}}^{ - 1} {\text{ K}}^{ - 1} $$ where [Al₂O₃], [CaO], and [SiO₂] are molar percent of Al₂O₃, CaO and SiO₂, respectively. The equation is suggested to cover the region of the following compositions: 8.0?mol pct < Al₂O₃ <21.0?mol pct, 31.5?mol pct < CaO <41.5?mol pct and 43.0?mol pct < SiO₂ <58.1?mol pct. The addition of Al₂O₃ to slags with a constant CaO/SiO₂ molar ratio resulted in an increase in thermal conductivity. In contrast, the effects of addition of SiO₂ and CaO are found to be insignificant.