Temperature Dependence of Diffuse Scattering in PZN

Structural disorder seems to relate to the useful physical properties of ferroelectrics and relaxors. One such material is PZN, PbZn_1/3Nb_2/3O₃.\hbox{PbZn}_{1/3}\hbox{Nb}_{2/3}\hbox{O}_{3}. To explore what aspects of the disorder are specific to the polarized state, the temperature dependence of diffuse scattering in PZN has been investigated. The data were collected using both neutron and X-ray single crystal experiments in a range of temperatures from 50 K to 500 K (−223 °C to 227 °C). It has been found that some features, like the diffuse scattering from the B-site ordering, remain unchanged with change of temperature in terms of both intensity and peak shape. However, other diffuse scattering features evolve with T, for example the size effect scattering around the Bragg peaks. The size-effect becomes less pronounced with increasing temperature, with the diffuse scattering becoming more symmetric around the Bragg peaks. The diffuse rods caused by the planar domains change only slightly with temperature. This finding indicates that the planar domains persist into the paraelectric state but that the correlation between lead displacement and the average separation of adjacent lead atoms becomes weaker, suggesting that this size effect may be crucial to the ferroelectric properties.     

Precipitation in Nb-Stabilized Ferritic Stainless Steel Investigated with <Emphasis Type="Italic">in-situ</Emphasis> and <Emphasis Type="Italic">ex-situ</Emphasis> Transmission Electron Microscopy

A Nb-stabilized Fe-15Cr-0.45Nb-0.010C-0.015N ferritic stainless steel is studied with transmission electron microscopy (TEM) to investigate the morphology and kinetics of precipitation. Nb_x(C,N)_y\hbox{Nb}_{x}\hbox{(C,N)}_y and MnS precipitates are present in the steel in the initial condition. Ex-situ TEM analysis is performed on samples heat treated at 973 K, 1073 K, 1173 K, and 1273 K (700 °C, 800 °C, 900 °C, and 1000 °C). Within this temperature range, both Fe₂Nb\hbox{Fe}_2\hbox{Nb} and Fe₃Nb₃X_x\hbox{Fe}_{3}\hbox{Nb}_{3}\hbox{X}_{x} (with X = C or N) precipitates form. Fe₂\hbox{Fe}_2Nb is observed at 1073 K (800 °C).   Fe₃Nb₃X_x\;\hbox{Fe}_{3}\hbox{Nb}_{3}\hbox{X}_{x} precipitates form at the grain boundaries between 973 K and 1273 K (700 °C and 1000 °C). Up to at least 1173 K (900 °C) their fraction increases with time and temperature, but at 1273 K (1000 °C) they lose stability with respect to Nb_x(C,N)_y.\hbox{Nb}_{x}\hbox{(C,N)}_{y}. With in-situ TEM, no phase transition is observed between room temperature and 1243 K (970 °C). At 1243 K (970 °C) the precipitation of Fe₃Nb₃X_x\hbox{Fe}_{3}\hbox{Nb}_{3}\hbox{X}_{x} is observed in the neighborhood of a dissolving Nb₂\hbox{Nb}_2(C,N) precipitate. For sections of grain boundaries where no Nb_x(C,N)_y\hbox{Nb}_x\hbox{(C,N)}_y precipitates are present, Fe₃Nb₃X_x\hbox{Fe}_3\hbox{Nb}_3\hbox{X}_{x} does not form. It is concluded that the precipitation of Fe₃Nb₃X_x\hbox{Fe}_{3}\hbox{Nb}_{3}\hbox{X}_x is directly related to the dissolution of Nb₂\hbox{Nb}_2(C,N) through the redistribution of C or N.     

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.     

Inverse Magnetocaloric Effect Related to the Magnetostructural Phase Transition in Quaternary Ni–Co–Mn–Al Alloy

The inverse magnetocaloric effect of Ni–Co–Mn–Al quaternary alloy with the relatively low material cost is achieved firstly in a theoretical study (V. Sokolovskiy et al.: J. Appl. Phys., 2020, vol. 127, p. 163901). To investigate and prove this study, the exact composition of \(\hbox{Ni}_{{40}}\hbox{Co}_{{10}}\hbox{Mn}_{{36}}\hbox{Al}_{{14}}\) alloy is selected and and explored by the combination of X-ray diffraction, scanning electron microscopy, resistivity, and magnetic studies. The quaternary alloy reveals that the main phase is associated with a martensitic L1₀ phase structure with some austenitic B2 phase in the vicinity of room temperature. The results show that the alloy maintains both Austenite and Martensite phases and has a grand scale change in magnetization of approximately 95 emu \(\hbox{g}^{-1}\) around the Martensitic phase transition (in the range of 20 K) that exhibits a first-order magnetic transition from ferromagnetic to non-ferromagnetic state. The alloy reveals the inverse magnetic entropy change of about 12 and 8 J \(\hbox{kg}^{-1}\,\hbox{K}^{-1}\) and the relative cooling power of 125 and 76 J kg⁻¹ under only 15 and 10 kOe, respectively. Likewise, the MR value of 11.5 pct obtains in the external magnetic field source of 10 kOe in the heating direction. The experimental results support the referenced theoretical study and make this material prominent in future magnetocaloric and magnetoresistivity studies.

     

First-Principles Calculations on Stabilization of Iron Carbides (Fe3C, Fe5C2, and η-Fe2C) in Steels by Common Alloying Elements

The control of carbide formation is crucial for the development of advanced low-alloy steels. Hence, it is of great practical use to know the (de)stabilization of carbides by commonly used alloying elements. Here, we use ab initio density functional theory (DFT) calculations to calculate the stabilization offered by common alloying elements (Al, Si, P, S, Ti, V, Cr, Mn, Ni, Co, Cu, Nb, Mo, and W) to carbides relevant to low-alloy steels, namely cementite $(\hbox{Fe}_{3}\hbox{C}),$ H?gg $(\hbox{Fe}_{5}\hbox{C}_{2}),$ and eta-carbide $(\eta{\text{-}}\hbox{Fe}_{2}\hbox{C})$ . All alloying elements are considered on the Fe sites of the carbides, whereas Al, Si, P, and S are also considered on the C sites. To consider the effect of larger supercell size on the results of (de)stabilization, we use both 1?×?1?×?1 and 2?×?2?×?2 supercells in the case of $\hbox{Fe}_{3}\hbox{C}.$      

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. $$\begin{gathered} \log {\text{ }}\gamma _{{\text{Pb}}} {\text{ = - 0}}{\text{.255 }}N_{Bi}^2 {\text{ }}N_{{\text{Bi}}} {\text{< 0}}{\text{.8}} \hfill \\ \log {\text{ }}\gamma _{{\text{Tl}}} {\text{ = - 0}}{\text{.805 }}N_{Bi}^2 {\text{ }}N_{{\text{Bi}}} {\text{< 0}}{\text{.7}} \hfill \\ \end{gathered} $$      

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.     

Predictions for the Sound Velocity in Various Liquid Metals at Their Melting Point Temperatures

Sound velocity values for 32 liquid metals at their melting point temperatures have been predicted using two models that we presented; most of these metals are transition and rare earth metals. The sound velocities for most of these liquid metals have yet to be measured experimentally. Dimensionless common parameters, denoted by x_\textT^1/2 \xi_{\text{T}}^{1/2} and x_\textE^1/2 , \xi_{\text{E}}^{1/2} , were determined on the basis of the predicted sound velocities. These common parameters, which characterize the liquid state (i.e., an atom’s hardness or softness and its anharmonic motions), allow for better predictions of several thermophysical properties (e.g., surface tension, viscosity, self-diffusivity, volume expansivity) of liquid metallic elements. The values of both the common parameters x_\textT^1/2 \xi_{\text{T}}^{1/2} and x_\textE^1/2 \xi_{\text{E}}^{1/2} vary periodically with atomic number. Using our viscosity model in terms of the parameter x_\textT^1/2 , \xi_{\text{T}}^{1/2} , values of melting point viscosity were calculated for liquid molybdenum and platinum. The agreement obtained between calculated and experimental values is good when using predicted values of x_\textT^1/2 \xi_{\text{T}}^{1/2} to calculate their viscosities.     

Thermodynamics of the system NaCl-AlCl3

Heat capacities of melts were measured in the range 400 to 1100 K and 0.48 < N_AlCl3 < 0.62, the results being expressed by C_p = 40.96 – 0.0295T + 2.01 × 10^?5 T ² J K^?1 g·atom^?1 i.e., AlCl₃ contains 4 atoms, and so forth). This equation was used in interpreting literature vapor pressure data. Measurements were made of the emf of the concentration cell $$AL\left| {_{AlCl_3 }^{NACl(sat)} } \right.\left| {_{(Na^ + )}^{Pyrex} } \right.\left| {_{AlCl_3 }^{NACl(sat)} } \right.\left| {_{AlCl_3 }^{NACl} } \right.\left| {AL} \right. $$ at temperatures 473 to 623 K, and the results were correlated with the vapor pressure data to yield activities of NaCl and AlCl₃. Measurements with a sodium electrode confirmed the accepted values for the free energy of formation of A1C1₃ within about 1.5 kJ mol^?1. The activities were used to analyze the phase diagram. Direct measurement of the eutectic temperature with a concentration-cell technique (which avoids supercooling) gave 386 K; the eutectic composition is 60.0 mol pct A1C1₃. The standard entropy of NaAlCl₄(s) is S _298.15 ^° = 199.1 J K^?1 mol^?1. The free energy for NaAlCl₄(l) = NaAlCl₄(g) is ΔG° = 82740 ?63.66T J mol^?1 at around 950 K.     

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.     

Enthalpy of mixing of liquid Ni-Zr and Cu-Ni-Zr alloys

The partial (Δ and the integral (ΔH) enthalpies of mixing of liquid Ni-Zr and Cu-Ni-Zr alloys have been determined by high-temperature isoperibolic calorimetry at 1565 ± 5 K. The heat capacity (C _p) of liquid Ni₂₆Zr₇₄ has been measured by adiabatic calorimetry (C _p=53.5±2.2 J mol⁻¹ K⁻¹ at 1261±15 K). The integral enthalpy of mixing changes with composition from a small positive (Cu-Ni, ΔH (x _Ni=0.50, T=1473 to 1750 K)=2.9 kJ mol⁻¹) to a moderate negative (Cu-Zr; ΔH(x _Zr=0.46, T=1485 K)=−16.2 kJ mol⁻¹) and a high negative value (Ni-Zr; ΔH(x _Zr=0.37, T=1565 K)=−45.8 kJ mol⁻¹). Regression analysis of new data, together with the literature data for liquid Ni-Zr alloys, results in the following relationships in kJ mol⁻¹ (standard states: Cu (1), Ni (1), and Zr (1)):for Ni-Zr (1281≤T≤2270 K),

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.     

Evaluation of unified interaction parameter model parameters for calculating activities of ferromanganese alloys: Mn-Fe-C, Mn-Fe-Si, Mn-C-Si, and Mn-Fe-C-Si systems

Interaction parameters for Mn-based alloys were evaluated using both carbon solubility and activity data for species in binary and ternary manganese alloys. The parameters at 1400 °C are the following

Thermodynamics of iron oxide in Fe x O-dilute CaO+Al2O3+Fe x O fluxes at 1873 K

The distribution of iron between Fe _x O-dilute CaO+Al₂O₃+Fe _x O fluxes and Pt+Fe alloys, as well as the redox equilibrium of iron ions in these fluxes, was experimentally investigated in pressure-controlled CO₂/CO gas at 1873 K. Total iron content in flux (pct Fe ^T ) and the ratio of (pct Fe²⁺) to (pct Fe ^T ) in fluxes with constant can be related to the activity of iron, α _Fe, and the partial pressure of oxygen, a _Fe, using the following equation:

Hydrogen Solubility in Pd/V2O5 Composites Formed by Internal Oxidization and Hydrogen Isotherms in Un-oxidized Pd-V Alloys

Pd-V alloys were internally oxidized (IOed) resulting in composites of nano-particle V₂O₅ precipitates within Pd matrices. These composites were found to interact with H₂ to form hydrogen bronzes, H _x V₂O₅, within the Pd matrix where x can vary between 1.65 and 2.20. Relative partial molar enthalpies for H intercalation into the H-bronze within the Pd/V₂O₅ composite were measured calorimetrically as a function of the H content of the bronze, and these molar enthalpies decrease in magnitude from about ?75 to ?20 kJ/mol H as the H content increases. H₂ isotherms have also been measured in disordered, fcc Pd_0.96V_0.04, Pd_0.945V_0.055, and Pd_0.93V_0.07 alloys from 273 K to 343 K (0 °C to 70 °C). Thermodynamic data have been derived from these isotherms. The relative partial molar enthalpies at infinite dilution of H, $\Updelta H_{\hbox{H}}^\circ,$ increase with atom fraction V, X $_{\hbox{V}},$ while the corresponding standard partial molar entropies, $\Updelta \hbox{S}_{\hbox{H}}^\circ,$ decrease with $\hbox{X}_{\hbox{V}}.$ The first-order term, g₁, in a polynomial expansion of the excess or non-ideal chemical potential of H in r = H-to-metal, mol ratio, decreases in magnitude with $\hbox{X}_{\hbox{V}}$ at a given temperature.     

Calculation of the product phase grain boundary area during solid state transformations

A general formal expression is derived for the calculation of product phase grain boundary area per unit volumeA _v at any time during solid state transformations occurring by nucleation and growth process. It is assumed that the spatial distribution of the product phase particles is random and dis-crete product phase particles have a spherical shape. The analysis is applicable to any arbitrary nu-cleation and growth kinetics. For the sizeindependent growth rate,A _v is given by: whereV _{V ex} is the extended volume fraction of the product phase, andS _{V ex} is the total extended product phase-matrix interfacial area per unit volume. If the growth rate depends onparticle size or particle size and time, then, The results are applicable to any arbitrary functional form of nucleation rate. The result for size de-pendent growth rate is approximate; however, the error involved in this approximation is less than ±10 pct. The analysis demonstrates thatA _v , and also the grain size of transformed structure, are basically determined by thepath of microstructural evolution described by the variation of product phase-matrix interfacial area per unit volume with the product phase volume fraction, and do not explicitly depend on any other variables. The analysis is also applicable to nonisothermal and con-tinuous cooling transformations. On leave from the Department of Metallurgical Engineering, Indian Institute of Technology, Kanpur-208016, India