Magnesium-rare earth phase diagrams: Experimental investigation of the Ho-Mg system

The Ho-Mg phase diagram was determined in the range 0 to 100 at.% Mg by differential thermal analysis (DIA), X-ray powder diffraction, optical and scanning electron microscopy, and electron probe microanalysis (EPMA). Four intermediate phases were found to exist, and their crystal structures were confirmed or determined as the following: (1) β phase, Ho₂Mg, cubic,cI2, W type, peritectic formation 1170 °C; (2) HoMg, cubic,cP2, CsCl type, peritectic formation 845 °C; (3) HoMg2, hexagonal,hP12, MgZn₂ type, peritectic formation 695 °C; and (4) Ho₅Mg₂₄, cubic, cI58, α Mn-type, peritectic formation 600 °C. The β phase undergoes a eutectoidal decomposition at 685 °C and 29.0 at.% Mg. A eutectic reaction was observed to occur at 555 °C and 90.0 at.% Mg. The data obtained in this study are compared with those of other previously studied RE-Mg phases and briefly discussed. “Mischmetal” is the tradename of an inter-rare earth alloy ideally containing: 27 at.% La, 48 at.% Ce, 5 at.% Pr, 16 at.% Nd, and 4 at.% other rare earths. This composition is close to that of a typical rareearth ore.     

Kinetics of Cr/Mo-rich precipitates formation for 25Cr-6.9Ni-3.8Mo-0.3N super duplex stainless steel

The amount and composition of Cr-rich (σ) and Mo-rich (χ) precipitates in super duplex stainless steels was analyzed. An isothermal heat treatment was conducted at temperatures ranging from 700 °C to 1000 °C for up to 10 days. A time-temperature transformation (TTT) diagram was constructed for the mixture of σ and χ phases. The mixture of the σ and χ phases exhibited the fastest rate of formation at approximately 900 °C. Minor phases, such as Cr₂N, M₂₃C₆, and M₇C3, were also detected using a transmission electron microscopy (TEM). Also, a continuous cooling transformation (CCT) diagram was constructed for the mixture of σ and χ phases using the Johnson-Mehl-Avrami equation. Compared with the known CCT diagram of the σ phase, this study revealed faster kinetics with an order of magnitude difference and a new CCT diagram was also developed for a mixture of σ and χ phases. The calculated fraction of σ and χ phases obtained at a cooling speed of 0.5 °C/s was in good agreement with the experimental data.     

A contribution to the rare earth intermetallic chemistry: Praseodymium-magnesium alloy system

Pr-Mg alloys were studied in the range 0–100 at. % Mg. By using X-ray powder diffraction, optical and scanning electron microscopy, electron probe micro-analysis and differential thermal analysis, the different intermediate phases were identified and their crystal structures confirmed or determined. The following phase equilibria have been also determined: PrMg (cubic, cP2, CsCl type, melting point 765°C), PrMg₂ (cubic, cF24, MgCu₂ type, peritectic formation 740°C), PrMg₃ (cubic, CF16-BiF₃ type, melting point 790°C), Pr₅Mg₄₁ (tetragonal, tI92 Ce₅Mg₄₁-type, peritectic formation 575°C) and PrMg₁₂ (tetragonal, tI26 ThMn₁₂-type, peritectic formation 565°C) PrMg₂ undergoes a eutectoidal decomposition at 670°C. Three eutectic reactions were observed to occur at 735°C and 40·0 at. % Mg, at 725°C and 59·5 at. % Mg and at 560°C and 95·0 at. % Mg, respectively. The (β-Pr) terminal solid solution was observed to decompose eutectoidally at 510°C and 19·5 at. % Mg. The data obtained in this study are compared with those relating to other previously studied R-Mg systems.The crystallochemical characteristics of the binary phases formed by Mg with the rare earths and with the alkaline earths are briefly discussed.     

The phase evolution and microstructural stability of an orthorhombic Ti-23Al-27Nb alloy

The phase evolution and microstructural stability were studied for an orthorhombic Ti-23Al-27Nb alloy. Monolithic sheet materials were produced through conventional thermomechanical processing techniques comprising nonisothermal forging and pack rolling. Phase evolution studies showed that, depending on the heat treatment schedule, this alloy may contain several constituent phases including: α₂ (ordered close-packed hexagonal D0₁₉ structure), B2 (ordered body-centered cubic, bcc), β (disordered bcc), and O (ordered orthorhombic based on Ti₂AlNb). Differential thermal analysis studies indicated that the B2 transus temperature was 1070 °C. Heat treatment and transmission electron microscopy studies showed that the α₂+B2 phase field extended between 1010 and 1070 °C. From 875 to 975 °C, a two-phase O+B2 field existed. Sandwiched between these two-phase regimes was a narrow three-phase α₂+B2+O field. Below 875°C, an O+β field existed. All heat treatments at or above 875°C, followed by quenching, resulted in equiaxed microstructures. However, below 875 °C, the B2 phase transformed into a mixture of O and bcc phases with lenticular morphologies. Cellular precipitation of O+β platelets at O/B2 and α₂/B2 grain boundaries occurred depending on solutionizing and aging temperatures, which is explained by the compositional gradient between the bcc phases.     

The phase evolution and microstructural stability of an orthorhombic Ti-23Al-27Nb alloy

The phase evolution and microstructural stability were studied for an orthorhombic Ti-23Al-27Nb alloy. Monolithic sheet materials were produced through conventional thermomechanical processing techniques comprising nonisothermal forging and pack rolling. Phase evolution studies showed that, depending on the heat treatment schedule, this alloy may contain several constituent phases including: α₂ (ordered close-packed hexagonal D0₁₉ structure), B2 (ordered body-centered cubic, bcc), β (disordered bcc), and O (ordered orthorhombic based on Ti₂AlNb). Differential thermal analysis studies indicated that the B2 transus temperature was 1070 °C. Heat treatment and transmission electron microscopy studies showed that the α₂+B2 phase field extended between 1010 and 1070 °C. From 875 to 975 °C, a two-phase O+B2 field existed. Sandwiched between these two-phase regimes was a narrow three-phase α₂+B2+O field. Below 875°C, an O+β field existed. All heat treatments at or above 875°C, followed by quenching, resulted in equiaxed microstructures. However, below 875 °C, the B2 phase transformed into a mixture of O and bcc phases with lenticular morphologies. Cellular precipitation of O+β platelets at O/B2 and α₂/B2 grain boundaries occurred depending on solutionizing and aging temperatures, which is explained by the compositional gradient between the bcc phases.     

Determination of the Fe-rich portion of the Fe-Ni-S phase diagram

The low-temperature, Fe-rich portion of the Fe-Ni-S phase diagram was determined from Fe-Ni-S alloys (2.5,5,10,20, and 30 wt.% Ni, 10 wt % S, balance Fe) heat treated at 100 °C intervals from 900 to 300 °C. The microstructure and microchemistry of the phases in the heat treated Fe-Ni-S alloys were studied using a high-resolution field-emission gun (FEG) scanning electron microscope (SEM), electron probe microanalyzer (EPMA), and analytical electron microscope (AEM). Tieline compositions were obtained by determining the average phase composition and by measuring compositional profiles across interphase interfaces with the EPMA and AEM. At 600 °C and below, at least one phase was <1 Μm in size requiring the use of the AEM for analysis. The measured α + FeS, γ+ FeS, and α + γ + FeS boundaries in the Fe-rich corner of the Fe-Ni-S isotherms are consistent with previous studies. However, two new phases were observed for the first time coexisting with γ and FeS phases: FeNiγ′′ (∼52 wt.% Ni) at 600 and 500 °C and Ni ₃ Fe, ordered Ll ₂,γ′ (∼64 wt.% Ni) at 400 °C. New ternary isotherms are given at 600,500, and 400 °C that include the newly determined γ+γ′′ + FeS and the γ + γ′ + FeS three-phase fields. The effects of S on the phase boundaries of the α + γ phase field and the application of the Fe-Ni-S phase diagram to explain the microstructure and microchemistry of the metallic phases of stony meteorites are also discussed.     

The Al-Er-Mg ternary system Part I: Experimental investigation

Investigation of phase equilibria in the ternary system Al-Er-Mg has been carried out by means of differential thermal analysis (DTA), powder x-ray diffraction (XRD), light optical microscopy (LOM), scanning electron microscopy (SEM), and quantitative electron probe microanalysis (EPMA). The isothermal section at 400 °C has been established. An extended homogeneity region with Al substitution for Mg at a constant Er content has been found for (Mg_1−x Al _x )Er (0≤x≤0.78); a few other boundary binary phases give lower ternary solubility. A ternary compound, τ, of Al_66.7Er₁₀Mg_23.3 stoichiometry, has been found to exist in the isothermal section at 400 °C.     

The Al-Er-Mg ternary system Part I: Experimental investigation

Investigation of phase equilibria in the ternary system Al-Er-Mg has been carried out by means of differential thermal analysis (DTA), powder x-ray diffraction (XRD), light optical microscopy (LOM), scanning electron microscopy (SEM), and quantitative electron probe microanalysis (EPMA). The isothermal section at 400 °C has been established. An extended homogeneity region with Al substitution for Mg at a constant Er content has been found for (Mg_1−x Al _x )Er (0≤x≤0.78); a few other boundary binary phases give lower ternary solubility. A ternary compound, τ, of Al_66.7Er₁₀Mg_23.3 stoichiometry, has been found to exist in the isothermal section at 400 °C.     

On the Quaternary System Ti-Fe-Ni-Al

The homogeneity ranges of the Laves phases and phase relations concerning the Laves phases in the quaternary system Ti-Fe-Ni-Al at 900 °C were defined by x-ray powder diffraction (XPD) data and electron probe microanalysis (EPMA). Although at higher temperatures the Laves phase forms a continuous solid solution, two separate homogeneity fields of TiFe₂-based (denoted by λ_Fe) and Ti(TiNiAl)₂-based (denoted by λ_Ni) Laves phases appear at 900 °C. The relative locations of Laves phases, G phase, Heusler phase, and CsCl-type phase as well as the associated tie-tetrahedra were experimentally established in the quaternary for 900 °C and presented in three-dimensional (3D) view. Furthermore, a partial isothermal section TiFe₂-TiAl₂-TiNi₂ was constructed, and a connectivity scheme, derived for equilibria involving Laves phases in the Ti-Fe-Ni-Al quaternary system at 900 °C was derived. As a characteristic feature of the quaternary phase diagram, the solid solubility of fourth elements in both the TiFe₂-based and Ti(NiAl)₂-based Laves phases is limited at 900 °C and is dependent on the ternary Laves phase composition. A maximum solubility of about 8 at.% Ni is reached for composition Ti_33.3Fe_33.3Al_33.4. Structural details have been evaluated from powder x-ray and neutron diffraction data for (i) the Ti-Fe-Ni ternary and the Ti-Fe-Ni-Al quaternary Laves phases (MgZn₂-type, space group: P6₃/mmc) and (ii) the quaternary G phase. Atom site occupation behavior for all phases from the quaternary system corresponds to that of the ternary systems. For the quaternary Laves phase, Ti occupies the 4f site and additional Ti (for compositions higher than 33.3 at.%Ti) preferably enters the 6h site. Aluminum and (Fe,Ni) share the 6h and the 2a sites. The compositional dependence of unit cell dimensions, atomic coordinates, and interatomic distances for the Laves phases from the quaternary system is discussed. For the quaternary cubic G phase, a centrosymmetric as well as a noncentrosymmetric variety was observed depending on the composition: from combined x-ray and neutron powder diffraction measurements Ti_33.33Fe_13.33Ni_10.67Al_42.67 was found to adopt the lower symmetry with space group .     

The Fe-A1-Ti system

An extensive experimental investigation of the Fe-Al-Ti system by metallography, microprobe analysis, and XRD on quenched specimens and on diffusion couples is presented. Two isothermal sections at 800 and 1000 °C were established; they differ substantially from the existing (800 °C) or partly determined (1000 °C) diagrams. From these results, existence of the τ₁ phase (Fe₂AlTi) can be ruled out. Existence of the ternary compounds, τ₂ (Al₂FeTi) and τ₃ (Al₂₂Fe₃Ti₈), is confirmed. The composition limits of both phases were determined; they differ considerably from those given in earlier reports. The τ₂ phase apparently exists in a cubic and a tetragonal polymorph, depending on composition. The cubic form exists at high titanium contents. At 1000 °C, the two polymorphs are separated by a miscibility gap. At compositions where the “X phase” (Al₆₉Fe₂₅Ti₆) was previously reported, single-phase samples were obtained at both temperatures. From the present results, there is no evidence to assume that this is a new ternary phase rather than the ternary homogeneity range of the Al₃Fe phase. In addition, extensions of the binary intermetallic phases into the ternary system were determined.     

The antimony-iron-niobium (Sb-Fe-Nb) system

Phase equilibria were established in the Nb-Fe-Sb ternary system for an isothermal section at 600 °C. Investigation of the phase relations was based on light optical microscopy, electron probe microanalysis, and X-ray diffraction experiments on arc-melted bulk alloys, which were annealed up to 1400 h. One ternary compound was observed: NbFeSb (MgAgAs type) without a significant homogeneity region at 600 °C. Except for NbFe_2−y , mutual solid solubilities at 600 °C were found to be very low, for example, <1 at.% Fe and <1 at.% Nb in the binary Nb antimonides and Fe antimonides, respectively. The binary Laves phase NbFe_2−y with the MgZn₂ type exhibits an extended homogeneity region dissolving at 600 °C up to 7 at.% Sb in the ternary without change of its structure type.     

Phase equilibria and multiphase reaction diffusion in the Cr-C and Cr-N systems

The phase formation in the Cr-C and Cr-N systems was investigated using reaction diffusion couples. The carbides were prepared by reaction of chromium metal with graphite powder in the range 1143 to 1413 °C in argon atmosphere; the nitride samples by reaction of the metal with N₂ (≤31 bar) in the range 1155 to 1420 °C. While the carbide samples showed the three chromium carbide phases in form of dense diffusion layers between 1100 and 1400 °C, porosity occurred at temperatures above 1400 °C. The composition of the phase bands was measured by the means of electron probe microanalysis. For the Cr₂₃C₆ phase, a slightly higher C composition was found than given in the literature. In Cr-N diffusion couples both the δCrN_1−x and βCr₂N formed phase bands at T≥1150 °C. Because decomposition processes occurred upon cooling, quenching experiments were carried out in the range 1370 to 1420 °C at 31 bar N₂ to stabilize the phases. The EPMA investigations of the homogeneity ranges yielded a large increase of the homogeneity range for δCrN_1−x with increasing temperature. The nonmetal diffusion coefficients in all phases of both systems were calculated from layer growth and/or from concentration profiles. In δCrN_1−x the N diffusivity was found to be strongly dependent on the composition. The Vickers microhardnesses of the various phases were obtained by measuring the diffusion layers.     

Phase equilibria and multiphase reaction diffusion in the Cr-C and Cr-N systems

The phase formation in the Cr-C and Cr-N systems was investigated using reaction diffusion couples. The carbides were prepared by reaction of chromium metal with graphite powder in the range 1143 to 1413 °C in argon atmosphere; the nitride samples by reaction of the metal with N₂ (≤31 bar) in the range 1155 to 1420 °C. While the carbide samples showed the three chromium carbide phases in form of dense diffusion layers between 1100 and 1400 °C, porosity occurred at temperatures above 1400 °C. The composition of the phase bands was measured by the means of electron probe microanalysis. For the Cr₂₃C₆ phase, a slightly higher C composition was found than given in the literature. In Cr-N diffusion couples both the δCrN_1−x and βCr₂N formed phase bands at T≥1150 °C. Because decomposition processes occurred upon cooling, quenching experiments were carried out in the range 1370 to 1420 °C at 31 bar N₂ to stabilize the phases. The EPMA investigations of the homogeneity ranges yielded a large increase of the homogeneity range for δCrN_1−x with increasing temperature. The nonmetal diffusion coefficients in all phases of both systems were calculated from layer growth and/or from concentration profiles. In δCrN_1−x the N diffusivity was found to be strongly dependent on the composition. The Vickers microhardnesses of the various phases were obtained by measuring the diffusion layers.     

Experimental Phase Diagram Investigations in the Ni-Rich Part of Al-Fe-Ni and Comparison with Calculated Phase Equilibria

The phase equilibria in the section Ni₃Fe-Ni₃Al and phase boundaries of γ′(Ni₃Al) at 1000 °C were studied by a combination of powder X-ray diffraction (XRD), differential thermal analysis (DTA), and electron probe microanalysis (EPMA). The existence of a continuous solid solution at 450 °C was confirmed by a linear decrease of the lattice parameter from Ni₃Al to Ni₃Fe. The phase boundaries of γ′ with γ and the B2-type phase were determined at 1000°C by EPMA. A vertical section of the phase diagram from Ni₃Al to Ni₃Fe above 450 °C, including the liquidus temperatures, is proposed based on the DTA investigations. The invariant four-phase equilibrium U: L + γ′ = γ + B2 is found to occur at 1366 ± 1 °C. The experimental data are compared with a calculated phase diagram obtained by extrapolation from the corresponding binary data sets.     

<Superscript>63</Superscript>Cu NMR spectra,magnetic susceptibility,and transmission electron microscopy of the rapidly quenched alloy Ti<Subscript>50</Subscript>Ni<Subscript>25</Subscript>Cu<Subscript>25</Subscript>

Methods of transmission and scanning electron microscopy and nuclear magnetic resonance (NMR) at ⁶³Cu nuclei, as well as measurements of the static magnetic susceptibility χ(T) have been used to study a shape-memory alloy (SMA) Ti₅₀Ni₂₅Cu₂₅, which experiences a thermoelastic martensite transformation. The alloy was obtained from an amorphous ribbon in a bimodal nano- and submicrocrystalline state via a crystallization annealing for 1 h at 770 K with a subsequent quenching to room-temperature water. The resultant B2 austenite is characterized by a fine structure of the ⁶³Cu NMR spectra, which is connected with the different distribution of ⁶³Cu atoms on the second coordination shell. The evolution of the shape of the spectra with decreasing temperature reveals a structural transition B2 → B19. In addition, the ⁶³Cu NMR spectra, just as the transmission electron microscopy, indicate the presence of phase separation in the alloy, with the precipitation of a TiCu (B11) phase. The temperature dependence of the static magnetic susceptibility χ(T) also indicates the occurrence of a structural transition and has a hysteretic nature of “stepped” type. The discovered stepped nature of the χ(T) dependence is explained by the bimodal size distribution of grains of the B2 phase due to the size effect of the martensitic transformation.     

The Y–Cu–Mg system in the 0–66.7 at.% Cu concentration range: The isothermal section at 400 °C

Synthesis and characterization of about fifty alloys were performed in order to construct the isothermal section of the Y–Cu–Mg ternary system at 400 °C in the 0–66.7 at.% Cu concentration range. Scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDXS) and X-ray powder diffraction (XRPD) techniques were used to examine microstructures, identify phases and define their compositions and crystal structures. Phase equilibria in the investigated compositional region are characterized by the absence of extended ternary solid solutions and by the presence of at least ten ternary phases. Crystal structures of the previously reported Y₂Cu₂Mg, Y₅Cu₅Mg₈, Y₅Cu₅Mg₁₃, Y₅Cu₅Mg₁₆ and YCuMg₄ phases were confirmed. A ternary phase with homogeneity range around the YCu₄Mg stoichiometry was found, crystallizing in the cF24--MgCu₄Sn structure type; at 400 °C this phase coexists with a ternary solid solution based on the binary Laves phase Cu₂Mg, which dissolves about 5 at.% Y. The equiatomic YCuMg phase was also found to exist: from the analysis of X-ray powder patterns it is suggested to crystallize in the hP9--ZrNiAl structure type (a = 0.74449(4) nm, c = 0.39953(2) nm). Two other stoichiometric ternary phases were detected, of approximate compositions Y₂₅Cu₁₈Mg₅₇ and Y₁₃Cu₉Mg₇₈, whose crystal structures are still unknown. In the Mg-rich region, a ternary phase forms characterized by a large homogeneity region.     

The antimony-iron-zirconium (Sb-Fe-Zr) system

Phase equilibria were established in the Fe-Sb-Zr ternary system below 60 at.% Sb for an isothermal section at 800 °C; the very Sb-rich region was studied at 600 °C. Investigation of the phase relations was based on light optical microscopy, electron probe microanalysis, and X-ray diffraction experiments on arc melted bulk alloys, which were annealed up to 350 h. Three ternary compounds were observed: ZrFe_1−x Sb (0.3<x<0.5; defect TiNiSi Type), Zr₆Fe_1−x Sb_2+x (0<x<0.24; ordered Fe₂P type) and Zr₅Fe _x Sb_3−x (x=0.44; W₅Si₃ type). Whereas Zr₅Fe_0.44Sb_2.56 at 800 °C formed at a given composition without a significant homogeneity region, a rather extended solid solution up to about 9 at.% Sb was observed for the Laves phase Zr(Fe_1−x Sb _x )_2−y . At 800 °C, binary ZrFe_2−y was only observed with the cubic MgCu₂ type; Sb content of more than about 3 at.% Sb stabilized the hexagonal MgNi₂ type in the Zr-poor end of the homogeneity region. The MgCu₂ type prevails at higher Sb content of up to 9 at.% Sb at the Zr-rich side. Zr₃Sb (Ni₃P type) seems to dissolve up to 3.5 at.% of Fe replacing Zr in the structure. Binary Zr₅Sb_3+x (Ti₅Ga₄ type or filled Mn₅Si₃ type) dissolves up to 11 at.% Fe by gradually replacing the Sb atoms in the octahedral sites with Fe; thus the boundary of the homogeneous region on the Fe-rich side essentially corresponds to the ternary limit at Zr₅FeSb₃. For stoichiometric binary Zr₅Sb₃ the authors observed only a small solubility (<1 at.% Fe) with the Yb₅Sb₃ type. At 800 °C there is practically no solid solubility of Zr in the iron antimonides and of Fe in the zirconium antimonides richer than 40 at.% Sb.     

The antimony-iron-zirconium (Sb-Fe-Zr) system

Phase equilibria were established in the Fe-Sb-Zr ternary system below 60 at.% Sb for an isothermal section at 800 °C; the very Sb-rich region was studied at 600 °C. Investigation of the phase relations was based on light optical microscopy, electron probe microanalysis, and X-ray diffraction experiments on arc melted bulk alloys, which were annealed up to 350 h. Three ternary compounds were observed: ZrFe_1−x Sb (0.3<x<0.5; defect TiNiSi Type), Zr₆Fe_1−x Sb_2+x (0<x<0.24; ordered Fe₂P type) and Zr₅Fe _x Sb_3−x (x=0.44; W₅Si₃ type). Whereas Zr₅Fe_0.44Sb_2.56 at 800 °C formed at a given composition without a significant homogeneity region, a rather extended solid solution up to about 9 at.% Sb was observed for the Laves phase Zr(Fe_1−x Sb _x )_2−y . At 800 °C, binary ZrFe_2−y was only observed with the cubic MgCu₂ type; Sb content of more than about 3 at.% Sb stabilized the hexagonal MgNi₂ type in the Zr-poor end of the homogeneity region. The MgCu₂ type prevails at higher Sb content of up to 9 at.% Sb at the Zr-rich side. Zr₃Sb (Ni₃P type) seems to dissolve up to 3.5 at.% of Fe replacing Zr in the structure. Binary Zr₅Sb_3+x (Ti₅Ga₄ type or filled Mn₅Si₃ type) dissolves up to 11 at.% Fe by gradually replacing the Sb atoms in the octahedral sites with Fe; thus the boundary of the homogeneous region on the Fe-rich side essentially corresponds to the ternary limit at Zr₅FeSb₃. For stoichiometric binary Zr₅Sb₃ the authors observed only a small solubility (<1 at.% Fe) with the Yb₅Sb₃ type. At 800 °C there is practically no solid solubility of Zr in the iron antimonides and of Fe in the zirconium antimonides richer than 40 at.% Sb.     

Experimental determination and thermodynamic calculation of phase equilibria in the Fe−Mn−Al system

The phase equilibria among the face-centered cubic (fcc), body-centered cubic (bcc), and βMn phases at 800, 900, 1000, 1100, and 1200 °C were examined by electron probe microanalysis (EPMA), and the A2/B2 and B2/D0₃ ordering temperatures were also determined using the diffusion couple method and differential scanning calorimetry (DSC). The critical temperatures for the A2/B2 and B2/D0₃ ordering were found to increase with increasing Mn content. Thermodynamic assessment of the Fe−Mn−Al system was also undertaken with use of experimental data for the phase equilibria and order-disorder transition temperatures using the CALPHAD (Calculation of Phase Diagrams) method. The Gibbs energies of the liquid, αMn, βMn, fcc, and ε phases were described by the subregular solution model and that of the bcc phase was represented by the two-sublattice model. The thermodynamic parameters for describing the phase equilibria and the ordering of the bcc phase were optimized with good agreement between the calculated and experimental results. This paper was presented at the International Symposium on User Aspects of Phase Diagrams, Materials Solutions Conference and Exposition, Columbus, Ohio, 18–20 October, 2004.