Phase equilibria in niobium rich Nb-Al-Ti alloys

The phase equilibria in the Al–Nb–Hf ternary system at 600 °C and 400 °C were experimentally investigated by X-ray diffraction (XRD), scanning electron microscope (SEM) and electron probe microanalysis (EPMA/WDS). In each isothermal sections, 13 three-phase regions were measured. And their phase region boundaries were precisely determined. Among the actually measured binary compounds, only Al₃Nb, AlNb₂ and AlNb₃ have a large range of solid solubility. In addition, two stable ternary compounds τ1-Al₁₁Nb₄Hf₅ and τ2-Al₂NbHf₂, which had certain solid solubility, were newly discovered. On the basis of the experimental results and reasonable inference, the isothermal sections of ternary system at 600 °C and 400 °C were constructed.

In the present work, a new assessment of the Al–Nb system, using a combined first-principle and CALPHAD approach, is presented. Formation enthalpies of all ordered configurations of the intermetallic phases (σ, D0₂₂ and A15) were estimated from ab initio calculations. The liquid, fcc and bcc phases are described by a substitutional solution model. The intermetallic phases D0₂₂ and A15 are described with the New Approach to the Compound Energy Formalism (NACEF). To model the σ phase, four descriptions are applied. In first, σ phase is described with the five-sublattice (5SL) model using combined CEF model where a modification was applied which allows to respect the formation enthalpies of end-members obtained from ab initio calculations. Then, a comparison with the results obtained using the NACEF approach applied to the five-sublattice (5SL), three-sublattice (3SL) and two-sublattice (2SL) models is presented. In all cases, assessments derived from the different descriptions of σ phase show very good agreement with the available experimental and ab initio data with a limited number of used parameters. In this sense, self-consistent thermodynamic descriptions of the Al–Nb system is provided. Moreover, this work shows from the example of the σ phase that the NACEF approach, contrary to what is commonly considered, allows the compatibility between different descriptions using different numbers of sublattices. This is particularly interesting for the construction of multi-component databases.

The microstructural evolution and mechanical behavior of Nb–35Ti–4C, Nb–35Ti–4C–15Al, Nb–25Ti–8C and Nb–25Ti–8C–15Al (at. %) alloys prepared using traditional arc-melting process were investigated. Nb–Ti–Al–C alloy consists of Nbss phase and eutectic phases of Nbss and (Nb, Ti)C in as-cast state. The addition of 8 at. % C leads to the generation of massive primary (Nb, Ti)C while the addition of 15 at.% Al results in the formation of Nb₃Al phase. After heat treatment, secondary (Nb, Ti)C and secondary Nbss precipitate within Nbss and (Nb, Ti)C, respectively. Meanwhile, massive Nb₃Al phase precipitates from parent Nbss in the alloy containing Al element. Both the precipitation of secondary Nbss and Nb₃Al follows a strict orientation relationship with the parent phase. Mechanical properties of the as-cast and heat-treated specimens at room and evaluated temperatures are significantly increased by adopting C and Al elements, as more (Nb, Ti)C and Nb₃Al acting as the strengthening phase are introduced. Whereas, the plasticity of the alloys decreases owing to the formation of massive brittle carbide and Nb₃Al.

Nb-Ti-Al alloys have attracted much attention as a potential candidate for high-temperature structural applications beyond the scope of Ni-based superalloys. Previously, Nb-Ti-Al alloys were prepared by arc-melting or hot-pressing. In this study, laser forming of the Nb-23Ti-15Al alloy is explored, and the microstructure and mechanical behaviors are systematically investigated. The results indicate that the nearly defectless Nb-Ti-Al alloy with fine dendrites could be obtained through laser forming. Three phases, β, δ and Ti(O,C), present in the alloy, and the β/δ matrix displays a columnar microstructure. Moreover, the Ti(O,C) phase, a solid-solution of TiO or TiC, appears in the alloy. Ti(O,C) is formed due to the introduction of the O and C from the elemental powder, which has a face-centered cubic (FCC) structure and a moderate lattice parameter between that of TiO and TiC. Two morphologies of Ti(O,C), large cobblestone-like particles and small dispersive particles, are observed in the alloy. Forbidden reflections of δ phase occur due to double diffraction and slight superlattice reflections of β phase arise. Furthermore, refined β/δ phases and dispersed small Ti(O,C) result in higher microhardness and fracture toughness. In brief, our results indicate that laser forming is a potential method for manufacturing Nb-Ti-Al alloys with prominent properties.

The formation of Kirkendall porosity and its role in determining β → δ transformation types in Nb-Ti-Al alloys has been reported. The Kirkendall porosity appears in Nb-Ti-Al alloys due to the discrepancy in diffusion rates of Nb, Ti and Al atoms. It is proposed that Kirkendall porosity can serve as a criterion in judging the types of β → δ phase transformations: The appearance of Kirkendall porosity indicates the long-range diffusional transformation and excludes the massive transformation. Moreover, the criterion is also applicable in some other alloy systems in judging the phase transformation types. Besides, Kirkendall porosity in Nb-Ti-Al alloys with different hot-pressing durations exhibits different dimensions.

The isothermal sections of the phase diagram and the liquid projection of the Ti-Al-Nb ternary system as well as the phase transformations of a Ti-46.5Al-16.5Nb alloy were calculated using Pandat software. The solidification sequence and subsequent phase transformations of the alloy were verified by microstructural observations and thermo-analysis (TA). The tie-triangle of the three-phase region of α + β + γ at 1300 °C and the tie line of the α and β phases at 1400 °C were identified using electron probe microanalysis (EPMA). Neither the phase equilibria of the α, β and γ phases at 1300 °C nor the tie line of the α and β phases at 1400 °C are in agreement with the calculated phase diagrams reported in the literature. The solidification sequence and subsequent solid-state phase transformations of this alloy were determined as follows: liquid → liquid + β → β → α + β → α + β + γ → β + γ (The β phase and the B2 phase were not differentiated in the present paper.) The phase transformation temperatures were also measured by differential scanning calorimetry (DSC).