In the field of aerospace engineering, the demand for lightweight, high-strength, and high-temperature-resistant materials has driven extensive research into intermetallic compounds. Among these, titanium aluminide (TiAl) alloys stand out due to their low density, excellent specific strength, and superior creep and oxidation resistance at elevated temperatures, making them ideal candidates for replacing nickel-based superalloys in components such as high-pressure compressor and low-pressure turbine blades. However, the intrinsic brittleness of TiAl alloys often necessitates precision casting as the primary manufacturing route for structural casting parts. A critical challenge arises from the as-cast microstructure, which typically consists of coarse columnar grains forming a fully lamellar (FL) structure. This coarse microstructure leads to low strength and poor ductility, severely limiting the practical application of these casting parts. Therefore, heat treatment is essential to refine the microstructure and enhance mechanical properties. In this study, I focus on the mechanisms of microstructure control through heat treatments for TiAl alloy casting parts, specifically examining the evolution of duplex (DP) and nearly lamellar (NL) microstructures. By optimizing heat treatment processes, I aim to achieve significant grain refinement and elucidate the underlying mechanisms, providing a foundation for improving the performance of TiAl alloy casting parts in service.
The material investigated in this work is a Ti–47Al–2Cr–2Nb alloy, with atomic percentages of 47.01% Al, 2.02% Cr, 2.09% Nb, 0.13% O, and balance Ti. The alloy was prepared through double vacuum arc melting and a skull melting process, followed by investment casting into preheated ceramic molds to produce cylindrical casting parts. These casting parts were then sectioned into specimens for heat treatment and analysis. To determine appropriate heat treatment parameters, differential scanning calorimetry (DSC) was used to identify the eutectoid temperature ($T_e$) and alpha transus temperature ($T_\alpha$), measured at 1130 °C and 1320 °C, respectively. Based on these critical temperatures, a series of heat treatments were designed and conducted in a vacuum tube furnace under high-purity argon atmosphere, with a heating rate of 10 °C/min and furnace cooling to room temperature. The heat treatment processes included both conventional and optimized schemes. Conventional treatments involved isothermal holdings in the α + γ phase region at temperatures ranging from 1185 °C to 1310 °C for 4 hours, while optimized treatments consisted of a two-step process: first, a DP microstructure treatment at 1185 °C for 6 hours, followed by a short-term NL microstructure treatment at 1330 °C for varying times (0.25 h, 0.5 h, 1 h). Microstructural characterization was performed using optical microscopy (OM) and scanning electron microscopy (SEM) in backscattered electron (BSE) mode. Specimens for OM were etched with Kroll’s reagent, while SEM specimens were examined without etching. Quantitative analysis of grain size and volume fractions was carried out using image analysis software.
The as-cast microstructure of the TiAl alloy casting parts revealed a typical FL structure composed of coarse columnar lamellar colony grains, as observed under OM. These grains exhibited a directional growth pattern from the edges toward the center, with an average grain size of approximately 561.03 µm. Additionally, a small fraction of equiaxed γ grains was present, primarily located at lamellar colony boundaries. SEM-BSE imaging confirmed the presence of γ (TiAl), α₂ (Ti₃Al), and minor B2 phases, with lamellar colonies consisting of alternating γ and α₂ layers. Elemental mapping indicated Al segregation due to dendritic solidification during casting, which influences subsequent phase transformations. This coarse as-cast microstructure underscores the necessity for heat treatment to refine the grains and improve mechanical properties for casting parts.
To assess the effectiveness of heat treatments, I first conducted conventional heat treatments in the α + γ phase region. The microstructural evolution was strongly dependent on temperature. As the heat treatment temperature decreased from 1310 °C to 1185 °C, the volume fraction of equiaxed γ grains increased, while the lamellar colony grain size decreased. Specifically, at 1310 °C, the microstructure remained nearly lamellar (NL) with equiaxed γ grains predominantly at colony boundaries. At lower temperatures, such as 1185 °C, a duplex (DP) microstructure emerged, with equiaxed γ grains distributed both at boundaries and within lamellar colonies. The quantitative data are summarized in Table 1, which shows the variation in grain size and volume fraction with temperature. The equivalent grain size ($d_{eq}$) was calculated using the formula:
$$ d_{eq} = d_l \times V_l + d_\gamma \times V_\gamma $$
where $d_l$ is the lamellar colony grain size, $V_l$ is the volume fraction of lamellar colonies, $d_\gamma$ is the equiaxed γ grain size, and $V_\gamma$ is the volume fraction of equiaxed γ grains. This formula allowed for a comprehensive evaluation of grain refinement. As shown in Table 1, the equivalent grain size decreased with lower heat treatment temperatures, with the most significant refinement observed at 1185 °C. However, even at this temperature, the refinement was limited, indicating the need for optimized processes.
| Heat Treatment Temperature (°C) | Microstructure Type | Lamellar Colony Volume Fraction (%) | Equiaxed γ Volume Fraction (%) | Equivalent Grain Size (µm) | Refinement Relative to As-Cast (%) |
|---|---|---|---|---|---|
| 1310 | NL | 85.2 | 14.8 | 488.01 | -13 |
| 1270 | NL | 72.5 | 27.5 | 363.78 | -35 |
| 1240 | NL | 68.9 | 31.1 | 340.09 | -39 |
| 1200 | NL | 64.3 | 35.7 | 301.92 | -46 |
| 1185 | DP | 65.3 | 34.7 | 259.15 | -54 |
Given the limitations of conventional treatments, I optimized the heat treatment processes to achieve greater grain refinement. For DP microstructure, a prolonged treatment at 1185 °C for 6 hours was applied. This resulted in a significant increase in equiaxed γ volume fraction to 55.04% and a reduction in equivalent grain size to 137.39 µm, representing a 75.51% refinement compared to the as-cast state. For NL microstructure, a two-step process was employed: first, the DP microstructure was obtained via 1185 °C/6 h, followed by a short-term treatment at 1330 °C for 0.25 hours. This yielded an NL microstructure with an equivalent grain size of 335.46 µm, a refinement of 40.21%. Longer durations at 1330 °C (0.5 h and 1 h) led to excessive grain growth, emphasizing the importance of precise time control. The optimized heat treatment parameters and outcomes are summarized in Table 2, highlighting the efficacy of these processes for refining casting parts.
| Heat Treatment Process | Microstructure Type | Equiaxed γ Volume Fraction (%) | Equivalent Grain Size (µm) | Refinement Relative to As-Cast (%) |
|---|---|---|---|---|
| 1185 °C/6 h/Furnace Cooling | DP | 55.04 | 137.39 | 75.51 |
| 1185 °C/6 h + 1330 °C/0.25 h/Furnace Cooling | NL | 6.34 | 335.46 | 40.21 |
| 1185 °C/6 h + 1330 °C/0.5 h/Furnace Cooling | FL | ~0 | 849.46 | -51.45 (Growth) |
| 1185 °C/6 h + 1330 °C/1 h/Furnace Cooling | FL | ~0 | 1550.28 | -176.51 (Growth) |
The mechanical properties of the heat-treated casting parts were evaluated through room-temperature tensile tests. As presented in Table 3, the DP microstructure exhibited superior strength and ductility compared to the NL microstructure. Specifically, the DP casting parts had a yield strength of 377 MPa and an elongation of 2.5%, while the NL casting parts showed 353 MPa and 1.6%, respectively. This enhancement is attributed to the finer grain size and higher volume fraction of equiaxed γ grains in DP microstructure, which facilitate more uniform deformation and reduce stress concentration. These results underscore the importance of microstructure control in optimizing the performance of TiAl alloy casting parts for aerospace applications.
| Microstructure Type | Tensile Strength, Rₘ (MPa) | Yield Strength, Rₚ₀.₂ (MPa) | Elongation, A (%) |
|---|---|---|---|
| DP | 490 | 377 | 2.5 |
| NL | 428 | 353 | 1.6 |
To elucidate the grain refinement mechanisms, I analyzed the microstructural transformations during heat treatment. For DP microstructure formation during the 1185 °C/6 h treatment, the refinement mechanism involves the precipitation of numerous γ grains within the coarse lamellar colonies. This process is driven by two factors: Al element segregation from the as-cast state and continuous coarsening of γ lamellae. During heating, defects at lamellar interfaces, such as ledges and terminations, act as nucleation sites for γ grains via diffusion-controlled processes. The nucleation rate can be described by the classical nucleation theory:
$$ \Delta G^* = \frac{16\pi\gamma^3}{3(\Delta G_v)^2} $$
where $\Delta G^*$ is the critical Gibbs free energy for nucleation, $\gamma$ is the interfacial energy, and $\Delta G_v$ is the volumetric free energy change. In casting parts, Al segregation lowers $\Delta G_v$, promoting γ nucleation. Additionally, the coarsening of γ lamellae disrupts the original lamellar structure, leading to fragmentation of colonies. As a result, the coarse lamellar colonies are broken into smaller units, refining the overall microstructure. This mechanism is particularly effective in casting parts due to their inherent elemental heterogeneity from solidification.
For NL microstructure formation via the two-step treatment, the refinement mechanism involves a phase transformation in the α single-phase region. During the short-term holding at 1330 °C, the γ→α transformation occurs, dissolving the equiaxed γ grains and forming new α grains. The pinning effect of remaining equiaxed γ grains restricts the growth of α grains, as described by the Zener pinning model:
$$ D \propto \frac{r}{f} $$
where $D$ is the grain size, $r$ is the particle radius, and $f$ is the volume fraction of pinning particles. In this case, the equiaxed γ grains act as pinning particles, limiting α grain growth. Upon cooling, these α grains transform into fine lamellar colonies, resulting in a refined NL microstructure. However, if the holding time is too long, complete dissolution of γ grains eliminates pinning, leading to abnormal grain growth—as observed with longer treatments. This highlights the critical balance required in heat treatment design for casting parts.
Further insights into the kinetics of these transformations can be gained from diffusion equations. For instance, the growth of γ grains during DP treatment can be modeled using the parabolic growth law:
$$ x = k\sqrt{t} $$
where $x$ is the growth distance, $k$ is the rate constant, and $t$ is time. Similarly, the dissolution of γ during NL treatment follows an exponential decay function, dependent on temperature and time. These kinetic considerations are vital for optimizing heat treatment schedules for casting parts, ensuring reproducible microstructures.
In practice, the success of heat treatment for TiAl alloy casting parts depends on precise control of parameters. The optimized processes I developed demonstrate that prolonged treatment in the α + γ phase region enhances γ precipitation, while short-term exposure in the α phase region leverages phase transformations for refinement. These findings are applicable to other TiAl-based casting parts, offering a pathway to overcome the limitations of as-cast microstructures. Moreover, the use of computational models, such as phase-field simulations, could further refine these processes by predicting microstructure evolution under varying conditions.
In conclusion, through systematic heat treatment studies, I have demonstrated effective microstructure control for TiAl alloy casting parts. The optimized DP microstructure treatment (1185 °C/6 h) refined the grain size by 75.51%, while the NL microstructure treatment (1185 °C/6 h + 1330 °C/0.25 h) achieved 40.21% refinement. The grain refinement mechanisms involve γ precipitation driven by Al segregation and lamellar coarsening for DP, and γ→α transformation with γ pinning for NL. These processes significantly enhance the mechanical properties, making the casting parts suitable for high-temperature applications. Future work should focus on integrating these heat treatments with advanced manufacturing techniques to produce high-performance TiAl alloy casting parts for the aerospace industry. By mastering microstructure control, we can unlock the full potential of TiAl alloys in next-generation engineering components.