Insights into Optimizing Heat Treatment for Hot Isostatic Pressing of Ti-48Al-3Nb-1.5Ta Alloy Powder

Abstract

In this study, various characterization techniques were utilized to investigate the effects of heat treatments on the microstructure and mechanical properties of Ti-48Al-3Nb-1.5Ta (at. %) alloy prepared by the supreme-speed plasma rotating electrode process and hot isostatic pressing. By comparing the microstructures of the alloy under different heat treatments conditions, it was found that the nearly lamellar structure with a size of about 145 μm is formed by a simple heat treatment (1400 °C/10 min, FC to 1300 °C, AC, 850 °C/3 h/FC). Under this heat treatment condition, the alloy exhibited satisfied mechanical properties, with a tensile fracture strain of 1.2% at room temperature and a tensile fracture strain of 7.5% at 750 °C. No fracture occurred after 225 h when creeping at 750 °C/250 MPa. Ta inhibited the growth of lamellae and the expansion of pores, thereby improving creep performance. In summary, the TiAl alloy with satisfied performance was obtained through a simple heat treatment process, which provides a significant idea for engineering application.

1. Introduction

γ-TiAl alloys have attracted great attention in high temperature structural material because of their excellent performance, such as low density, superior oxidation resistance, and outstanding creep performance [1,2]. In 2011, the TiAl4822 alloy was applied on a large scale to low-pressure turbine blades of the GEnX engine by GE [3]. In 2014, the TNM alloy was applied to the geared turbofan engine by Pratt & Whitney [4]. However, these alloys can only be used below 700 °C due to their insufficient high temperature performance and oxidation resistance. A strong resistance to creep allows for exceptional performance in structural applications at high temperatures [5]. The improvements of high temperature performance and room temperature plasticity for TiAl alloys have been the subject of much research [6,7,8]. The mechanical properties of TiAl alloys can be improved by the addition of Ta and Nb [9,10], especially Ta, which results from microstructure refinement and solution strengthening. The Ti-46Al-8Ta alloy manifests satisfied comprehensive performance [11].

Nevertheless, the high-Ta content raises both the cost and density of the alloy and can lead to severe segregation. The newly developed TiAl alloys with a medium-Nb and low-Ta content has exhibited attractive high temperature oxidation resistance and similar mechanical properties comparable to those of high-Ta alloys [12]. Research shows that the fracture strain of an as-cast Ti-48Al-3Nb-1Ta alloy can reach 1.7% through five-cycle heat treatment [13], and a Ti-48Al-3Nb-1.5Ta alloy can reach 1.0% through multi-step heat treatment [14]. However, complex heat treatments are not conducive to engineering applications, especially for complex parts such as low-pressure turbine blades, in which it is difficult to achieve uniform composition and structure through cyclic heat treatment. It is anticipated that TiAl alloy’s service temperature will rise even further due to improved Nb and Ta segregation, enhanced microstructure homogeneity and grain size refinement achieved through the powder metallurgy (PM) process [15,16], and improved mechanical performance at both room and high temperatures.

Numerous investigations have demonstrated that γ-TiAl alloy’s strength and ductility at both room temperature and high temperature can be enhanced by grain refining [13,17]. One way that phase transition and deformation can improve the strength and toughness significantly is by creating γ/γ twin grain boundaries (GBs) [18]. On the other side, numerous γ/γ twisted GBs can be created during rapid cooling to achieve microstructure refinement, as the initial coarse grains were divided into refined grains by freshly formed GBs [19,20]. When rapidly cooled, the Ta-containing TiAl alloy strongly encourages metastable massive transformation, which breaks up big grains and aids in microstructure refinement and strengthening [19,21]. According to Dey’s proposal [20], the phase transition that occurs during the cooling of γ-TiAl alloys is intimately linked to the development of various types of γ/γ twisted GBs, such as Σ3, Σ5, and Σ11 coincidence site lattice (CSL) GBs. Lower energy GBs are more likely to generate new refined grains and related microstructures during cooling [20,22]. However, up to now, although the Ti-48Al-3Nb-1.5Ta (at. %) alloy has promising performance, the optimal heat treatments system has not yet been determined, especially for the PM alloy, and the effects of different heat treatments on its microstructure and properties remain unclear.

In this work, the effects of different heat treatments on the microstructure and mechanical properties of the hot isostatic pressing of the Ti-48Al-3Nb-1.5Ta alloy powder were studied. Moreover, the role of Ta to obtain excellent high temperature performance was investigated. By optimizing the heat treatment regime, it is expected to provide a reference for the engineering application of the alloy with a satisfied performance prepared by a simple process.

2. Materials and Methods

2.1. Material Preparation

The Ti-48Al-3Nb-1.5Ta (at. %) alloy was prepared by supreme-speed plasma rotating electrode process and hot isostatic pressing (SS-PREP&HIP) by Xi’an Sino-Euro Materials Technologies Co., Ltd. (Xi’an, China). The typical rotating speed of traditional PREP is about 15,000 rpm, while the maximum rotating speed of SS-PREP^® can reach 40,000 rpm, which can obtain the required particle size of powder more flexibly, especially being able to obtain finer particles with the size of minus 53 μm. It can be applied to various powder metallurgy processes and obtain higher tap density. Compared with gas atomization, SS-PREP^® powders show better sphericity and fewer hollow powders, guaranteeing excellent performance in PM parts. The powders were mostly perfectly spherical, and there were many fine grains in each particle. The chemical composition of the powder is shown in

Table 1. The chemical composition of the powder.

at./%				wt./ppm
Ti	Al	Nb	Ta	Fe	C	O	N	H
Bal.	48.7	2.9	1.4	480	46	340	60	<10

. The typical morphology and laser particle size distribution of powders with a range of 45–150 μm are shown in Figure 1, with a tap density of 2.7 g/cm³. Other details about Ti-48Al-3Nb-1.5Ta powders can be found in the previous study [23].

The powders were filled into steel envelopes with the size of φ89 × 240 mm, and then the degassing process was conducted at 400 °C for 4 h in order to remove the air inside the envelopes and prevent the introduction of impurities on the surface of powders. The sealed envelopes after degassing were hot isostatic pressed in an Ar atmosphere for 4 h at 1260 °C and 150 MPa with a heating rate of 4 °C/min and a cooling rate of 5 °C/min to ambient temperature. The temperature and pressure were increased simultaneously.

After the envelopes were removed by wire cutting, the SS-PREP&HIP alloy was cut into 10 × 10 × 10 mm specimens, which were then used for heat treatments in a precision high temperature treatment furnace following the regimes listed in

Table 2. Heat treatment regimes.

No.	Heat Treatment Regimes
T1	1420 °C/10 min, furnace cooling (FC) to 1300 °C, then air cooling (AC)
T2	1420 °C/10 min, FC to 1260 °C, then AC
T3	1400 °C/10 min/AC
T4	1400 °C/10 min, FC to 1340 °C, then AC
T5	1400 °C/10 min, FC to 1300 °C, then AC
T6	1400 °C/10 min, FC to 1300 °C, AC, 850 °C/3 h/FC
T7	1380 °C/10 min /AC
T8	1380 °C/3 min/AC × 3 times + 1280 °C/2 h/AC + 1330 °C/10 min/AC

. According to the metallographic microstructures after quenching, as shown in Figure 2, the α single phase’s transition temperature T_α→α+γ is 1370–1380 °C. In the metallographic structure obtained after holding at 1370 °C and then quenching in water, there are still lamellar structures of α and γ phases. However, the metallographic structure obtained after holding at 1380 °C and then quenching in water is an α single-phase structure. After cooling from the α single phase, nearly lamellar or fully lamellar structures are anticipated; the microstructure will be influenced by the holding temperature, duration, and cooling rate, which leads to the development of the regimes displayed in Table 2.

2.2. Specimen Preparation

A φ5 × 20 sample of the SS-PREP&HIP alloy was used for porosity analysis. The 10 × 10 × 10 mm specimens were halved after heat treatments, ground with abrasive papers, and a Vibromet 2 vibratory polisher (Buehler, Lake Bluff, IL, USA) was used for final polishing. The prepared samples were subsequently utilized for electron backscatter diffraction (EBSD) and scanning electron microscope (SEM) analyses. The samples for transmission electron microscopy (TEM) analysis underwent initial slicing to obtain 400 μm thick slices and thinned to 60 μm with sandpaper. Then, they were thinned using a PIPS II 695 precision ion-beam thinning device (Gatan, Pleasanton, CA, USA) at 3–5 keV. For the specimens after heat treatments, they were cut from the middle of the 10 × 10 × 10 mm samples, while for the specimens after creep testing without rupture, they were cut from the middle cross-section of the gauge. Cylindrical specimens with 2 mm machining allowance for mechanical performance analysis were subjected to heat treatment. The samples after heat treatments were machined into samples with a gauge diameter of 5 mm, an original gauge length of 25 mm, and a total length of 71 mm, for tension testing at room and high temperatures. The samples after heat treatments were machined into samples with a gauge diameter of 5 mm, an original gauge length of 25 mm, and a total length of 75 mm, for creep testing at elevated temperatures. The number of tests for mechanical properties is shown in

Table 3. The number of tests for mechanical properties.

Testing Item	Testing Number	Data Selection
Tensile tests at different temperatures	3	Plotting with the middle data of fracture strain
Creep tests at 750 °C/250 MPa	2	The same results and choose one for plotting

.

2.3. Characterization Methods

The porosity of the SS-PREP&HIP alloy was determined by micro-computed tomography (Micro–CT, GE V|Tome x m, Boston, MA, USA) with an accelerating voltage of 200 kV and current of 100 μA, and 1800 projection results with a resolution ratio of 14 μm were obtained. Phase composition within the 20° to 90° 2θ range was analyzed using a D8 X-ray diffractometer (XRD, Bruker, Billerica, MA, USA) with Co-Kα radiation at a scan rate of 2.5 °/min and room temperature. The microstructure in backscattered electron (BSE) mode, element distribution, and crystallographic features were examined using a Sigma 300 SEM (ZEISS, Oberkochen, Germany) equipped with an energy-dispersive X-ray spectrometer (EDS) and an EBSD system at 20 kV. The scanning step of EBSD was 0.5 μm. Further microstructure analysis was performed via Thermo Fisher Talos F200X TEM (FEI, Hillsboro, OR, USA) at 200 kV. The tensile tests at room temperature, 650 °C, and 750 °C were conducted using a CMT5105 electron universal testing machine (Sansi Yongheng, Ningbo, China) at a strain rate of $1\times {10}^{-4}$ s⁻¹ with an extensometer according to GB/T 228.1-2021 [24] and GB/T 228.2-2015 [25]. The creep performance test was conducted using a QBJ-30 electronic high temperature creep testing machine (Qianbang, Changchun, China ) with an extensometer according to GB/T 2039-2012 [26]. The creep testing was carried out at 750 °C and 250 MPa.

3. Results

3.1. Microstructures of the SS-PREP&HIP Alloy

The density of the SS-PREP&HIP alloy is 4.3 g/cm³, measured by Archimedean method. And there is no porosity detected by Micro-CT for the SS-PREP&HIP alloy, as illustrated in Figure 3. Figure 4 shows the XRD diffractograms of both the powders and alloy after HIP. The powders primarily consist of the α₂ phase, while the γ phase becomes dominant after HIP at 1260 °C.

Figure 5a,b shows the BSE-SEM microstructures of the SS-PREP&HIP alloy, manifesting a near-γ structure with localized duplex phases. The prior particle boundaries can be seen in the microstructure, illustrated in Figure 5a. The different grayscales in Figure 5b, such as sites 2 and 3, manifest the γ phase with different orientations, which show similar chemical composition, as demonstrated in

Table 4. EDS results of marked area in Figure 5b (wt. %).

Site	1	2	3
Al	25	31	31
Ti	61	55	55
Nb	6.5	7.5	7.5
Ta	7.5	6.5	6.5

. The alloy consists of 94% γ and 6% α₂ phases, as shown in Figure 5c, and γ twins can be seen in EBSD IPF of Figure 5d. The average grain size, about 5 μm by EBSD analysis, reveals that the SS-PREP&HIP structure is significantly finer compared with the cast alloy. The EDS results of sites 1, 2, 3 (Table 4) and the mapping of the whole area (Figure 5e–h) in Figure 5b show the γ (Al-rich) phase and α₂ (Al-poor) phase. Ta prefers to exist in α₂-Ti₃Al, and Nb shows no obvious preference, as illustrated in Figure 5g,h.

3.2. Microstructures of the SS-PREP&HIP Alloy After Heat Treatments

The SS-PREP&HIP alloy with different microstructures obtained after different heat treatments according to Table 2, is shown in Figure 6. The alloy characterizes fully coarse lamellar (FL) structures after heat treatments according to T1 and T2, as shown in Figure 6a,b, with the average lamellar colony sizes of 506 μm and 464 μm, respectively. When the alloy is heat treated according to T3, the nearly lamellar (NL) is achieved, characterized by crisscrossed lamellar colonies and metastable structures such as Widmanstätten, feather, and massive γ in the local area. However, the boundaries are not sharp, and there are thick lamellae after precipitation from γ, and part of γ is equiaxial at the junctions. When the heat treatment follows T4, the proportion of lamellae increases, with narrower lamellae and sharper boundaries. The average lamellar colony size is about 269 μm. When the heat treatment follows T5, the transformation of γ lamellae precipitated from α is more complete, the lamellar structures are sharper, and the lamellae are slightly coarsened. As the γ lamellae further precipitated from the initial α₂, α₂ is cut into more lamellae, and the average lamellar colony size decreases to 165 μm. When the heat treatment follows T6, which adds stress relief annealing based on T5, there is no significant variation in the microstructure characteristics compared with T5, and the average lamellar colony size is about 145 μm. When the heat treatment follows T7, the typical thick lath formed by the α phase precipitates from the γ phase, and the residual approximately triangular γ phase can be seen. When the heat treatment follows T8, the convolved NL structure is obtained, that is, the intermittent α₂ and γ lamellae with a specific rotation and inter-arrangement. The average lamellar colony size is about 37 μm, with some equiaxial γ structures also present.

Figure 7 presents the TEM analysis of the sample following the T6 heat treatment. In Figure 5a, the bright field (BF) image is shown, with typical sites marked. This image reveals that the matrix is composed of the γ phase, which contains numerous nanotwins, dislocations, and stacking faults. Stacking faults are exhibited at site 1 in Figure 7a, and this can be verified by Selected Area Electron Diffraction (SAED) patterns in Figure 7b, which correspond to a single γ phase. Similarly, the BF images in Figure 7g,h also clearly demonstrate the significant presence of dislocations and stacking faults, especially stacking faults entangled by dislocations. Nanotwins are displayed at site 2 and site 3 in Figure 7a, and this is confirmed by the SAED patterns in Figure 7c,d. The atomic arrangements in the high-resolution results presented in Figure 7e,f also indicate twin features.

XRD-based phase analysis of the heat-treated alloy indicates that, as depicted in Figure 8, the α₂ + γ two-phase structure resulting from different heat treatment processes is predominantly composed of the γ phase. As indicated by the patterns excluding T6, the planes of (002) γ and (200) γ emerge following high-temperature heat treatment and air cooling. Conversely, as demonstrated by the pattern of T6, these planes vanish after stress relief annealing. When compared with T5, the proportion of γ phase in T6 increases further after stress relief annealing. This is attributed to the fact that the annealing temperature of 850 °C falls within the γ single-phase field. In contrast to T1, the α→γ transformation in T2 is more sufficient after the holding time in the two-phase region is relatively prolonged. When held at lower temperatures within the α single-phase region, compared to T1, T5 yields a greater amount of γ phase.

The crystal characteristics of the heat-treated alloy were analyzed by EBSD. The IPFs, as shown in Figure 9, show that there are orientation differences among different laminates, which is consistent with the contrast difference in laminates in BSE-SEM images of Figure 6. In addition, the presence of twins can also be observed in IPFs. CSL data analyzed by EBSD is shown in

Table 5. CSL data of different heat-treated samples compared with as-HIP material.

Billets No.	As-HIP	T1	T2	T6	T8
Σ3	8.36	80.6	6.55	5.86	8.20
Σ5	0.07	0.00	0.01	0.14	0.04
Σ11	0.00	0.00	0.00	0.00	0.01

, manifesting that Σ3 twins exist in the materials with different heat treatments. The preferred orientation characteristics of the alloy can be quantitatively verified by calculating the texture index (TI) through orientation distribution function (ODF), as shown in

Table 6. Texture index of different heat-treated samples compared with as-HIP material.

Billets No.	As-HIP	T1	T2	T6	T8
TI	3.96	21.22	11.75	10.42	6.82

, and smaller TI values imply higher crystal orientation randomness.

3.3. Tensile Properties of the SS-PREP&HIP Alloy After Heat Treatments

In order to analyze the effects of heat treatments on mechanical performance, the samples with typical structures, including coarse FL structure (T1), relatively fine NL structure without/with stress relief annealing (T5, T6), and fine convolved NL structure (T8), were subjected to tensile testing at room temperature, as illustrated in Figure 10 and

Table 7. Results of tensile test at room temperature of the alloy after different heat treatments compared with as-HIP material.

Sample	As-HIP	T1	T5	T6	T8
Rm/MPa	${357}_{-2}^{+3}$	${325}_{-2}^{+2}$	${433}_{-4}^{+1}$	${407}_{-2}^{+0}$	${397}_{-3}^{+2}$
A/%	${0.3}_{-0.1}^{+0}$	${0.5}_{-0}^{+0}$	${0.9}_{-0}^{+0.1}$	${1.2}_{-0.1}^{+0}$	${0.3}_{-0}^{+0.1}$

, compared with as-HIP material. After the T1 heat treatment, the fracture strain is 0.5%, while after the T5 heat treatment, the fracture strain is increased to 0.9%, and the tensile strength is increased to 433 MPa. After the T6 heat treatment, the tensile strength is 407 MPa, and the fracture strain is 1.2%. The tensile fracture strain at room temperature is only 0.3% after the T8 heat treatment, which is similar to as-HIP material. Figure 11 shows the fracture morphology of room temperature tensile testing. It can be seen that the fractures of the alloys under all heat treatment conditions exhibit a typical river pattern, indicating obvious brittle fracture characteristics, which is consistent with the elongation of the alloy. Additionally, grain boundary cracks were found in the T6 sample.

As the plasticity of TiAl alloy is quite important and the T6 alloy shows attractive room temperature plasticity, the high temperature plasticity of the T6 alloy is further investigated at different temperatures. The tensile results of the alloy after the T6 heat treatment tested at 650 °C and 750 °C are shown in Figure 12 and

Table 8. Results of high-temperature tensile tests of the T6 alloy.

Test Temperature	650 °C	750 °C
Rm/MPa	${419}_{-1}^{+3}$	${460}_{-3}^{+2}$
A/%	${1.2}_{-0.1}^{+0}$	${7.5}_{-0.1}^{+0}$

. The strength is 419 MPa and 460 MPa, respectively, and the fracture strain is 1.2% and 7.5%, respectively. Figure 13 shows the fracture morphology of the tensile tests conducted at high temperatures. The fracture morphology at 650 °C mainly exhibits river patterns, whereas the fracture morphology at 750 °C shows a large number of dimples, corresponding to the improved ductility of the alloy at this temperature.

3.4. Creep Performance of the SS-PREP&HIP Alloy After Heat Treatments

The creep testing result of the alloy after the T6 heat treatment at 750 °C/250 MPa is shown in Figure 14. The test was stopped without rupture when creeping for 225 h, with strain of 11%. The creep process involves three stages: the initial, steady state, and accelerated creep stage. In the initial creep stage, the creep rate is initially high but gradually decreases over time until it reaches a steady value. In the steady-state creep stage, the creep rate remains nearly constant, and the creep strain is approximately linear with time, with its slope representing the steady-state creep rate. For the T6 alloy, the steady-state creep rate is 8.92 × 10⁻⁶. The third stage is the accelerated creep stage, where the creep rate increases rapidly.

Figure 15 shows the unruptured section morphology as well as its EDS mapping results, where obvious holes can be seen. From Figure 15a, it can be found that most of the holes are located at the edges of the lamellae. Ta is mainly enriched at the edge of the lamellae, as shown in Figure 15a,b, while there is no obvious preference for Nb, as shown in Figure 15e. TEM analysis of the microstructure after creep testing is shown in Figure 16, revealing the presence of a large number of nanotwins.

4. Discussion

4.1. Effects of Heat Treatments on Microstructures

The process of SS-PREP^® powder preparation involves rapid solidification and insufficient diffusion [23]. Due to the partial phase transition of α → α₂ + γ under rapid cooling, the powders primarily consist of the α₂ phase. When subjected to hot powder isostatic pressing at 1260 °C, the metastable α₂ phase almost completely transforms into the γ phase, as illustrated in Figure 4. The oxide scale on the surface of particles impedes the metallurgical bonding and the element diffusion [27], resulting in the prior particle boundaries, as shown in Figure 5a. Table 6 shows that the TI value of as-HIP material is just 3.96, suggesting the random orientation. The HIP alloy powder has isotropic properties and no preferred orientation since the powders that are loaded into the envelope undergo isotropic pressure and heat throughout the HIP process.

γ-TiAl alloys can exhibit four typical microstructures [5]: FL, NL, duplex, and near γ, forming under different heat treatment processes. Generally, the near-γ microstructure has coarse grains and poor performance at both room and high temperatures, while the NL structure has poor plasticity at room temperature. In contrast, FL and duplex structures have the ideal combination of ductility and toughness at room temperature.

When the alloy is heat treated according to T1 and T2, the alloy almost entirely transforms into the α phase during 1420 °C holding in α single-phase field, and the γ phase precipitates from the α phase to form the FL structure in the subsequent cooling process. Compared to as-HIP material, the TI values rise following various heat treatments, especially T1, primarily because of the identified area having almost just one coarse lamellar colony. T1 exhibits a relatively high degree of supercooling and a significant phase transformation driving force. This promotes the rapid nucleation and growth of the γ phase through twinning, thereby producing a large number of Σ3 twinned boundaries, as illustrated in Table 5. When annealing at a lower temperature above T_α for T3 heat treatment, the NL structure is obtained after air cooling, and γ phase is distributed between γ/α₂ lamellae. The fraction of lamellae grows significantly, the lamellar structures sharpen, and the lamellae are slightly coarsened with adding furnace cooling before air cooling, as reported by T4 and T5. According to T6, adding stress relieving annealing does not result in any appreciable changes. For peritectic TiAl alloys, after treating in the γ single phase region (during the HIP process) to obtain near γ structure, and then treating in α + γ two-phase region during the cooling process, the Widmanstätten α₂ lath will be precipitated along four groups close-packed planes of γ phase, as the microstructure after T7 heat treatment.

The α-dominated convolved NL structure is obtained through heat holding in α single-phase region when the alloy is heat treated according to T8. Adding a small amount of Ta can facilitate the formation of metastable structures by relaxing the cooling rate requirements for metastable transition [28]. The interfaces between γ/metastable γ are low-energy interfaces, including Σ3 (massive γ), Σ5 (feather), and Σ11 (Widmanstätten) CSL interfaces [29]. Ta atoms can enhance the nucleation and twinning of metastable structures by reducing the interface energy [28] and the phase transition resistance by substituting Al atoms on or near the interface, such as Σ3 interfaces, 120° rotation interfaces, and anti-phase domain boundaries. Due to the slow diffusion rate of the Ta element, the α→γ_m massive transition can be completed at a slower cooling rate, and the massive structure can be obtained from air cooling in the α single phase [19]. When the metastable γ_m ages in the α₂ + γ two-phase region, the newly precipitated α₂ maintains a strict orientation relationship with γ: {111}γ ∥ {0001}α₂, <110>γ ∥ <11$\overline{2}$0>α₂. α₂ can nucleate on four {111}γ planes at the same time, and the massive transition leads to the formation of a large number of defects in γ_m, such as anti-phase domain boundaries, stack faults, twins, and dislocation steps [22], which provide a large number of sites for α₂ nucleation. Therefore, α₂ can nucleate in many places and orientations, eventually forming a fine structure composed of convolved α₂ + γ lamellae.

According to previous studies [13,14], the fine fully lamellar structure can be achieved through cyclic heat treatments of cast TiAl-Ta alloys, taking into account the massive transformation characteristics of Ta-containing TiAl alloys from the α single-phase region during air cooling. Fortunately, powder metallurgy TiAl alloys exhibit a relatively uniform composition and fine microstructure in their initial state. Therefore, by brief heat holding in α single-phase region, to obtain an α single-phase dominant microstructure, γ lamellae and metastable structures can precipitate from α phases during slow cooling to the α + γ two-phase region. These structures have different crystal orientations compared to the parent phases [20], resulting in refined lamellae and metastable structures during the cooling process. Subsequently, air cooling until ambient temperature can prevent lamellar coarsening and grain growth. The γ lamellar growth is governed by long-range diffusions, while massive γ structures are controlled by short-range diffusions [30]. Since lamellar growth relies on diffusions [31], it is expected that Ta, being a slow diffuser, will decelerate the growth of lamellae during air cooling [32]. Recovery and recrystallization occur during stress relief annealing at 850 °C, stimulated by internal stress and stored energy of metastable structures, thereby stabilizing metastable structures with the elimination of internal stress and defects [14]. With the extension of annealing duration, the grain size is basically unchanged, and low ΣCSL GBs, particularly the coherent Σ3 grain boundaries, increase further. The coherent Σ3 boundaries can create the dislocation slip path, improve the dislocation slip, and uniform the plastic deformation, thereby contributing to plastic deformation [33,34]. As a result, the dislocations easily slide along the Σ3 boundaries, avoiding the hardening resulting from the dislocation accumulation [35]. Finally, the NL Ti-48Al-3Nb-1.5Ta alloy with a relatively fine microstructure was obtained by a simple T6 heat treatment, which met the application requirements.

4.2. Effects of Heat Treatments on Tensile Performance

Although the grain size is small after the T8 heat treatment, the structure is a convolved lamella + near-γ structure, and the tensile fracture strain at room temperature is as low as the as-HIP material. Similarly, the fracture strain with a coarse FL structure after T1 heat treatment does not show significant improvement. In contrast, the tensile property is improved significantly at room temperature after T5 heat treatment, which is a NL structure with a moderate grain size. After stress relief annealing based on T5, the density of dislocation and other defects are reduced due to the recovery, the tensile strength is reduced by 26 MPa, and the fracture strain is increased to 1.2%, which is better than Ti46Al8Nb alloy and Ti46Al8Ta alloy [11]. Nanotwins, as shown in Figure 7, contribute to improved mechanical properties by hindering dislocation movement [36]. Even so, the fractures with different heat treatments are all river patterns, as shown in Figure 11. Combined with the microstructures and grain sizes after different heat treatments in Chapter 3.2, it can be seen that the mechanical properties are affected not only by the size of the microstructure but also by the microstructure type [1,4].

Compared with as-HIP material, the remarkable increases in strength and elongation of heat-treated alloys are mainly due to the refinement and morphological regulation. The refined microstructure is obtained through the metastable structure evolution according to the designed heat treatment, improving plastic incompatibilities and yielding uniformity during the plastic deformation [37]. In addition, more interfaces can hinder the dislocation movements, thus increasing strength [38]. In contrast, for the coarse microstructure, stress is more likely to concentrate in certain places, such as triple junctions, resulting in local strain and crack initiation. In thick lamellae, once cracks occur, they rapidly extend through coarse colonies or along boundaries, leading to macro scale failures [39].

Compared with the cast Ti-48Al-3Nb-1.5Ta alloy, whose fracture strain reaches 1.7% by five-cycle heat treatment [13] or 1.0% by multi-step heat treatments [14], the SS-PREP&HIP Ti-48Al-3Nb-1.5Ta alloy can achieve similar plasticity at room temperature by a simple heat treatment, meeting the service requirements. The convenient and controllable heat treatment process is more conducive to engineering applications.

In high temperature tensile testing, the T6 alloy exhibits the highest strength and plasticity at 750 °C. Compared with different TiAl alloys, as illustrated in Figure 17, the T6 alloy demonstrates good strength and outstanding plasticity at 750 °C, whose fracture strain is better than that of additive manufacturing TiAl4822 [40], additive manufacturing TiAl4822-8Ta [9], directionally solidification Ti-45Al-7Nb, Ti-45Al-8Nb [41], and Ti-47Al-2Cr-2Nb [42] alloys. As illustrated in Figure 13, the fracture morphology at 750 °C tensile exhibits dimples, but the fracture morphology at 650 °C is primarily cleavage, indicating that the tough-brittle transition temperature is between 650 and 750 °C. In the tensile test conducted at 650 °C, the fracture occurs on the lamellar plane, as depicted by the fracture morphologies presented in Figure 13a,b. The slip deformation encounters significant difficulty in traversing the lamellae (owing to the hard orientation). During the tensile test at 750 °C, with the transition from brittle fracture to tough fracture, the dislocation climbing in the γ phase and α₂ phase introduces an additional deformation mechanism. This mechanism effectively alleviates the plastic incoordination between the two phases. Correspondingly, the fracture surface exhibits a substantial number of dimples, as demonstrated in Figure 13c,d.

4.3. Creep Behavior Analysis

The primary deformation mechanisms of TiAl alloys during creep are dislocation slip and deformation twinning. As a result of external forces, dislocations originate from the lamellar interfaces and proceed toward the lamellar borders during the creep process after being initially triggered in the α₂/γ lamellae. Due to the lack of deformation coordination between neighboring lamellar colonies, dislocations accumulate at the edges of the lamellar colonies [43]. When the dislocations become entangled or aggregated to a certain extent, the internal stress can be relieved by twinning.

Deformation twins alleviate stresses during creep and improve the total strain [44]. Refractory elements Ta, Nb, Mo, and W can increase the creep resistance and decrease the creep rate of TiAl alloy by strengthening the solution and decreasing diffusion rates. For Ti-43Al-9V-Y alloy [45], the creep rupture life of a duplex structure at 700 °C/250 MPa is 140.25 h with steady-state creep rate of 1.51 × 10⁻⁷, and the creep rupture life of a FL structure is 182 h with steady-state creep rate of 1.22 × 10⁻⁷. Despite a somewhat lower steady-state creep rate, the creep rupture life of both structures is lower than that of the same stress at 750 °C in this study.

The surface damage caused by oxidation and the volume defect caused by cavity nucleation will increase the creep rate [43], and the creep rate will increase with the extension of time until the specimen ruptures [46]. The fracture of the specimen occurs in the final stage of creep caused by overloading, and the crack propagation path extends preferentially to intergranular and transgranular. Numerous cavities in the gauge section indicate a vigorous coalescence process during creeping, especially in the third stage. This cavitation mechanism progressively reduces the effective cross-section of the creep specimen, potentially causing local overloading that leads to fracture [12].

Sneary and Tian [47,48] found that Ti-44Al-8Nb-0.2W-0.2B-0.1Y (at. %) alloy and Ti-47Al-2Cr-2Nb (at. %) alloy preferentially nucleate along GBs inclined 45° to the stress axis under the dynamic recrystallization, high shear stress, and deformation twins. Du [49] showed that the cavity growth along the lamellar boundaries is governed by diffusions in Ti-46.5Al-2Cr-3Nb-0.2 W (at. %) alloy. As shown in Figure 15, Ta is mainly enriched at the edge of the lamellae, which is similar to the distribution of τ phase rich in Ta for the high-Ta TiAl alloy [12]. Due to low diffusion rate of Ta, the expansion of holes is hindered. In contrast, the distribution tendency of Nb is not obvious, and the effect of Ta on improving the creep resistance of TiAl alloys is more notable. Adding alloying elements and adjusting the grain boundary characteristics can lower the grain boundary diffusion, and the pore growth can be inhibited by adding Ta and Nb to reduce the diffusion coefficient.

5. Conclusions

A Ti-48Al-3Nb-1.5Ta alloy was prepared by SS-PREP&HIP and heat treated through different regimes. The microstructure and mechanical performance of the SS-PREP&HIP Ti-48Al-3Nb-1.5Ta alloy after heat treatments were analyzed and elucidated. The important findings are as follows:

(1)

The NL structure, with the lamellar colony size of about 145 μm for this PM alloy, was first obtained by simple heat treatment, with aa tensile fracture strain of 1.2% at room temperature and 7.5% at 750 °C, superior to other TiAl alloys. Furthermore, the simple regime is conducive to large-scale application.

(2)

The designed alloy shows excellent creep performance, with a rupture life at 750 °C/250 MPa that is superior to Ti-43Al-9V-Y alloys, due to the presence of Ta.

(3)

The Ta element gathers at the lamellar edge during creep process, which can inhibit lamellar growth and hinder pore expansion, thus improving creep performance.