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High-temperature tensile properties and deformation fracture mechanism of a TiAl-Nb alloy

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Published/Copyright: April 14, 2026
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High Temperature Materials and Processes
From the journal High Temperature Materials and Processes Volume 45 Issue 1

Abstract

A high-performance TiAl-Nb alloy was obtained by forging and heat treatment. The tensile properties of the alloy were tested, and the microstructure of the alloy was observed via transmission electron microscopy (TEM) and scanning electron microscopy (SEM) with electron backscatter diffraction (EBSD). The results revealed that after the alloy undergoes tensile fracture at the temperature of 820 °C, the yield strength was 420 Mpa and the elongation rate is 3.6 %. After tensile fracture at 880 °C, the yield strength of the alloy was 340 Mpa and the elongation rate is 9.8 %.The brittle-ductile transition temperature (BDTT) of the alloy ranged from 820 °C to 880 °C. In the brittle fracture process, the alloy fractures by a mixed mode of lamellar fracture and lamellar colony boundary fracture. In lamellar colonies, cracks propagate mainly along the γ/γ interface. Among the lamellar colony boundaries, the passivated serrated lamellar colony boundary has better crack propagation resistance than the straight strip lamellar colony boundary. In the process of plastic fracture, lamellar colony boundary fracture is the main damage mechanism. The repeated recrystallization of γ grains and abundant B2 phase that precipitates within and between lamellar colonies are the main reasons for the formation and propagation of cracks.

Keywords: TiAl-Nb alloy; high-temperature tensile; B2 phase; fracture mechanism

1 Introduction

As a new generation of lightweight, high-temperature structural materials, TiA-Nb alloys have been extensively studied in recent years, and great theoretical breakthroughs have been made [1], [2], [3], [4], [5]. Zheng et al. prepared a Ti-43.5Al-4Nb-1Mo-0.1 B alloy with a three-phase trimodal structure by solution and ageing heat treatment, and the results revealed that the pearlitic-like microstructure content has an important effect on the mechanical properties of the alloy [6]. Dai et al. investigated a γ-Ti-45Al-8.5Nb-(W, B, Y) alloy and reported that a two-step heat treatment improved the tensile properties of the alloy due to microstructure refinement and a reduction in the amount of remaining B2 phase located at lamellar colony boundaries [7].

The β(B2) phase has an important effect on the mechanical properties of TiAl-Nb alloys. Under high-temperature conditions, the β phase is a disordered body-centred cubic structure with comparably more slip systems, which can decrease the strength and increase the deformation ability of the alloy. During the cooling process, the body-centred atoms and vertex atoms in the β phase lattice shift, breaking the structural symmetry of the body-centred cubic lattice and forming the ordered B2 phase. The B2 phase is hard and brittle, which is not conducive to room-temperature plasticity [8], 9]. Masahashi et al. reported that the β phase can precipitate at grain boundaries during thermal deformation and inhibit grain growth so that alloy grains can be refined. In addition, the β phase can promote grain boundary slip, thus improving the plastic deformation ability of the alloy [10]. Liu et al. studied the thermal deformation behaviour of a Ti-47Al-7Nb-0.4W-0.15 B alloy and concluded that the β phase can enhance the slip and migration ability of grain boundaries and that the β phase at the grain boundaries has priority in deformation and recrystallization, thus significantly improving the thermal deformation ability of the alloy. In addition, during the deformation process, the β phase can decompose into the α 2 and γ phases, thus alleviating stress concentrations and making the alloy less prone to cracking during thermal deformation [11]. Liang et al. investigated the tensile creep behaviour of a heat-treated β-solidified γ-TiAl alloy and reported that the B2 phase can be involved in deformation under specific creep conditions and that voids can nucleate at the γ-B2 and B2-colony interfaces [12].

In this work, a TiAl-Nb alloy with a nominal composition of Ti-44Al-8Nb-0.2W-0.2B-0.1Ywas treated by forging and heat treatment, and a microstructure with a uniform grain size and regular lamellar boundaries was obtained. Because Nb is a β-stabilized element, the alloy studied in this work has more β phases and has a complex microstructure transition during the tensile process [9], 11], [13], [14], [15], [16]. The TiAl-Nb alloy demonstrates a significant improvement in high temperature strength and the upper limit of service temperature compared to conventional TiAl alloys. The addition of Nb and W elements promoted the formation of the β phase, not only enhancing the hot working properties of the alloy, but also simultaneously improving its high-temperature strength and oxidation resistance. In addition, the trace addition of elements B and Y can effectively refine the crystal grains and further improve the hot formability of the alloy. Previous studies generally suggest that in the intergranular fracture process of TiAl-Nb alloys, cracking mainly occurs along the γ/α 2 interface. However, the microstructure analysis in this study indicates that the fracture actually occurred at the γ/γ interface. This phenomenon reveals a new fracture mechanism, providing a new research direction for a deeper understanding of the deformation and fracture behavior of this type of alloy and for conducting interface optimization design.

2 Experimental materials and methods

A high-purity master alloy was prepared as an alloy casting blank with a nominal composition of Ti-44Al-8Nb-0.2W-0.2B-0.1Y (atomic percentage, at.%) in a vacuum shell furnace. The alloy casting billet was isothermal forging at 1,250 °C and a strain rate of 10−3/s. The forged alloy was subsequently heat treated in a furnace at 1,380 °C for 2 h and allowed to cool. After heat treatment, the samples were cut into I-shaped plates with cross-sectional dimensions of 4.5 mm × 2.5 mm and a standard length of 30 mm. After mechanical grinding and polishing, the tensile properties of the alloy were tested at different temperatures and a rate of 1 × 10−3/s via a WDW-100 electronic universal testing machine, the experimental process is shown in Figure 1. A Talos F200X transmission electron microscopy (TEM) system and a Thermo S field emission electron microscope with electron backscatter diffraction (EBSD) were used to observe the microstructures of the original sample and the sample after tensile fracture.

Figure 1: The experimental process: (a) original forged alloy (a) tensile test sample (c) mount the test sample (d) tensile the sample.
Figure 1:

The experimental process: (a) original forged alloy (a) tensile test sample (c) mount the test sample (d) tensile the sample.

3 Experimental results and analysis

3.1 Microstructure of the alloy

The morphology of the alloy determined via a backscattered electron (BSE) detector is shown in Figure 2. Figure 2(a) shows the microstructure of the alloy under low magnification. The alloy has a fully lamellar structure, and the lamellar colonies are relatively uniform in size. Ten images were selected and analyzed using Image J, resulting in an average size of the lamellar colonies of approximately 163 μm. Bright white precipitates exist within and on the boundaries of the lamellar colonies of the alloy, as shown by the white and black arrows in Figure 2(a), respectively. Since the average atomic number of the B2 phase is significantly higher than that of the γ phase and the α 2 phase, according to the BSE imaging principle, the B2 phase region will generate a stronger backscattered electron signal, which appears as a bright white color. Therefore, it can be determined that this phase in the figure is the B2 phase. Figure 2(b) is a higher magnification view of the local area shown in Figure 2(a). The B2 phase usually occurs in the form of strips or blocks at the interface between the α 2/γ lamellar colonies, which is known as β segregation. Alternatively, it can also be distributed in the form of B2/γ lamellae within the α 2/γ lamellar colonies, known as α segregation. The bright white regions shown in the figures are the α segregation and β segregation regions, as shown by the white and black arrows, respectively. Similar morphologies have been reported in other studies [17], 18].

Figure 2: Microstructure of the alloy: (a) Low-magnification morphology and (b) high-magnification morphology.
Figure 2:

Microstructure of the alloy: (a) Low-magnification morphology and (b) high-magnification morphology.

Figure 3 shows the EBSD analysis of the alloy near the lamellar colony boundaries, and Figure 3(a) and (b) show the IQ map and phase map, respectively. The volume fractions of the γ phase, α 2 phase, and B2 phase in this region are 90.4 %, 8.8 %, and 0.8 %, respectively. The alloy has two kinds of lamellar colony boundaries. The three lamellar colonies with different orientations are defined as A, B, and C. The lamellar boundary between lamellae A and C is a straight strip, whereas the lamellar colony boundary between lamellae A and B is serrated.

Figure 3: EBSD observation results for the alloy near the lamellar colony boundaries: (a) IQ map and (b) phase map.
Figure 3:

EBSD observation results for the alloy near the lamellar colony boundaries: (a) IQ map and (b) phase map.

Figure 4 shows the TEM bright field image of the alloy. Figure 4(a) shows a TEM image of the lamellar region of the alloy. The structure is composed of alternating lamellae with flat phase boundaries and few dislocations within the lamellae. Figure 4(b) shows another TEM image of the lamellar region of the alloy. The equiaxed grains that precipitate in the lamellar structure are small, ranging from a single γ grain, as shown by the black arrow, or grains subjected to α segregation, as shown by the white arrow in Figure 4(b).

Figure 4: TEM microstructure of the alloy: (a) Lamellar region and (b) precipitated phase.
Figure 4:

TEM microstructure of the alloy: (a) Lamellar region and (b) precipitated phase.

3.2 Tensile properties of the alloy

Figure 5 shows the tensile curves obtained by tensile fracture of the alloy at different temperatures. Figure 5(a) shows the tensile fracture curves of the alloy at 800 °C, 820 °C, 880 °C, and 920 °C, the tensile properties data of the alloy are shown in Table 1, which indicate that the alloy enters the plastic zone above 820 °C. Under lower temperature conditions, the yield strength increases sharply with increasing time, and the sample breaks quickly; under higher temperature conditions, the strength of the alloy increases rapidly with increasing time before gradually decreasing after the peak value is reached. Figure 5(b) shows the yield strength and elongation of the alloy at different temperatures. The brittle-plastic transition temperature of the alloy ranges from 820 °C to 880 °C, and the phenomenon of brittle-ductile transition has been reported in many studies [19], [20], [21], [22], [23], [24].

Figure 5: Tensile properties of the alloy: (a) Stress as a function of elongation and (b) stress and elongation as a function of temperature.
Figure 5:

Tensile properties of the alloy: (a) Stress as a function of elongation and (b) stress and elongation as a function of temperature.

Table 1:

The tensile properties data of the alloy.

Parameters temperature 800 °C 820 °C 880 °C 920 °C
Tensile strength (MPa) 461 420 340 325
Elongation (%) 4.3 4.7 9.8 11.2

3.3 Deformation and fracture characteristics of the alloy in the brittle stage

The deformation and fracture characteristics of the alloy in the brittle stage are shown in Figure 6–9. The SEM morphology of the alloy sample after tensile fracture at 800 °C is shown in Figure 6, and the direction of the applied stress is marked by the white double-headed arrows in Figure 6(a). During the brittle stage, cracks can propagate within lamellar colonies (crack 1 in Figure 6(a)) and at the lamellar colony boundaries (crack 2 and crack 3 in Figure 6(a)). When a crack expands into an adjacent lamellar colony, if the lamellar orientation difference from the adjacent lamellar colony is large, crack growth is blocked and stops, as shown in the white box in Figure 6(a). The crack can extend from the lamellar colony boundary into the lamellar colony and terminate within the lamellar colony, as shown in crack 2 in Figure 6(a). If the tensile stress continues to increase, cracks overcome the resistance and coalesce. The microstructures of the different cracks connected together are shown in Figure 6(b). Crack deflection is an important toughening mechanism of an alloy.

Figure 6: Propagation of cracks in the brittle stage: (a) Lamellar colonies with extreme orientation differences, (b) cracked connection, (c) toughened shear band, and (d) lamellar colonies with less extreme orientation differences.
Figure 6:

Propagation of cracks in the brittle stage: (a) Lamellar colonies with extreme orientation differences, (b) cracked connection, (c) toughened shear band, and (d) lamellar colonies with less extreme orientation differences.

Figure 7: EBSD results of the lamellar colony area of the alloy after tensile fracture at 800 °C: (a) IQ map and straight strip boundary, (b) phase map, and (c) IPF map.
Figure 7:

EBSD results of the lamellar colony area of the alloy after tensile fracture at 800 °C: (a) IQ map and straight strip boundary, (b) phase map, and (c) IPF map.

Figure 8: EBSD observations of lamellar colony boundaries near the fracture after tensile fracture at 800 °C. (a)-(b) crack propagation at the lamellar colony boundary and (c)-(d) passivated serrated lamellar colony boundary.
Figure 8:

EBSD observations of lamellar colony boundaries near the fracture after tensile fracture at 800 °C. (a)-(b) crack propagation at the lamellar colony boundary and (c)-(d) passivated serrated lamellar colony boundary.

Figure 9: TEM morphology of the alloy after tensile fracture at 820 °C: (a) Straight strip boundary, (b) passivated serrated boundary, (c) steps in the lamellar area, and (d) twinning in the lamellar region.
Figure 9:

TEM morphology of the alloy after tensile fracture at 820 °C: (a) Straight strip boundary, (b) passivated serrated boundary, (c) steps in the lamellar area, and (d) twinning in the lamellar region.

A crack can also extend from within the lamellae and terminate at the lamellae boundary, as shown by the crack in Figure 6(c). When two closely spaced cracks in the lamellae grow to meet, they can join together by cutting and tearing the ligament, as shown by the white arrow in Figure 6(c). When the lamellar orientation difference between adjacent lamellar colonies is small, cracks can propagate along the same crack plane with less propagation resistance, resulting in a linear crack propagation path, as shown in Figure 6(d).

To further analyse crack propagation in different regions, the microstructures of lamellar colonies and lamellar colony boundaries were observed. Figure 7 shows the EBSD analysis of crack propagation in lamellar colonies. Figure 7(a)–(c) shows the IQ map, phase map and IPF map of the lamellae region after tensile fracture at 800 °C. Many studies have suggested that cracks easily occur at the γ/α 2 interface due to incompatibility, which is inconsistent with the results observed in this work [25], [26], [27]. Cracks 1 and 2 propagate along the γ/γ interface. At the junction of cracks 1 and 2, the γ/γ and γ/α 2 interfaces are broken, as shown by the yellow arrows in Figure 7(a)–(c), which is similar to the results observed in Figure 7(c). Large holes can form in the lamellar colony without spreading or expanding, as shown by the white circles in Figure 7(a)–(c).

Figure 8 shows the EBSD analysis of the lamellar colony boundaries near the fracture zone after tensile fracture at 800 °C. Figure 8(a)–(b) shows the IQ map and IPF map of the alloy. The four lamellar colonies in the figure are denoted lamellar colonies A, B, C, and D. To distinguish between different types of lamellar colony boundaries, the left and right parts of lamellar colony D are defined as lamellar colony D1 and lamellar colony D2, respectively. As mentioned in Section 3.1, the lamellar colony boundaries of the alloy can be divided into two types: straight strips and passivated serrations. The boundary between lamellar colonies A and D1 is a straight strip boundary, whereas the boundary between lamellar colonies A and D2 is a passivated serrated boundary. Among the three cracks, the crack widths decrease in the order of crack 1, crack 2, and crack 3; this shows that cracks propagate from left to right and that the crack propagation resistance in lamellar colonies is the lowest, followed by that of straight strip lamellar colony boundaries and, finally, that of passivated serrated lamellar colony boundaries. The crack terminates at the passivated serrated lamellar colony boundary between lamellar colonies A and D2, with the same microstructure observed in Figure 8 (c). Figure 8(c)–(d) shows the IPF map and phase map of a passivated serrated lamellar colony boundary at high magnification. Holes were generated at the lamellar colony boundary but did not expand. In addition, microcracks almost perpendicular to the lamellar colony boundary pass through the hole and end at the lamellar boundary, as shown by the white arrow in the image. This finding indicates that the passivated serrated lamellar colony boundary has high strength and can inhibit crack propagation.

The TEM morphology of the alloy after tensile fracture in the brittle stage (820 °C) is shown in Figure 9. Figure 9(a) shows the TEM bright field image of the interface between the lamellar colony and lamellar colony boundary near the fracture after tensile fracture. The lamellar colony boundary in Figure 9(a) is a straight strip. The dislocation density is low both in the lamellar colony and at its boundary, and no dislocation plugging is found at the interface between the lamellar colony boundaries. The lamellar colony boundary in Figure 9(b) is passivated by the serrated lamellar colony boundary, and dislocation plugging occurs at the interface between the lamellar colony and lamellar colony boundary, as shown by the black arrow in Figure 9(b). This is the main reason why the passivated serrated lamellar boundary has greater strength than the straight strip lamellar boundary. As shown in Figure 9(c), there are dislocations and interface steps in the γ lamella, as shown by the black and white arrows in Figure 9(c), respectively. The dislocation density in the γ lamella is low. Figure 9(d) shows the TEM bright-field image of the other lamellae region under these conditions. There are relatively high-density twins in the γ lamellae, and the twin propagation direction is perpendicular to the phase boundary, which provides the alloy with shear stress in the vertical lamellar direction. The above morphologies reveal that the main deformation mechanism of the alloy in the γ lamellae is twin deformation and that interface migration and dislocation slip are secondary deformation mechanisms.

3.4 Deformation and fracture characteristics of alloys at the plastic stage

TiAl alloys have poor plasticity at room temperature and require hot working above the brittle‒ductile transition temperature. Therefore, studying the deformation behaviour of TiAl alloys at the plastic stage can provide a theoretical basis for forging, extrusion, rolling, and other hot working processes.

The BSE morphology of the alloy after tensile fracture at 920 °C is shown in Figure 10. Figure 10(a) shows the morphology under low-magnification conditions. After tensile fracture at 920 °C, a large amount of the B2 phase precipitated into lamellar colonies and lamellar colony boundaries, and many cracks and holes were found in the lamellar colony boundaries. For more detail, a magnified image is shown in Figure 10(b), which reveals that the newly precipitated B2 phase exists in bulk form in the lamellar colonies.

Figure 10: Deformation features of the alloy after fracture at 920 °C. (a) Low-magnification backscatter morphology and (b) high-magnification backscatter morphology.
Figure 10:

Deformation features of the alloy after fracture at 920 °C. (a) Low-magnification backscatter morphology and (b) high-magnification backscatter morphology.

Figure 11 shows the EBSD analysis of the alloy after tensile fracture at 920 °C. Figure 11(a)–(c) shows the IQ map, phase map, and IPF map of the region near the lamellar colony boundary. Both the γ phase and the B2 phase of the lamellar colony boundary recrystallize, and similar observations have been reported in many studies [28], [29], [30], [31], [32]. The γ phase fully recrystallizes with small and uneven grain sizes and different grain orientations. The degree of recrystallization of the B2 phase is relatively low, and the orientation of the recrystallized grains is consistent, as shown by the white arrows in Figure 11(a)–(c). Elongated cracks and holes formed at the interface between the lamellar colony and its boundary, as shown by the black arrow. Figure 11(d) shows the Kernel Average Misorientation (KAM) map, which reveals that the KAM value of large γ grains is low and that the KAM values of small grains and B2 grains are high. The higher the KAM value is, the greater the degree of lattice distortion must be, i.e., the higher the dislocation density. The low KAM value of small γ grains indicates that small γ grains are transformed from large γ grains through continuous recrystallization, which consumes dislocations. The main reason for the low KAM value of the B2 grains is that the B2 phase is softer than the γ and α 2 phases are at high temperatures and is therefore deformed more easily [11].

Figure 11: EBSD results of the alloy near the fracture zone after tensile fracture at 920 °C: (a) IQ map, (b) phase map, (c) IPF map and (d) KAM map.
Figure 11:

EBSD results of the alloy near the fracture zone after tensile fracture at 920 °C: (a) IQ map, (b) phase map, (c) IPF map and (d) KAM map.

EBSD analysis was performed on the lamellar colony regions near the fracture zone after tensile fracture at 920 °C to further analyse the B2 phase deformation mechanism during the high-temperature tensile process. Figure 12(a)-(b) shows the IQ map and phase map of the precipitation in the lamellae. The precipitate in the lamellae is a mixture of the γ phase and B2 phase, namely, a β segregation phase. The B2 phase is a complete grain with no recrystallization, which means that the β segregation phase in lamellae does not deform significantly because of the restriction of surrounding lamellae.

Figure 12: EBSD observations of the precipitated phase in the lamellar colony after tensile fracture at 920 °C: (a) IQ map and (b) phase map.
Figure 12:

EBSD observations of the precipitated phase in the lamellar colony after tensile fracture at 920 °C: (a) IQ map and (b) phase map.

The lamellar colony boundary near the crack was cut by a focused ion beam (FIB), and the cutting position is shown in the black square wire frame in Figure 13(a). Figure 13(b) shows the overall morphology of the cut region. The γ grains at the lamellar colony boundary are equiaxed, and the B2 grains are irregularly nodular. Figure 13(c) shows an enlarged view of area A, where several recrystallized grains precipitated within the original grains. Selected area electron diffraction (SAED) pattern analysis was performed on the material and the original matrix, as shown in the upper right corner and lower right corner of the figure. The orientation of the γ phase was deflected by 45° in this grain. Figure 12(d) shows an enlarged view of area B, wherein the recrystallized grains precipitated at the grain boundaries, as shown by the black arrow. The B2 and γ phases can also arrange in alternating forms, as shown by the white circles in Figure 13(d). Figure 13(e) shows an enlarged view of area C, wherein the B2 phase also recrystallizes to form new grains by discontinuous recrystallization, and thus, the grains are equiaxed. Figure 13(f) shows an enlarged view of area D, which reveals that cracks formed in the γ grains.

Figure 13: Microstructure of the alloy after tensile fracture at 920 °C: (a) Cutting position, (b) overall morphology, (c) local magnification of area A, (d) local magnification of area B, (e) local magnification of area C, and (f) local magnification of area D.
Figure 13:

Microstructure of the alloy after tensile fracture at 920 °C: (a) Cutting position, (b) overall morphology, (c) local magnification of area A, (d) local magnification of area B, (e) local magnification of area C, and (f) local magnification of area D.

4 Discussion

This analysis revealed that the alloy has a fully lamellar structure. During the forging process, due to the significant plastic deformation that occurs, there may be a relatively high density of dislocations in the alloy. However, during the subsequent heat treatment process, due to the higher temperature, the atomic vibrations intensify at high temperatures, and the diffusion rate significantly increases. The dislocations gain sufficient energy through sliding and climbing to overcome the energy barriers and achieve rearrangement or annihilation, thereby resulting in a lower dislocation density in the alloy after heat treatmen, as shown in Figure 4. The uniform size of the lamellar colony boundaries reduces the degree of fracture instability caused by the nonuniform microstructure during the tensile process. The alloy underwent a brittle-ductile transition between 820 °C and 880 °C.

In the brittle stage, owing to the short duration of tensile fracture and the small number of dislocation slip planes, the alloy did not experience obvious dislocation proliferation during the tensile process. Similarly, there is less interface migration. Therefore, the alloy deforms mainly by twinning. The fracture mode of the alloy is a mixed mode of lamellar fracture and lamellar boundary fracture, and the cracks start mainly at the γ/γ interface and propagate along the interface. The propagation resistance of cracks is closely related to the orientation of adjacent lamellar colonies; when the orientation difference between adjacent lamellar colonies is small, the propagation resistance is also low, and the crack extends directly along the lamellar interface or lamellar colony boundary of the adjacent lamellar colony. When the difference in orientation between adjacent lamellar colonies is large, the propagation resistance is high, and the crack deflects and spreads along the lamellar colony boundary. Crack deflection can alleviate the stress concentration at the crack tip and increase the fracture area, the surface energy required for crack propagation, and the fracture toughness of the alloy. Two cracks expanding along different lamella in the same lamellar colony can be connected by shear ligaments. Because the resistance of these ligaments must be overcome during the fracture process of the alloy, plastic deformation occurs at the ligament connection, which consumes the elastic strain energy and improves the fracture toughness of the alloy. This mechanism is called the ligament toughening mechanism [33], [34], [35].

There are two types of lamellar colony boundaries in the alloy. Compared with straight strip lamellar colony boundaries, passivated serrated lamellar colony boundaries have high strength and can inhibit crack propagation. The main reason for this phenomenon is that straight boundaries can serve as rapid pathways for crack propagation, reducing lateral plasticity or fracture toughness. For the passivated serrated lamellar colony boundaries, dislocation plugging occurs at the passivated serrated lamellar colony boundary and produces a strengthening effect. On one hand, the zigzag boundary can effectively hinder the movement of dislocations and the propagation of cracks, thereby enhancing strength and toughness. On the other hand, the sawtooth shape can disperse local stress and delay failure. Straight boundaries can serve as rapid pathways for crack propagation, reducing lateral plasticity or fracture toughness.

At the plastic stage (920 °C), more B2 phases precipitate in both the lamellar colonies and the lamellar colony boundaries during the tensile process. The B2 phase that precipitates in the lamellar colonies experiences a small deformation due to the limitation of the surrounding lamella and does not exert a great influence on the fracture mechanism of the alloy [36], [37], [38], [39]. In the later stages of the tensile test, lamellar colony boundary sliding causes high local stress concentrations at lamellar colony boundaries. The B2 phase at the lamellar colony boundary does not rotate, and the degree of recrystallization is relatively low. However, the B2 phase is softer than the γ and a2 phases at high temperatures, which can promote lamellar boundary slipping during deformation, alleviate the stress concentration at the lamellar colony boundary, and improve the deformation ability of the lamellar colony boundary. The γ phase at the lamellar colony boundary has low hardness and good plasticity and forms smaller grains through repeated recrystallization, which is conducive to lamellar boundary sliding and a decrease in flow resistance during the alloy tensile process, thus coordinating tensile deformation and reducing the work hardening effect. At the late stage of the tensile process, the γ grains in the lamellar grain boundary break to form holes that connect to form large cracks, which are typical cavity aggregation fractures.

5 Conclusions

  1. The alloy is composed of lamellar colonies and lamellar colony boundaries. The lamellar boundaries have two different forms: straight strips and serrations.

  2. The main fracture mechanism in the brittle stage is cracking of the γ/γ phase boundary in lamellar colonies. The tensile fracture resistance of the serrated lamellar colony boundary is greater than that of the straight strip lamellae boundary because of the strengthening effect of interfacial dislocation entanglement.

  3. The main characteristic of the alloy in the plastic phase is the precipitation of the B2 phase in the lamellar colony and at the lamellar colony boundary. The alloy fractures along the lamellar colony boundary, which is due primarily to the deformation of the B2 precipitated phase and the repeated recrystallization of γ grains at the lamellar colony boundaries.

Corresponding author: Shunke Zhang, Guizhou Communications Polytechnic University, Qingzhen, 551400, China, E-mail:

Acknowledgments

This work was supported by the Scientific Research Project of Guizhou Communications Polytechnic University (2024ZD03ZK, 2024ZD04ZK), Guizhou Province Science and Technology Project (qiankehejichuMS[2026]236), Scientific Research Project of Guizhou Communications Polytechnic (2022ZD01KJ), Natural Science Research Project of Guizhou Higher Education Institutions of China (QJJ[2023]047), Intelligent Transportation Equipment Manufacturing Technology Innovation Team Project (qianjiaoji[2023]1100) and Guizhou Province Science and Technology Plan Project (QKHJC-ZK[2024] yiban604).

  1. Funding information: This work was supported by the Scientific Research Project of Guizhou Communications Polytechnic University (2024ZD03ZK, 2024ZD04ZK), Guizhou Province Science and Technology Project (qiankehejichuMS[2026]236), Scientific Research Project of Guizhou Communications Polytechnic (2022ZD01KJ), Natural Science Research Project of Guizhou Higher Education Institutions of China (QJJ[2023]047), Intelligent Transportation Equipment Manufacturing Technology Innovation Team Project (qianjiaoji[2023]1100) and Guizhou Province Science and Technology Plan Project (QKHJC-ZK[2024] yiban604).

  2. Author contribution: All authors have accepted responsibility for the entire content of this manuscript and consented to its submission to the journal, reviewed all the results and approved the final version of the manuscript. Ning Tian: study design, experimental operation, data collection, and drafting of the manuscript; Peng Zhang: data analysis, interpretation of results, and manuscript revision; Shunke Zhang: study conception and design, critical revision and final approval of the manuscript, and funding acquisition; Shulei Sun: sample collection and experimental support; Xiaojuan Shang: literature review and data verification; Aiquan Peng: experimental assistance and data organization; Danping Dang: data analysis and figure preparation.

  3. Conflict of interest: Authors state no conflict of interest.

  4. Data availability statement: The datasets generated during and/or analyzed during the current study are available from the corresponding author on reasonable request.

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Received: 2025-06-10
Accepted: 2026-01-06
Published Online: 2026-04-14

© 2026 the author(s), published by De Gruyter, Berlin/Boston

This work is licensed under the Creative Commons Attribution 4.0 International License.

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https://doi.org/10.1515/htmp-2025-0115
Keywords for this article
TiAl-Nb alloy; high-temperature tensile; B2 phase; fracture mechanism
Creative Commons
BY 4.0

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