Fractographic study on the competitive bending fracture mechanisms of High-Nb TiAl alloys

To address the challenges of limited room-temperature ductility and low damage tolerance in high-Nb TiAl alloys and to improve their overall mechanical performance, this study investigates their fracture mechanism and intrinsic brittleness through bending tests combined with detailed fractographic analysis. Through bending and unloading experiments combined with scanning electron microscopy (SEM) observations of fracture surfaces and specimen morphologies, the results revealed that cracks initiated abruptly and led to catastrophic fracture once the applied bending load reached a critical level. In addition to the main crack initiation source, one or two secondary initiation sites were also identified on the fracture surfaces of the unloaded specimens. Variations in the crack lengths observed under different external unloading stresses indicate differences in the location and size of the cleavage initiation zone, rather than implying a progressive crack initiation–coalescence–fracture process. The borides within the alloy acted as dislocation obstacles or ligament bridges, effectively impeding further microcrack propagation or altering crack paths. These mechanisms enhanced the alloy’s resistance to crack initiation and propagation, thereby improving its fracture toughness. The fracture process can be summarized as follows: when the grains at the notch root are favorably oriented, cracks initiate within these grains and propagate along favorable colony boundaries or interlamellar interfaces, releasing stored strain energy and triggering cleavage across the entire specimen. In contrast, when the notch-root grains are unfavorably oriented, strain energy accumulates near the crack tip until a critical threshold is reached, resulting in instantaneous brittle cleavage fracture.

Introduction

Due to their low density, high oxidation resistance, and excellent specific strength, TiAl alloys have emerged as promising lightweight structural materials for high-temperature applications in the range of 600 °C–900 °C (Dong et al., 2024; Ma et al., 2024; Han et al., 2018). They have been increasingly employed as structural components in the aerospace and automotive industries (Sun et al., 2025; Huang et al., 2024). However, the dominant deformation mechanisms of TiAl alloys during creep involve dislocation slip, climb, and twinning. Dislocations typically originate at phase boundaries and migrate into the γ phase under the combined effects of stress and thermal activation. Consequently, the limited high-temperature strength and poor room-temperature ductility of TiAl alloys have restricted their broader engineering applications (Tian et al., 2025; Liu et al., 2025).

To overcome these limitations, high-Nb TiAl alloys have been developed on the basis of conventional TiAl alloys. Owing to their superior high-temperature mechanical properties and oxidation resistance, these alloys have attracted considerable attention from both domestic and international researchers (Zhang H et al., 2023; Xue et al., 2024; Yang et al., 2025). Nevertheless, high-Nb TiAl alloys still exhibit several drawbacks, including low room-temperature ductility, limited damage tolerance, and a relatively high crack-propagation rate (Li et al., 2025; Xu et al., 2023; Yang et al., 2023). Therefore, it is necessary to investigate the intrinsic brittleness of these materials in greater depth to enhance their application potential while maintaining their advantageous properties. Chen performed uniaxial tensile tests on high-Nb TiAl alloys and found that both U-shaped and V-shaped notched specimens exhibited higher tensile strengths than smooth specimens, with notch-sensitivity ratios (NSR>1), indicating that the alloy was not notch-sensitive (Chen et al., 2024). Wang reported that increasing the applied stress reduced the number of microcracks and voids at the fracture surface, enlarging the brittle-fracture zone (Wang et al., 2025). The simultaneous addition of boron and carbon in the form of B₄C improved the room-temperature mechanical properties of as-cast Ti-48Al alloys, achieving a tensile strength of 517 MPa and an elongation of 0.47%. Xu demonstrated that heat-treated high-Nb TiAl alloys exhibited enhanced room-temperature strength and ductility compared with as-cast alloys, undergoing a ductile-to-brittle transition between 800 °C and 900 °C, where the fracture mechanism changed from brittle transgranular fracture to void nucleation and coalescence (Xu et al., 2025). More recently, several researchers have reported that grain refinement and the addition of stabilizing elements such as W and Cr can significantly enhance both the plasticity and strength of high-Nb TiAl alloys (Guo et al., 2024). The presence of large, low-specific-surface micro-segregation regions reduces the rate of cavity and crack formation along grain boundaries, which is unfavorable for boundary strengthening and toughness improvement (Wei et al., 2024). Zhang observed that with increasing applied stress, the number of microcracks and voids at the fracture surface decreased while the brittle-fracture zone expanded; however, the addition of B₄C slightly reduced both the room-temperature tensile strength and strain of high-Nb TiAl alloys (Zhang S et al., 2023).

Considering the limited damage tolerance, low fracture toughness, and high crack-propagation rate that still constrain the use of high-Nb TiAl alloys (Zhong and Lin, 2023; Liang et al., 2024; Ding et al., 2025), an in-depth study of their intrinsic brittleness is essential to improve their room-temperature mechanical performance while retaining their inherent high-temperature characteristics. In this work, the damage and fracture behavior of high-Nb TiAl alloys were systematically investigated through bending tests and detailed SEM observations of fracture surfaces and cross-sections.

Experimental materials and methods

The experimental material used in this study was a fully lamellar high-Nb TiAl alloy. The microstructure of the alloy at room temperature is shown in Figure 1. The average colony size of the lamellar structure was approximately 100–150 μm. The alloy was prepared by vacuum arc melting and then heat treated under a vacuum to achieve a fully lamellar microstructure. The chemical composition and the microstructural morphology of the alloy are listed in Table 1 and illustrated in Figure 1, respectively.

Figure 1

Figure 1. Microstructure of fully lamellar high-Nb TiAl alloy (Lamellar colony size ≈100–150 μm, SEM micrograph).

Table 1

Table 1. Chemical composition of high-Nb TiAl alloy.

The bending specimens were machined using a CKX-2AJ precision wire-cutting machine. Two types of specimens were prepared: (i) three-point bending specimens with a 3.0-mm-long straight notch and (ii) four-point bending specimens with a V-notch, as shown in Figures 2a–c. The geometric dimensions of the single-notched and double-notched four-point specimens are presented in Figures 2b,c, respectively.

Figure 2

Figure 2. Shape and dimensions of (a) three-point bending specimen and (b) single-notch four-point bending specimen, (c) double-notch four-point bending specimen (unit: mm).

The bending tests were conducted on an Instron 1341 testing machine manufactured by the Instron Engineering Corporation (USA). All tests were carried out at room temperature with a loading rate of 0.5 mm min^-1. From the automatically recorded load–displacement curves, the fracture load (P_f), ultimate tensile strength (σ_UTS), yield load (P_Y), and fracture work per unit area (W′) were determined. Based on Equations 1, 2, the maximum stress intensity factor (K_max) for the high-Nb TiAl alloy was calculated. Here, K_max represents the mode I stress intensity factor corresponding to the maximum applied load; (W) denotes the specimen height, (B) the specimen thickness, and (a) the notch length including the precrack. The span length was defined as (S = 4W), and the geometric correction factor was expressed as Y_(a/W).

$K_{\max }=\left(\frac{\mathrm{S}·P_{\max }}{B·W^{3/2}}\right)·Y_{\left(a/w\right)}(1)$

$Y_{\left(a/w\right)}=\left[\frac{3{\left(a/w\right)}^{1/2}\left\{1.99-a/w\left(1-a/w\right)\times \left[2.15-3.93\left(a/w\right)+2.7{\left(a/w\right)}^2\right]\right\}}{2\left(1+2a/w\right){\left(1-a/w\right)}^{3/2}}\right](2)$

To investigate the influence of different loading levels, five distinct unloading conditions were established. The ratios of the maximum applied load to the ultimate fracture load (σ_UTS) were set to 0.9, 0.8, 0.7, 0.6, and 0.5, respectively. Here, σ_UTS denotes the maximum bending stress measured in the fracture test. Each specimen was loaded to the prescribed fraction of σ_UTS, held for a short period, and then fully unloaded. After unloading, the middle portion of each specimen (1.1 mm thick) was cut from the region near the notch root using a wire electrical discharge machining (EDM) system. The detailed specimen preparation process is illustrated in Figures 3a,b.

Figure 3

Figure 3. Treatment of unloading specimens after bending: (a) unloading schematic; (b) cutting orientation (unit: mm).

After bending and unloading, the fracture surfaces and cross-sections of all specimens were examined using a JSM-6700F scanning electron microscope (SEM). The observations focused on the morphological features of the crack-initiation sites and propagation paths. The relationships between the applied stress, crack morphology, and fracture mechanisms were analyzed in detail. Additionally, the damage characteristics and crack-propagation behaviors were compared between specimens under different unloading conditions. Representative features were photographed and statistically analyzed to reveal the evolution of the fracture process in the high-Nb TiAl alloy.

Results and discussion

Analysis of bending fracture test

The results of the three-point bending fracture tests for the high-Nb TiAl alloy specimens are presented in Table 2. From Table 2, it can be seen that the average fracture work per unit area (W′) of the fully lamellar structured specimens is 3.479 × 10⁻³ J/mm², and the average maximum fracture toughness (K_max) is 16.337 MPa·m1/2. Among all the samples, 3 PB-03 exhibited the highest fracture work and fracture toughness. This may be related to the fact that specimen 3 PB-03 experienced more interlamellar fracture during the fracture process. Since interlamellar cracking requires greater energy for propagation, specimens with higher proportions of translamellar fracture surfaces tend to exhibit higher fracture energy and fracture toughness.

Table 2

Table 2. Data of three-point bending fracture test of high-Nb TiAl alloy.

As shown in Figure 4, the load–displacement curves of specimens 3 PB-01, 3 PB-02, 3 PB-03, and 3 PB-05 have almost identical slopes, while specimen 3 PB-04 shows a similar trend but slightly lower load response. All curves are located within the elastic stage, indicating that the material underwent sudden cleavage fracture under a certain external load rather than progressive crack growth and coalescence with increasing load.

Figure 4

Figure 4. Load–displacement curves of three-point bending test.

To compare the fracture toughness of the materials, the maximum fracture toughness values of the fine-grained lamellar specimens and the coarse-grained lamellar specimens were analyzed (Shen et al., 2015). It was found that the fine-grained lamellar specimens exhibited higher maximum fracture toughness values than the coarse-grained ones, indicating that the fracture toughness of the alloy decreases with the refinement of lamellar colonies and the increase of colony size.

The fracture morphology of specimen 3 PB-01 under three-point bending is shown in Figure 5. As shown in Figure 5a, the specimen exhibits a main crack initiation site (indicated by solid arrow 1) at the notch root and a distinct crack propagation trace in the form of cleavage river patterns. This main crack was generated by stress concentration at the notch root. Although the specimen contained a machined notch, the stress was relatively uniformly distributed because of the square cross-section (approximately 6 mm thick). The weakest region near the notch root still contained more than 30 lamellar colonies, whose stresses were comparable; therefore, considerable energy accumulated ahead of the notch before fracture, resulting in the formation of a primary cleavage crack. At the same time, another secondary crack initiation site appeared at a certain distance from the notch root, as indicated by solid arrow two in Figure 5a.

Figure 5

Figure 5. Fracture surface of specimen 3 PB-01 under three-point bending: (a) overall fracture surface; (b) local fracture morphology; (c) river-like pattern on the fracture surface.

Compared with the tensile specimens studied by previous scholars (Wang et al., 2020), the three-point bending specimens also exhibited obvious initiation sites and river-like fracture patterns. However, the tensile specimens had a single crack initiation site, whereas the bending specimens exhibited one main and one or two secondary initiation sites. Figure 5b shows the fracture surface photographed from the region between the two initiation sites near the notch root. The river-like pattern extending from the main initiation site can be clearly observed, confirming that a dominant main crack initiated and propagated in this region. The distance between the main initiation site and the notch root was approximately 70 μm, and this crack played an important role in the final fracture. Further analysis revealed that many translamellar river-like patterns appeared on the lamellar planes of the fracture surface, as shown in Figure 5c. This indicates that the fracture process of the alloy during bending was not a gradual crack nucleation and growth process; instead, due to the concentrated main stress field localized in a narrow region ahead of the notch root, cracks preferentially initiated along weak lamellar cleavage planes and then propagated in a sudden brittle manner.

The fracture surface of specimen 3 PB-02 under three-point bending is shown in Figure 6. The crack initiation site is still located at the notch root, as indicated by the solid arrow in Figure 6a, and the size of the initiation region is approximately 200 μm × 88 μm. Once this initiation region-formed along lamellar interfaces—was activated, it triggered cleavage fracture throughout the entire specimen. The fracture surface consists of both interlamellar and translamellar facets, and a magnified view of its fracture path and river pattern direction is shown in Figure 6b. The macroscopic load–displacement curve of 3 PB-02 shows a slight deviation from the purely elastic stage (Figure 4), which is consistent with the presence of a well-defined cleavage origin at the notch root. Overall, 3 PB-02 failed in a brittle-dominated manner initiated at the notch root.

Figure 6

Figure 6. Fracture surface of specimen 3 PB-02 under three-point bending: (a) overall view of the fracture surface; (b) enlarged view of a section of the fracture path.

The four-point bending fracture test results for the high Nb–TiAl alloy specimens are summarized in Table 3. As shown in the table, the fracture work per unit area of the specimens ranges from 0.8 × 10⁻² to 1.0 × 10⁻² J/mm². The specimens with a single notch exhibited slightly higher fracture work per unit area than those with double notches, but the difference in overall bending performance was not significant. The corresponding load–displacement curves of the four-point bending fracture tests are shown in Figure 7, indicating that the load and displacement maintained a nearly linear relationship, and fracture occurred during the elastic stage.

Table 3

Table 3. Results of four-point bending fracture test for high Nb–TiAl alloy.

Figure 7

Figure 7. Load–displacement curves of four-point bending fracture test.

The fracture morphology of the single-notched four-point bending specimen 4 PB-02 is shown in Figure 8. Analysis revealed a main crack initiation site on the fracture surface, as indicated in Figure 8a, together with distinct river-like patterns, which resulted from a relatively uniform stress distribution at the notch root. Although the four-point bending specimen contained a notch and stress concentration did occur at the notch root, the specimen’s square cross-section (about 6 mm thick) contained more than 40 lamellar colonies, whose stresses were roughly equivalent. Consequently, before fracture, substantial energy was accumulated at the notch front, leading to the formation of a dominant cleavage crack that played a critical role in the overall fracture of the specimen. As shown in Figure 8b, secondary cracks were also observed near the notch root between two main initiation sites, indicating that crack propagation started close to the notch and extended along multiple lamellar boundaries. Numerous translamellar and interlamellar river patterns were observed across the fracture surface, as illustrated in Figure 8c, suggesting that the fracture process was not a gradual nucleation-and-growth mechanism. Instead, although the main stress field was confined to a narrow region ahead of the notch root, the crack preferentially initiated in locally weaker zones, where cleavage occurred along multiple lamellar planes. The river patterns displayed well-defined propagation directions, as shown in Figure 8d. Overall, the fracture process can be summarized as follows: after the initial crack initiation, continuous loading caused energy accumulation until a critical level was reached, at which point catastrophic cleavage fracture occurred throughout the specimen.

Figure 8

Figure 8. Fracture morphology of single-notched four-point bending specimen 4 PB-02: (a) macroscopic fracture morphology; (b) locally magnified view of fracture surface; (c) trans-lamellar and inter-lamellar fracture; (d) river pattern on the fracture surface.

During the analysis of the double-notched specimen 4 PB-10, it was found that the fracture surface no longer exhibited a single crack initiation site. Instead, two main crack initiation sources were formed, as indicated by the two small arrows at the notch roots in Figure 9. These cracks propagated along their respective paths independently: some propagated directly toward the specimen edges, while others intersected with each other, as shown by the long arrow in Figure 9. After intersecting, the cracks continued to propagate along their original directions, maintaining nearly constant propagation intensity and velocity, until they extended to the specimen boundaries and caused complete fracture. The fracture mechanism can thus be attributed to the competition between different crack initiation sites and their respective propagation paths, ultimately leading to a cleavage-type fracture of the specimen.

Figure 9

Figure 9. Fracture morphology of double-notched four-point bending specimen 4 PB-10.

Analysis of bending unloading test

The macroscopic unloading test results of notched three-point bending specimens of the high-Nb TiAl alloy are shown in Table 4.

Table 4

Table 4. Experimental data of three-point bending unloading tests for high-Nb TiAl alloy.

The SEM images of the cross-sections of partially unloaded three-point bending specimens are shown in Figure 10. The unloaded surface of specimen 3 PB-08, which was subjected to a loading force of 389 N, is presented in Figure 10a. No apparent damage or cracks can be observed on the specimen surface; however, numerous unfavorably oriented borides are distributed near the notch root of specimen 3 PB-08. These borides effectively hindered crack initiation and propagation at the notch root.

Figure 10

Figure 10. Cross-sectional morphologies of three-point bending specimens under different unloading loads: (a) specimen 3 PB-08 (389 N); (b) specimen 3 PB-09 (540 N); (c) specimen 3 PB-11 (706 N); (d) specimen 3 PB-12 (790 N).

With the increase of unloading load, material damage could already be observed in specimen 3 PB-09, which was subjected to a loading force of 540 N, as indicated by the arrows and the black-outlined region in Figure 10b. The damage appeared at a certain distance from the notch root rather than directly at the root, and it could only be clearly identified under higher magnification. As shown in Figure 10b, the absence of crack initiation from the notch root was mainly attributed to the unfavorable grain orientation in that region. It can also be observed that cracks tended to accumulate and align parallel to each other within the lamellar colonies that had favorable orientations, suggesting that this phenomenon might be related to interlamellar cracking. These lamellae were aligned in parallel, and when local stress concentration occurred, parallel microcracks were generated. However, careful examination revealed that the cracks were still very small.

At an unloading load of 706 N, clear crack morphology extending from the notch root can be observed in specimen 3 PB-11, as shown in Figure 10c. In the image, the crack initiated at the center of the notch root and propagated along a favorable orientation that was nearly parallel to the loading direction. This indicates that in the notched three-point bending tests of the high-Nb TiAl alloy, when the grains at the notch root were favorably oriented, crack initiation first occurred within those grains, followed by crack propagation along favorable colony boundaries or interlamellar regions. During crack propagation, the externally applied energy was released gradually. However, when the grain orientation at the crack tip was unfavorable, the accumulated strain energy could not be efficiently released, leading to a sudden and complete brittle cleavage fracture once the stored energy reached a critical level.

When the load increased to 790 N, a dense crack zone was observed on the specimen surface, as shown in Figure 10d. However, in specimen 3 PB-12, no cracks were found near the notch root; instead, a large number of cracks appeared at a certain distance away from it. This behavior was similar to that of specimen 3 PB-08, except that the grain misorientation near the notch root in 3 PB-12 was more pronounced, and the unfavorably oriented grains were larger in size. As a result, crack initiation at the notch root was extremely difficult, and cracks could only nucleate within favorably oriented grains located farther away from the notch. Consequently, specimen 3 PB-12 exhibited higher bending resistance than the other specimens.

From Figure 10 and Table 4, it can be concluded that the number of cracks in the notched high-Nb TiAl alloy specimens during the three-point bending unloading tests generally increased with the rise of unloading load.

Numerous borides of varying sizes were also observed on the surface of the unloaded specimens. To identify these borides, an XRD analysis was performed on the notch region of specimen 4 PB-11, as shown in Figure 11. The large-sized borides tended to induce dislocation pile-ups, and when these borides—acting as ligament bridges—were pulled out or fractured, they effectively hindered the further propagation of microcracks or altered their propagation paths. This behavior significantly increased the resistance to crack initiation and growth. In addition, the small-sized borides contributed to second-phase strengthening. Therefore, the dispersed distribution of borides markedly enhanced the fracture toughness of the high-Nb TiAl alloy. However, as shown in Figure 10, many damage sites were found within approximately 100 μm from the notch root. These microcracks tended to accumulate and align parallel to each other within lamellar colonies that had favorable orientations, which was likely related to interlamellar cracking. The parallel arrangement of these lamellae led to the formation of parallel cracks under local stress concentration. Nevertheless, detailed observation revealed that these cracks were still very small. Such damage and microcracks facilitated the partial release of the accumulated strain energy in the specimen. Comparison of these experimental results reveals that the fracture mode of the notched high-Nb TiAl alloy under bending is brittle cleavage fracture. During the fracture process, different main crack initiation sites propagated along their respective directions toward the specimen boundaries, indicating a cleavage fracture controlled by crack initiation.

Figure 11

Figure 11. XRD pattern of the notch region of bending specimen 4 PB-11.

To investigate the crystallographic features involved in the brittle fracture process of the high-Nb TiAl alloy, TEM analyses were conducted on the microstructures in the vicinity of the fracture surfaces of the bent specimens. The microscopic morphology at the interfaces between lamellar colonies and their colony boundaries is shown in Figure 12. Figure 12a shows the bright-field TEM image of the region near the fracture surface of specimen 3 PB-07 unloaded at a bending load of 204 N. At this stage, no obvious dislocations were observed in the material. The diffraction spots corresponding to the two alternating phases within the lamellar colonies are marked in the lower-left and upper-right corners of Figure 12a. As shown in all subfigures of Figure 12, the bright regions correspond to the γ phase (TiAl), whereas the dark regions correspond to the α₂ phase (Ti₃Al). In Figure 12b, a typical dislocation network (indicated by arrow A) and twin bands (indicated by arrow B) are present on the cross-section of specimen 3 PB-09 unloaded at 540 N. Figure 12c presents the TEM micrograph of the lamellar colony region near the fracture surface of specimen 3 PB-01 after bending fracture at 750 N. Numerous twins were emitted from the γ/α₂ lamellar interfaces and propagated for a short distance within the γ lamellae before terminating at the γ/α₂ interface. Most of the twins formed an angle of approximately 45° with the γ/α₂ interface. Figure 12d shows the microstructure near the fracture surface of specimen 3 PB-02 after bending fracture at 796 N. A large number of dislocations and dense twin bands can be observed, and the deformation of the lamellar colonies remained primarily concentrated within the γ lamellae. Both twinning and dislocation formation could occur at the γ/α₂ interfaces, as illustrated in Figure 12d. The interaction between twins and the lamellar interfaces could induce dislocations to bow out from the interface.

Figure 12

Figure 12. TEM analysis of the microstructure near the fracture surface of the bent specimens: (a) 3 PB-07 unloaded specimen; (b) 3 PB-09 unloaded specimen; (c) 3 PB-01 fractured specimen; (d) 3 PB-02 fractured specimen.

Based on the analysis, the deformation of the high-Nb TiAl alloy during bending was primarily governed by twinning and dislocation activity. Because the deformation of the γ lamellae was constrained by the neighboring α₂ lamellae, deformation preferentially occurred at the lamellar colony boundaries. Dislocations were unable to pass through the interfaces between the lamellar colonies and the colony boundaries, leading to dislocation pile-ups at these interfaces. As the applied stress increased, dislocations and dislocation arrays bowed out from the γ/α₂ interfaces, and twins were emitted from the same interfaces. This indicates that the γ and α₂ lamellae exhibit different deformation characteristics, resulting in significant interface stresses. The twins, dislocations, and dislocation arrangements generated at the interfaces effectively alleviated local stress concentrations, thereby enhancing the plasticity and damage tolerance of the high-Nb TiAl alloy.

From the above bending test results and analysis, it can be concluded that in the notched three-point bending tests of the high-Nb TiAl alloy, when the grains at the notch root were favorably oriented, crack initiation first occurred within those grains. The cracks then propagated along favorable colony boundaries or interlamellar regions, during which the external work was released gradually. However, when the orientation at the crack tip was unfavorable, the strain energy continued to accumulate near the crack front. Once the strain energy in this region reached a critical level, the crack nucleated and rapidly propagated, leading to fracture. The fracture mode conformed to the characteristics of a completely brittle cleavage fracture, and the corresponding crack propagation path is illustrated in Figure 13. The differences in surface crack length observed in the unloaded specimens under different external loads indicate that the locations of the crack initiation regions varied. Nevertheless, this does not imply that the fracture process of this material involves crack nucleation, coalescence, and gradual propagation until final failure. Rather, it suggests that under different unloading loads, only the size of the regions responsible for initiating cleavage fracture differs.

Figure 13

Figure 13. Fracture path of bending specimen of high-Nb TiAl alloy.

Conclusion

1. The bending damage and fracture mechanism of the alloy can be summarized as follows: When the grains at the notch root were favorably oriented, crack initiation first occurred within these grains and then propagated along favorable colony boundaries or interlamellar regions, releasing external work during propagation. Once a crack initiated along such lamellar planes, it led to complete cleavage fracture of the specimen. Conversely, when the grain orientation at the notch root was unfavorable, the strain energy at the crack tip continuously accumulated due to restricted propagation. When this local strain energy reached a critical level, a crack suddenly initiated and rapidly propagated, resulting in a completely brittle fracture mode.

2. A comparison between the three-point and four-point bending tests reveals that both exhibit a clear dependence of the fracture path and fracture toughness on grain orientation. The distinction lies in the fact that the stress in the three-point bending specimens is more highly concentrated, leading to a competition mechanism governed by grain orientation primarily within the crack-initiation region. In contrast, the stress distribution in the four-point bending specimens is relatively less concentrated, resulting in a higher likelihood that such grain-orientation-driven competing mechanisms occur not only at the crack-initiation zone but also throughout the entire fracture process.

3. The dispersed borides in the high-Nb TiAl alloy enhanced the resistance to crack initiation and propagation through dislocation pile-ups and the formation of ligament bridges, thereby significantly improving the fracture toughness of the alloy.

Data availability statement

The original contributions presented in the study are included in the article/supplementary material, further inquiries can be directed to the corresponding author.

Author contributions

YL: Validation, Methodology, Writing – original draft, Conceptualization. PL: Writing – review and editing, Resources. BL: Writing – review and editing, Investigation, Conceptualization.

Funding

The author(s) declared that financial support was received for this work and/or its publication. PL had been funded in part by the Texas A&M University and the China Scholarship Council. This work was supported by the National Natural Science Foundations of China (No. 50471109), the Guangdong Provincial Educational Science Planning Project (Higher Education Special Project) (No. 2025GXJK0067), the National Undergraduate Innovation and Entrepreneurship Training Program (No. 202412622003), the Zhanjiang University of Science and Technology Undergraduate Innovation and Entrepreneurship Training Program (No. 2024ZKDCJZ71; ZJKJXYPDJHA202406), and Zhanjiang University of Science and Technology 2024 Undergraduate Teaching Quality and Reform Project (No. WCRHJG-202405).

Conflict of interest

The author(s) declared that this work was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

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