As-cast ti-al alloy with alloying optimization and comprehensive mechanical properties and preparation method thereof

Abstract

The application belongs to the technical field of light high-temperature structural materials, and more particularly relates to a cast TiAl alloy with alloying optimized comprehensive mechanical properties and a preparation method thereof. Through multi-element synergistic alloying design of Zr, beta-stabilizing elements and C, the application realizes comprehensive optimization of room-temperature strength and plasticity and high-temperature creep performance without sacrificing the low-density advantage of the alloy, and good balance is achieved between various performance indexes.

Description

Technical Field

[0001] This invention belongs to the field of lightweight high-temperature structural materials technology, and more specifically relates to a cast TiAl alloy with optimized comprehensive mechanical properties through alloying and its preparation method. Background Technology

[0002] With the rapid development of the aerospace industry, stringent requirements have been placed on the thrust-to-weight ratio, service temperature, and creep performance of materials used in aircraft structural components. High-temperature titanium alloys have a thermal barrier temperature of approximately 600 °C, above which their mechanical properties decline sharply. While nickel-based superalloys exhibit excellent high-temperature performance, their excessive density fails to meet lightweight requirements. TiAl-based intermetallic compounds, on the other hand, have a density of only 3.7–4.1 g / cm³. 3 It possesses high specific strength, high specific stiffness, excellent flame retardancy, and outstanding high-temperature oxidation and creep resistance. In the 600–900 °C range, its specific strength surpasses that of nickel-based superalloys and traditional titanium alloys, and its specific modulus is 50%–70% higher than that of titanium alloys, making it an ideal candidate material for lightweight structural components in aero-engines. Currently, GE's Ti4822 alloy is used in the low-pressure turbine blades of the GEnx engine, Germany's TNM alloy has been applied to the PW-1100G engine, and TiAl alloy has moved from theoretical research to practical application.

[0003] Despite significant progress in TiAl alloy research, existing technologies still face several prominent challenges: First, TiAl alloys exhibit severe intrinsic brittleness, which cannot be fundamentally addressed by current single-element alloying strategies. Second, the operating temperatures of existing mature grades are limited; Ti4822 operates below 700 °C, and TNM below 750 °C, making it difficult to meet higher-temperature requirements. Third, Nb, as a traditional β-stabilizing element, readily induces β-segregation during solidification, requiring prolonged isothermal treatment in the α-phase region to eliminate it, resulting in high processing costs. Furthermore, Nb is a rare strategic metal with limited resources and high prices. Fourth, while Zr can improve brittleness and reduce segregation tendency, it shrinks the α-phase region and increases the difficulty of heat treatment; currently, a complete and mature alloy system to replace Nb has not yet been developed. Fifth, existing multi-element alloying research lacks sufficient understanding of the synergistic mechanisms between elements and lacks systematic composition optimization design methods, making it difficult to achieve a comprehensive optimal match between room-temperature strength and plasticity and high-temperature creep performance. Summary of the Invention

[0004] The purpose of this invention is to provide an as-cast TiAl alloy with optimized comprehensive mechanical properties through alloying and its preparation method, so as to solve the problems existing in the prior art.

[0005] To achieve the above objectives, the present invention provides the following solution: One of the technical solutions of this invention is to provide an as-cast TiAl alloy with optimized comprehensive mechanical properties through alloying. The composition, by atomic percentage, includes: 40-50% Al, 0-5% Zr, 0-3% β-stabilizing elements, and 0-2% C, with the balance being Ti. The β-stabilizing element includes at least one of Cr, Mn and Fe; Among them, the amounts of Zr, β-stable elements, and C are not zero.

[0006] Furthermore, by atomic percentage, the composition of the as-cast TiAl alloy comprises: 40-50% Al, 1-3% Zr, 1.5-2.5% β-stabilizing elements and 0.5-2% C, with the balance being Ti.

[0007] The second technical solution of this invention provides a method for preparing an as-cast TiAl alloy with optimized comprehensive mechanical properties through alloying, comprising the following steps: Prepare raw materials according to the composition of the as-cast TiAl alloy described above; The raw material is placed in a vacuum electric arc furnace and then subjected to electric arc melting under argon protection. After cooling, the melting is repeated at least once to obtain the as-cast TiAl alloy.

[0008] Furthermore, the raw materials include titanium, aluminum, zirconium, and graphite powder, as well as at least one of chromium, manganese, and iron.

[0009] Except for graphite, which is in powder form, all other raw materials are in granular form.

[0010] Optionally, the graphite powder has a particle size of 200-400 mesh, preferably 300 mesh.

[0011] Furthermore, the process prior to arc melting also includes a gas washing operation.

[0012] Optionally, the gas scrubbing operation includes the following steps: vacuuming (5×10⁻⁶). -3 After filling with argon gas (0.05 MPa), then evacuating to a vacuum (5 × 10⁻⁶ MPa), the pressure was reduced to 0.05 MPa. -3 After Pa), argon gas is introduced for protection (0.05 MPa).

[0013] Furthermore, the current increase rate of the electric arc melting is 10-30 A / s (preferably 20 A / s), and the current is 350-450A (preferably 400 A). After all the raw materials are melted, the current is kept constant for 1-3 minutes for heat preservation, and the furnace is cooled after the heat preservation is completed.

[0014] Evaporation loss is unavoidable during the smelting process. To compensate for the evaporation loss during smelting, 3 at.% Al particles and 5 at.% Mn particles are added to ensure uniform chemical composition (other β-stabilizing elements do not need to be added).

[0015] Furthermore, the repeated melting involves flipping the cooled ingot and melting it, and the number of times is repeated is 3-7.

[0016] The third technical solution of the present invention provides an application of the above-mentioned as-cast TiAl alloy as a high-temperature structural material.

[0017] Compared with existing technologies, the as-cast TiAl alloy and its preparation method provided by this technical solution have the following advantages: Significant improvement in room temperature mechanical properties: Zr element is dissolved in the γ phase, reducing its c / a axis ratio and effectively improving the intrinsic brittleness of TiAl alloy; β stabilizing element reduces stacking fault energy, promotes deformation twin formation, and activates more slip systems. The synergistic effect of the two makes the room temperature strength and plasticity of the alloy significantly improved.

[0018] Excellent high-temperature creep performance: C element, as an interstitial atom, generates a solid solution strengthening effect, which hinders dislocation movement; at the same time, C forms thermally stable carbides with Ti and Al, which effectively pins dislocations and inhibits grain boundary slip, significantly improving the alloy's creep resistance under high-temperature conditions.

[0019] Excellent high-temperature structural stability: During service, nano-carbides preferentially precipitate out of grain boundaries and phase interfaces, which can hinder interface migration during long-term high-temperature service, inhibit structural degradation, and ensure the structural reliability of the alloy in high-temperature environments.

[0020] Excellent overall mechanical properties: This invention achieves comprehensive optimization of TiAl alloy's room-temperature strength and high-temperature creep resistance through a multi-element synergistic alloying design using Zr, β-stabilizing elements, and C. Zr substitution of Ti introduces lattice distortion, reducing the c / a ratio of the γ phase to improve intrinsic brittleness, while simultaneously increasing the interstitial solid solubility of C. β-elements such as Cr / Mn synergistically promote the formation of bulk γ phase with Zr and synergistically stabilize the B2 phase with C. C and Zr together increase the atomic diffusion barrier, while the formed Ti2AlC / Ti3AlC carbides synergistically refine the lamellar clusters with the β-elements. All three elements work together to enhance room-temperature strength through multi-level superposition of solid solution strengthening, grain refinement strengthening, and second-phase strengthening, and synergistically improve creep resistance. This invention, without sacrificing the advantage of low density, balances high strength, good plasticity, and excellent creep resistance, making it suitable for hot-end components of aero-engines.

[0021] The preparation process is simple and the cost is controllable: the vacuum arc melting process is adopted, the preparation process is simple, the amount of alloying elements added is small, the impact on the alloy density is minimal, it has good compatibility with existing melting processes, and it is easy to promote and apply in engineering. Attached Figure Description

[0022] The accompanying drawings, which form part of this invention, are used to provide a further understanding of the invention. The illustrative embodiments of the invention and their descriptions are used to explain the invention and do not constitute an undue limitation of the invention. In the drawings: Figure 1 The images show the XRD patterns of the as-cast TiAl alloys in Examples 1 and 2.

[0023] Figure 2 The initial microstructures of the as-cast TiAl alloys of Examples 1 and 2 are shown, wherein (a) and (b) are TZC-0.5C, (c) and (d) are TZC-1C, (e) and (f) are TZC-1.5C, and (g) and (h) are TZC-2C.

[0024] Figure 3 The images are transmission electron microscopy (TEM) images of TZC-1C, where (a) shows the morphology of the blocky γ phase interface, (b)-(f) are the elemental distribution diagrams of (a), (g) shows the γ phase diffraction spots, (h) shows the internal morphology of the lamellar clusters, and (i) shows the α2 phase diffraction spots.

[0025] Figure 4 The images are transmission electron microscopy (TEM) images of TZC-1.5C, where (a) shows the γ-rays at the lamellar boundary. b (a) Phase B2; (b) Diffraction spots corresponding to region A; (c) Laminar cluster morphology; (d) Diffraction spots of region B; (e) Ti2AlC morphology; (f) Diffraction spots of region C; (g)-(l) Ti2AlC morphology and corresponding elemental distribution diagram inside laminar clusters.

[0026] Figure 5 A schematic diagram illustrating the effect of adding different C element contents on the microstructure morphology.

[0027] Figure 6 The room temperature compression curves are for the as-cast TiAl alloys of Examples 1, 2, and 1.

[0028] Figure 7 The room temperature compression curves are for the as-cast TiAl alloys of Comparative Examples 1-3.

[0029] Figure 8 Crack propagation and fracture morphology of the as-cast TiAl alloys of Examples 1 and 2 are shown, where (a) and (b) are TZC-0.5C, (c) is TZC-1C, (d) and (e) are TZC-1.5C, and (f) is TZC-2C.

[0030] Figure 9 The compression creep properties of different as-cast TiAl alloys in the examples are shown, where (a) is the creep strain-time curve and (b) is the corresponding creep rate-time curve.

[0031] Figure 10 To compare the compressive creep properties of different as-cast TiAl alloys in the comparative examples, (a) is the creep strain-time curve, and (b) is the corresponding creep rate-time curve.

[0032] Figure 11 The images shown are BSE morphology diagrams of different as-cast TiAl alloys after creep experiments in the examples. (a) and (b) are TZC-0.5C, (c) and (d) are TZC-1C, and (e) and (f) are TZC-1.5C.

[0033] Figure 12 The images show the BSE morphology of different as-cast TiAl alloys after creep experiments in the comparative examples. (a)-(c) represent TZ, (d)-(f) represent TZ-Cr, and (g)-(i) represent TZ-Mn.

[0034] Figure 13 These are TEM images of the microstructure of TZC-1C after creep testing, where (a) shows the morphology at the boundary of lamellar clusters, and (b) shows the γ-ray structure. b The precipitation of needle-like C14 phase in the phase, (c) shows the morphology of Ti3AlC in lamellar clusters, and (d) shows the γ-phase precipitation. b The precipitated phases in (d) are shown in (e)-(g), which are the diffraction results of different phases in (d), and (h)-(l) are the corresponding elemental distribution diagrams of (d).

[0035] Figure 14 The images are TEM images of the microstructure of TZ-Cr alloy after creep testing. (a) shows the bent lamellar structure, (b) shows the fractured α2 laths, (c) shows the diffraction results of the residual α2 phase, (d) shows the precipitates around the lamellar cluster boundaries, (e) shows the precipitates in the recrystallization region, (f) shows the precipitates around the blocky γ phase, and (g)-(j) are the elemental distribution diagrams of (d).

[0036] Figure 15 The images show the B2 phase morphology of the TZC-1C alloy after creep, where (a) is the B2 phase morphology, (b)-(f) are the element distribution diagrams of (a), (g) is the morphology of Ti3AlC, (h) is the diffraction result of the B2 phase, and (i) is the diffraction result of Ti3AlC.

[0037] Figure 16 The images show the microstructure of the TZC-1.5C alloy after creep testing, where (a) represents the γ... bThe morphology of the internal precipitated phases is shown in (b)-(d), which are the diffraction results of different phases in (a). (c) is the morphology of Ti3AlC in the lamellar clusters. (e)-(i) are the elemental distribution diagrams of a. (j)-(i) are the morphology and diffraction results of the B2 phase at the interface.

[0038] Figure 17 The images show the internal morphology of the lamellar clusters of the TZC-1.5C alloy after the creep test. Among them, (a) is the morphology of Ti2AlC, (b)-(f) are the element distribution diagrams of (a), (g) is the morphology of Ti3AlC in the lamellar clusters, (h) is the diffraction result of Ti3AlC, and (i) is the diffraction result of the γ phase.

[0039] Figure 18 The initial microstructure of the as-cast TiAl alloy in Example 3 is shown in (a) and (b), which are TZM-0.5C, and (c) and (d), which are TZF-0.5C.

[0040] Figure 19 The room temperature compression curve of the as-cast TiAl alloy in Example 3 is shown. Detailed Implementation

[0041] Various exemplary embodiments of the present invention will now be described in detail. This detailed description should not be considered as a limitation of the present invention, but rather as a more detailed description of certain aspects, features, and embodiments of the present invention.

[0042] It should be understood that the terminology used in this invention is merely for describing particular embodiments and is not intended to limit the invention. Furthermore, with respect to numerical ranges in this invention, it should be understood that each intermediate value between the upper and lower limits of the range is also specifically disclosed. Any stated value or intermediate value within a stated range, as well as each smaller range between any other stated value or intermediate value within said range, is also included in this invention. The upper and lower limits of these smaller ranges may be independently included or excluded from the range.

[0043] Unless otherwise stated, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art. While only preferred methods and materials have been described herein, any methods and materials similar or equivalent to those described herein may be used in the implementation or testing of this invention. All references to this specification are incorporated by way of citation to disclose and describe methods and / or materials associated with those references. In the event of any conflict with any incorporated reference, the content of this specification shall prevail.

[0044] Various modifications and variations can be made to the specific embodiments described in this specification without departing from the scope or spirit of the invention, as will be apparent to those skilled in the art. Other embodiments derived from this specification will also be apparent to those skilled in the art. This specification and embodiments are merely exemplary.

[0045] The terms “include,” “including,” “have,” “contain,” etc., used in this article are all open-ended terms, meaning that they include but are not limited to.

[0046] Unless otherwise specified, all raw materials and reagents involved in the specific embodiments of this invention are commercially available products.

[0047] Unless otherwise specified, room temperature and ambient temperature in the specific embodiments of this invention refer to 20-30 ℃.

[0048] It should be noted that any aspects not described in detail in this invention are conventional practices in the field and are not the focus of this invention.

[0049] In the specific embodiments of the present invention, the purity of the raw materials involved, namely aluminum particles, titanium particles, zirconium particles, chromium particles, manganese particles and iron particles, is not less than 99.9 wt%, and the particle size of the graphite powder is 300 mesh.

[0050] In a specific embodiment of the present invention, the vacuum arc furnace used for arc melting is a DHL-600 vacuum arc furnace.

[0051] Example 1 The preparation steps of the as-cast TiAl alloy (Ti-46Al-2Zr-2Cr-0.5C) with optimized comprehensive mechanical properties by alloying include: S1. Prepare titanium, aluminum, zirconium, chromium particles and graphite powder, with the following atomic percentages: It contains 46% Al, 2% Zr, 2% β-stabilizing element (Cr) and 0.5% C, with the balance being Ti.

[0052] Considering the oxidation loss of aluminum during the smelting process, a loss of 3% should be applied, and the actual weight should be weighed according to 103% of the aluminum mass.

[0053] S2. Place the raw materials in the crucible in order of increasing melting point (from bottom to top: Al, Ti, Zr, Cr, C), transfer to a vacuum arc furnace, and evacuate to 5 × 10⁻⁶. -3 Pa, then purged with argon gas at 0.05 MPa for gas washing, and then evacuated to 5 × 10 Pa. -3Argon gas at 0.05 MPa was introduced as a protective atmosphere. The current was then increased to 400 A at a rate of 20 A / s. After the raw material was completely melted, the current was kept constant for 3 minutes. Then, the current was reduced to 0 A at a rate of 50 A / s. After cooling, the ingot was flipped over and the above arc melting was repeated. The melting was repeated 5 times to obtain the as-cast TiAl alloy (Ti-46Al-2Zr-2Cr-0.5C), denoted as TZC-0.5C.

[0054] Example 2 The difference from Example 1 is that the amount of C used in step S1 is 1%, 1.5% or 2%.

[0055] When the amount of C is 1.0%, the product is denoted as TZC-1C.

[0056] When the amount of C is 1.5%, the product is denoted as TZC-1.5C.

[0057] When the amount of C is 2.0%, the product is denoted as TZC-2C.

[0058] Example 3 The difference from Example 1 is that the β-stabilizing element is replaced with an equal amount of Mn or Fe.

[0059] When the β-stabilizing element is Mn, the product is denoted as TZM-0.5C.

[0060] When the β-stabilizing element is Fe, the product is denoted as TZF-0.5C.

[0061] Comparative Example 1 Compared with Example 1, the difference is that the amount of C used in step S1 is 0%, and the product is Ti-46Al-2Zr-2Cr, denoted as TZ-Cr or TZC.

[0062] Comparative Example 2 The difference compared to Example 3 is that the amount of C used in step S1 is 0%. When the β-stabilizing element is Mn, the product is Ti-46Al-2Zr-2Mn, denoted as TZ-Mn.

[0063] When the β-stabilizing element is Fe, the product is Ti-46Al-2Zr-2Fe, denoted as TZ-Fe.

[0064] Comparative Example 3 The difference from Example 1 is that the atomic percentage in step S1 is 46% Al, 2% Zr, and the balance is Ti.

[0065] The product is Ti-46Al-2Zr, denoted as TZ.

[0066] Test case Samples were taken from the center and radial midpoint of the ingot using electrical discharge machining (EDM). The average chemical composition of the as-cast TiAl alloy was analyzed by energy dispersive spectroscopy (EDS). Three micro-regions were randomly selected from each sample for analysis, and the median value was chosen as the sample composition to avoid random errors. The results are shown in Table 1. The actual chemical composition of the as-cast alloy shows a slight deviation in atomic content from the designed composition, which is within the normal error range, indicating that the smelted ingots in the examples and comparative examples meet the design requirements.

[0067] Table 1 Actual chemical composition (at.%) The effect of C content on the microstructure and phase composition of TiAl alloys: Figure 1 The XRD patterns of the as-cast TiAl alloys in Examples 1 and 2 are shown. As shown, it can be seen that under different C additions, it is mainly composed of α2 and γ phases. The strongest diffraction peak appears at about 38°, corresponding to the (0001)α2 and (111)γ crystal planes, which conforms to the Blackburn orientation relationship. It is worth noting that when the C content is greater than 1.5 at.%, the diffraction peak corresponding to Ti2AlC begins to appear at about 39°, and the direction of this peak is (103). The diffraction peak corresponding to the 2C content is significantly enhanced compared to 1.5C. The (2021) crystal plane corresponding to α2 is near 41°. When the C content is less than 1.5 at%, the intensity of this peak gradually increases and then shows a decreasing trend. C atoms mainly occupy interstitial sites in the α2 phase, while α is transformed into α2 through solid-state phase transformation. The increase of α2 content means that interstitial carbon atoms can expand the α phase region. When excessive C addition causes C atoms to precipitate mainly in the form of carbides, the intensity of the corresponding peak begins to decrease. Phase B2 has no significant corresponding peak due to its low content.

[0068] Figure 2 The figures show the initial microstructures of the as-cast TiAl alloys in Examples 1 and 2, where (a) and (b) are TZC-0.5C, (c) and (d) are TZC-1C, (e) and (f) are TZC-1.5C, and (g) and (h) are TZC-2C. As can be seen from the figures, the TZC-0.5C alloy has a near-lamellar structure, mainly composed of α2 / γ lamellar clusters and massive γ phase (γ...). b Composed of lamellar clusters and γ b A small amount of B2 phase was present at the phase interface, consistent with the distribution and morphological characteristics of the B2 phase in TZC. γ b The γ phase is mainly distributed within the α2 / γ lamellae and at grain boundaries, appearing as irregular bands along the grain boundaries and as small elliptical shapes within the lamellar clusters. In TZC-1C, the γ phase... bThe phase content increases, of which γ b The increased number of phases forms a network-like structure. In TZC-1.5C, Ti2AlC phase precipitates, mainly distributed within the lamellar clusters as a needle-like structure. The addition of more C leads to γ... b The increased content leads to a network structure, with a small amount of B2 phase distributed at the interface. TZC-2C exhibits a higher amount of Ti2AlC precipitation, while γ... b It exhibits a similar network-like continuous structure, but γ b The phase is even larger.

[0069] The α2 and γ values ​​in the as-cast TiAl alloys prepared in Examples 1 and 2 b The EDS composition of the phase was analyzed, and the results are shown in Table 2.

[0070] Table 2. Composition of different phases (at.%) The γ phase in the TiAl alloy has an L10 structure, where Ti mainly occupies the 1a position and Al occupies the 1d position. Table 2 shows that the γ phase... b The Al content in the phase remains stable at approximately 50%, and the ratio of (Ti, Zr, Cr) to Al is close to 1:1. This indicates that in the TZ alloy, Zr and Cr atoms preferentially replace Ti atoms. The α2 phase exhibits D0... 19 The structure is such that Ti and Al occupy the 6h and 2c positions, respectively. The elemental distribution in the α2 / γ sheets exhibits a certain regularity: the Al content remains at approximately 45%, while the total content of Ti, Zr, and Cr atoms accounts for approximately 55%. Zr and Cr atoms preferentially replace Ti atom sites in the α2 phase.

[0071] To accurately assess the influence of C content on the phase content of as-cast TiAl alloys, the γ element content was quantitatively analyzed. b The volume fractions of Ti2AlC are shown in Table 3.

[0072] Table 3 Phase composition (volume fraction %) Table 3 shows that in TZC-0.5C and TZC-1C, carbon mainly exists in solid solution form. In TZC-1.5C, Ti2AlC begins to precipitate, with 2.52% carbides. In TZC-2C, the carbide content increases to 4.85%. Carbides are mainly distributed in large quantities within the lamellar clusters, and the carbide content is closely related to the carbon content. Ti2AlC may promote the nucleation of primary α-phase, thereby effectively reducing the size of the lamellar clusters.

[0073] The average lamellar cluster size and lamellar spacing of the as-cast TiAl alloys prepared in Examples 1 and 2 are shown in Table 4.

[0074] Table 4 Table 4 shows that the addition of carbon (C) has a significant impact on the average size of lamellar clusters and the interlamellar spacing. For TZC-0.5C and TZC-1C, the average lamellar cluster size decreases to some extent, while the interlamellar spacing increases slightly. For TZC-1.5C and TZC-2C, the average lamellar cluster size decreases significantly, while the interlamellar spacing increases even more significantly. With increasing C content, the average lamellar cluster size decreases while the interlamellar spacing gradually increases. The precipitation of Ti2AlC has a significant impact on both the average lamellar cluster size and the interlamellar spacing, but increasing the Ti2AlC content does not exacerbate this effect.

[0075] Figure 3 The images show transmission electron microscopy (TEM) images of TZC-1C, where (a) shows the morphology of the bulky γ phase interface, (b)-(f) are the elemental distribution diagrams of (a), (g) shows the γ phase diffraction spots, (h) shows the internal morphology of the lamellar clusters, and (i) shows the α2 phase diffraction spots. As can be seen from the figures, the addition of a small amount of C in the TZC-1C alloy did not change the phase structure type; the microstructure still contains γ phases. b Phases include α2 / γ lamellar structure and B2 phase. The B2 phase is located at the interface between the massive γ and lamellar structure, where Ti and Cr are enriched; Zr is enriched in the γ phase. b In the γ phase, Al and C are enriched in the lamellar structure. Diffraction analysis shows that the equiaxed γ phase is oriented at

[101] γ, and there are γ phases along the lamellar structure. The α2 phase is oriented in the α direction. C, as an α-stabilizing element, is enriched along the α phase, forming alternating α2 / γ lamellae upon cooling. Furthermore, the addition of C reduces the formation of twins in the equiaxed γ phase.

[0076] Figure 4 The images are transmission electron microscopy (TEM) images of TZC-1.5C, where (a) shows the γ-rays at the lamellar boundary. b (a) shows the morphology of the lamellar clusters; (b) shows the diffraction spots corresponding to region A; (c) shows the morphology of the lamellar clusters; (d) shows the diffraction spots of region B; (e) shows the morphology of Ti2AlC; (f) shows the diffraction spots of region C; (g)-(l) show the morphology of Ti2AlC inside the lamellar clusters and the corresponding elemental distribution diagrams. As can be seen from the figure, in (a), a black contrast precipitate can be seen at the interface between the lamellar structure and the blocky γ phase of the TZC-1.5C alloy. (b) shows the diffraction spots corresponding to region A in (a), and the black precipitate is further confirmed to be the B2 phase through diffraction spot analysis. (c) and (d) show the morphology inside the lamellar clusters and the corresponding diffraction spots. The diffraction analysis results show that the α2 phase with HCP structure and the γ phase with FCT structure are similar to the B2 phase. L The phases satisfy the Blackburn orientation relation: {0001}α2 / / {111}γ; (e) and (f) show the morphology of carbides precipitated in different regions. Analysis of the diffraction spots in this phase of Ti2AlC reveals the crystal orientation as […]. 110], needle-like. (g)-(l) are Ti2AlC inside the lamellar structure and corresponding EDS energy dispersive spectroscopy analysis. The carbides in the lamellar clusters are oriented at a certain angle to the lamellar structure. Elemental analysis shows that only C element is enriched in this phase.

[0077] Figure 5 A schematic diagram illustrating the effect of adding different C element contents on the microstructure morphology.

[0078] The addition of carbon (C) does not alter the basic solidification path of TZC alloys, but it significantly affects the phase transformation process. A small amount of C increases the nucleation temperature of the primary α phase, promoting the β→α peritectic reaction. Simultaneously, C enriches and displaces Zr atoms in the residual liquid phase, increasing Zr concentration and promoting γ phase nucleation, resulting in finer α and more γ phases. A high C content leads to the precipitation of Ti2AlC, which acts as nucleation sites, inducing rapid α phase growth, increasing the number of α phases, and stabilizing the α2 phase in subsequent transformations, slightly increasing lamellar thickness. In TZC-0.5 / 1C, C mainly exists in solid solution form. The addition of Zr, through lattice expansion and a reduction in solid solution free energy, increases the interstitial solid solution limit of C in γ-TiAl, enhancing solid solution strengthening and optimizing the deformation mechanism.

[0079] Room temperature compression performance: Sample: Ø6mm×9mm cylindrical specimen; Method: A Shimadzu AG-XPlus electronic universal testing machine was used to perform room temperature compression at a constant loading rate of 0.5 mm / min, and each ingot was repeated three times.

[0080] High-temperature creep performance: Conditions: 850℃ / 250MPa / 100h; Sample: Ø6mm×9mm cylindrical specimen; Method: The RDL100 electronic creep testing machine was used. The temperature was increased to 850℃ at 5℃ / min and held for 30min. The load was increased to 250MPa at 2kN / min and the temperature was controlled by multiple thermocouples (fluctuation ±1℃).

[0081] Figure 6The figures show the room temperature compression curves of the as-cast TiAl alloys in Examples 1, 2, and 1 (Comparative Example 1). As can be seen from the figures, the TZC alloy exhibits good plastic deformation capability, but its yield strength is relatively low. The TZC-0.5C alloy shows significantly lower room temperature plasticity and ultimate compressive strength compared to the TZC alloy, but its yield strength is improved to some extent. With the increase of carbon (C), the room temperature yield strength and ultimate compressive strength of the TZC-1 / 1.5C alloys gradually increase, while the plasticity also improves simultaneously. Although the TZC-2C alloy shows good strength and plasticity, its overall mechanical properties show a downward trend compared to TZC-1.5C. The addition of C can effectively improve the room temperature compressive strength, but it significantly reduces the material's plasticity. The addition of C, existing in the form of interstitial solid solution, reduces the material's plasticity but effectively improves its yield strength. The addition of C may regulate the material's mechanical properties by effectively refining the size of the lamellar clusters. However, adding excessive carbon will precipitate needle-like Ti2AlC, which will worsen the material's room temperature mechanical properties.

[0082] Table 5 Room Temperature Compression Performance Table 5 shows the room-temperature mechanical properties of TiAl alloys in different as-cast states. The TZC material has a yield strength of 769 MPa, an ultimate compressive strength of 2256 MPa, and a fracture strain of 35.75%. TZC-0.5C exhibits significantly reduced plasticity, with a fracture strain of 21.16%; however, its yield strength increases from 769 MPa to 828 MPa. With increasing carbon content, the yield strength, compressive strength, and fracture strain of the material all increase simultaneously. TZC-1.5C exhibits the best overall room-temperature mechanical properties, with a yield strength of 922 MPa, an ultimate compressive strength of 2371 MPa, and a fracture strain of 31.85%. Further increasing the carbon content worsens the mechanical properties, but its ultimate compressive strength and yield strength are still higher than those of the TZC alloy, with an ultimate compressive strength of 2302 MPa and a yield strength of 859 MPa. The combined effects of solid solution strengthening and grain refinement caused by carbon content increase the overall mechanical properties of the material. Excessive C addition leads to the precipitation of Ti2AlC, which deteriorates the overall mechanical properties of the material at room temperature.

[0083] Figure 7The figures show the room temperature compression curves of the as-cast TiAl alloys in Examples 1-3. As can be seen from the figures, the fracture strength of the TZ alloy is 1608 MPa, its yield strength is 523 MPa, and its fracture strain is 24.65%. The addition of Cr brings the most significant strengthening effect, increasing the yield strength to 769 MPa, followed by Mn (729 MPa) and Fe (642 MPa). The fracture strength shows a similar trend to the yield strength, with the order being TZ-Cr (2256 MPa) > TZ-Mn (2247 MPa) > TZ-Fe (2221 MPa). Furthermore, compared with the TZ alloy, the fracture strains of the TZ-Cr, TZ-Mn, and TZ-Fe alloys are increased by 1.43 times, 2.05 times, and 1.54 times, respectively.

[0084] The differences in the mechanical properties of TiAl alloys essentially stem from their different fracture modes. The alloys can exhibit intergranular brittle fracture, transgranular cleavage fracture, or a mixed fracture mode. Changes in the fracture path directly determine the crack propagation resistance and energy release rate, thus significantly affecting strength, plasticity, and toughness. By analyzing the differences in fracture modes, the intrinsic relationship between the microscopic mechanism and the differences in macroscopic mechanical properties is ultimately clarified.

[0085] Figure 8The figures show the crack propagation and fracture morphology of the as-cast TiAl alloys in Examples 1 and 2, where (a) and (b) are TZC-0.5C, (c) is TZC-1C, (d) and (e) are TZC-1.5C, and (f) is TZC-2C. As can be seen from the figures, (a) and (b) are TZC-0.5C cracks and their high-magnification morphology. The cracks mainly initiate and propagate along the lamellar layers, with a relatively smooth propagation path. (c) shows the TZC-1C crack morphology. The cracks mainly initiate along the lamellar layers and then propagate through a mixed fracture mode, forming a serrated crack. (d) shows the TZC-1.5C crack. The cracks initiate in the middle of the lamellar layers, propagate to another grain, and then fracture along the lamellar layers, with a significantly tortuous propagation path. (e) shows the high-magnification morphology of the TZC-1.5C crack. Crack initiation is observed at the lamellar interface, while Ti2AlC is distributed within the lamellar clusters at a certain angle to the lamellar orientation. When the crack propagates to Ti2AlC, it extends along this phase. (f) shows the morphology of the TZC-2C crack. Crack initiation occurs at the lamellar interface, and propagation mainly occurs through the lamellar interface. Ti2AlC accelerates crack propagation, and the crack path is relatively smooth. Carbon (C) as a solid solution atom increases lattice distortion, leading to improved compressive strength. Increased C content leads to an increase in the content of the bulk γ phase and effectively refines the average lamellar cluster size. Therefore, the material's plasticity increases with increasing C content. The addition of high C content induces the precipitation of the Ti2AlC phase. The precipitation of a small amount of Ti2AlC can significantly refine the average lamellar cluster size, effectively improving the material's strength and plasticity. However, the precipitation of excessive carbides will accelerate the propagation of cracks in the material during loading, thus reducing the strength and plasticity of TZC-2C.

[0086] The toughening mechanism is as follows: The strengthening mechanism of carbon (C) in TZC alloys mainly includes solid solution strengthening and lamellar refinement. Low-content C occupies octahedral interstitial sites in the γ-TiAl and α2-Ti3Al lattices as interstitial solid solutions, inducing lattice distortion, effectively pinning dislocations, and increasing yield strength. Simultaneously, C atoms segregate in stacking fault regions, reducing stacking fault energy, promoting hyperdislocation decomposition, and transforming screw dislocations into edge dislocations, increasing the difficulty of cross-slip, promoting coordinated deformation of multi-slip systems, and avoiding brittle fracture caused by early stress concentration. Lamellar refinement follows the Hall-Petch relationship; reducing lamellar size shortens dislocation slip length, making plastic deformation more uniformly distributed, delaying crack initiation, and achieving a synergistic improvement in strength and plasticity. High-content C, exceeding the solid solution limit, precipitates needle-like Ti2AlC carbides, which can generate second-phase strengthening, but consumes dissolved C atoms, weakening the solid solution strengthening effect. Ti₂AlC is a brittle phase. Its elongated morphology leads to severe stress concentration at the tips, becoming a preferential source for crack initiation and a rapid propagation channel. Transgranular cleavage fracture occurs within the carbides, accelerating macroscopic crack propagation and resulting in a decrease in the alloy's strength and ductility. Therefore, adding an appropriate amount of C can improve performance through the synergistic effect of solid solution strengthening and grain refinement, while excessive addition will worsen the mechanical behavior due to carbide precipitation.

[0087] Figure 9 The compression creep properties of different as-cast TiAl alloys in the examples are shown, where (a) is the creep strain-time curve and (b) is the corresponding creep rate-time curve. Figure 9 The compressive creep properties of as-cast TiAl alloys tested at 850 °C and 250 MPa for 100 hours are shown in (a). In (a), the TZC alloy exhibits the highest creep strain among all studied alloys, approximately 38.9%. In contrast, the addition of carbon effectively improves the creep resistance of the material. The highest creep deformations were TZC-0.5C (6.41%), TZC-1C (6.66%), and TZC-1.5C (10.17%). (b) reveals two distinct processes in the creep rate-time curve, referred to as the initial creep stage and the secondary creep stage. The average creep rate of the TZC alloy is 1.701 × 10⁻⁶. -6 s -1 The TZC-0.5C strain had the lowest average creep rate, at 1.711 × 10⁻⁶. -7 s -1 The creep rate of TZC-1C is slightly higher than that of TZC-0.5C, at 1.771 × 10⁻⁶. -7 s -1 The maximum value of TZC-1.5C in the addition of C element is 2.807 × 10⁻⁶. -7 s -1Carbon alloying effectively improves the creep resistance of materials, resulting in less deformation during creep compared to TZC matrix alloys. For TZC-1.5C alloys, the creep rate initially decreases and then increases, while a more pronounced steady-state creep stage is observed in TZC-0.5 / 1C alloys. Carbon addition reaches the lowest creep rate around 25 hours, similar to TZC alloys. After reaching the slowest creep rate, TZC-0.5 / 1C alloys maintain this creep rate and undergo slow deformation, while the creep rate of TZC-1.5C alloys increases slowly. Compared to TZC alloys, carbon addition reduces the rate of the initial creep stage, maintains steady-state creep in the secondary creep stage of low-carbon alloys, and suppresses the occurrence of secondary creep in TZC-1.5C alloys.

[0088] Figure 10 To compare the compressive creep properties of different as-cast TiAl alloys in the examples, (a) shows the creep strain-time curve, and (b) shows the corresponding creep rate-time curve. As can be seen from the figures, under the conditions of 850℃ / 250MPa, all three alloys underwent initial creep and accelerated creep stages before entering secondary decelerating creep without fracture. The TZ-Cr alloy exhibited the best creep resistance, with the lowest creep strain (38.9%) and the lowest creep rate (7.13×10¹⁰). -7 s -1 The creep rate is comparable to that of TZ alloy, but the creep rate increases slowly during the accelerated creep phase, reaching the maximum creep rate (1.77 × 10⁻⁶) in about 94 hours. -6 s -1 The TZ-Mn alloy exhibits the worst creep resistance, with a creep strain of 48.1% and a minimum creep rate of 1.26 × 10⁻⁶. -6 s -1 The creep rate was significantly higher than the previous two, reaching its minimum creep rate in just 12 hours and its maximum creep rate (4.15 × 10⁻⁶) in 36 hours. -6 s -1 The secondary creep stage is the most significant. Secondary creep is caused by the combined effects of increased cross-sectional area after sample deformation reducing effective stress, work hardening, and pinning of precipitated relative dislocations, which lead to a decrease in the creep rate.

[0089] Figure 11 The figures show the BSE morphology of different as-cast TiAl alloys after creep experiments in the examples. (a) and (b) represent TZC-0.5C, (c) and (d) represent TZC-1C, and (e) and (f) represent TZC-1.5C. As can be seen from the figures, (a) and (b) represent the morphology of TZC-0.5C after creep, showing mainly α2 / γ lamellars and γ... b Phase, and a small amount of B2 phase. Among them, γ... bThe phase precipitates a partially dispersed needle-like C14-TiCr2 phase, while the lamellar clusters show no significant deformation and no other phases precipitate. (c) and (d) show the corresponding morphology after TZC-1C creep. Further addition of C does not affect the microstructure composition after creep, but it can be seen that the increased C content leads to more C14-TiCr2 phase precipitation in the massive γ phase, while a small amount of precipitated phase is present in the lamellar clusters, distributed along the lamellar interface. (e) and (f) show the corresponding morphology after TZC-1.5C creep. It can be seen that after creep, the main phases are α2 / γ and massive γ, with a small amount of B2 and Ti2AlC phases. The increase in C content shows a further increase in the content of needle-like C14-TiCr2 phase in the massive γ phase, while a large number of fine precipitated phases are present in the lamellar clusters. In addition, the creep process does not affect the morphology of the Ti2AlC phase, which still exhibits a needle-like structure.

[0090] Figure 12 The images show the BSE morphology of different as-cast TiAl alloys after creep experiments in the comparative examples. (a)-(c) represent TZ, (d)-(f) represent TZ-Cr, and (g)-(i) represent TZ-Mn. In (a)-(c), the creep-treated samples show a significant microstructural evolution compared to the initial microstructure. For the TZ alloy, some α2 / γ lamellae exhibit significant bending due to prolonged thermal stress. Meanwhile, the Zr-rich bulk γ phase, characterized by its grayish-white color, remains relatively intact. Notably, numerous fine equiaxed grains are observed at the grain boundaries, primarily due to dynamic recrystallization (DRX) during creep. Furthermore, a small number of recrystallized grains are present within the lamellar structure, especially in regions adjacent to the bulk γ phase. In (d)-(f), the TZ-Cr alloy, in addition to lamellar bending, exhibits not only high-density recrystallized grains at grain boundaries but also a large number of recrystallized grains within the lamellae. Notably, fine, bright white particles precipitated within the Zr-rich bulk γ phase, while rod-shaped precipitates were observed within the lamellar structure, exhibiting a parallel orientation to the α2 / γ laths. Higher density precipitates were observed near the bulk γ phase, both within the lamellar structure and at grain boundaries. However, several micropores were observed around the precipitates near the bulk γ phase. (g)-(i) show a more significant microstructural evolution in the TZ-Mn alloy, with higher density bright white precipitates observed in the bulk γ phase and its adjacent regions. Similarly, relatively large pores were observed around the precipitates near the bulk γ phase. Based on these observations, it can be inferred that all three alloys underwent dynamic recrystallization (DRX) behavior during compression creep, with the TZ-Mn alloy exhibiting a more pronounced DRX. Furthermore, significant precipitates formed in both the TZ-Cr and TZ-Mn alloys, with the latter having a higher volume fraction.

[0091] Figure 13These are TEM images of the microstructure of TZC-1C after creep testing, where (a) shows the morphology at the boundary of lamellar clusters, and (b) shows the γ-ray structure. b The precipitation of needle-like C14 phase in the phase, (c) shows the morphology of Ti3AlC in lamellar clusters, and (d) shows the γ-phase precipitation. b The precipitates in (d) are shown in (e)-(g), and the diffraction results of different phases in (d) are shown in (h)-(l). As can be seen from the figures, (a) shows the interface between the lamellar clusters and the bulk γ phase after creep of the TZC-1C alloy. The bulk γ phase recrystallizes, while the lamellar clusters retain their original structure relatively well. In (b), recrystallization occurs in the bulk γ region, with a large amount of TiCr2 phase precipitating in needle-like shapes. (c) shows the internal morphology of the lamellar clusters. The lamellar clusters dissolve, the lamellar γ phase coarsens, and some lamellars precipitate Ti3AlC phase, distributed along the lamellar interface. (d)-(g) show the precipitation morphology of the C14 phase in the bulk γ phase and the corresponding phase diffraction spots. It can be seen that the C14 phase and Ti3AlC precipitate in the bulk γ phase in needle-like shapes. (h)-(l) are the EDS energy spectra corresponding to (d). Under the redistribution of supersaturated elements in the bulk γ, TiCr2 and Ti3AlC phases are precipitated. Among them, TiCr2 is formed in the Cr and Zr enriched regions, while Ti3AlC is formed in the Ti and C enriched regions.

[0092] Figure 14 The images show the TEM images of the microstructure of the TZ-Cr alloy after creep testing. (a) shows the bent lamellar structure, (b) shows the fractured α2 laths, (c) shows the diffraction results of the residual α2 phase, (d) shows the precipitates around the lamellar boundaries, (e) shows the precipitates in the recrystallization region, (f) shows the precipitates around the blocky γ phase, and (g)-(j) are the elemental distribution maps for (d). The images show significant microstructural degradation observed in the layered structure, characterized by lamellar bending, α2 lath fragmentation, γ lath spheroidization, and dynamic recrystallization (DRX) of γ grains. Figure 14 (a)-(c)). Notably, a complex phase transition occurred at the boundaries of adjacent lamellar layers of the original bulk γ ( Figure 14 (d)). Apart from recrystallized γ grains and residual α2 phase, according to the elemental distribution diagram ( Figure 14 (g)-(j)) also identified three different types of precipitates. The first type is rich in Cr and Zr, but lacks Ti relative to the γ matrix, and exists in the dynamic recrystallization (DRX) region around the grain boundaries, and is abundant in this region ( Figure 14 (e)). Furthermore, this precipitate phase is also distributed in the dynamic recrystallization (DRX) region of the lamellar structure and the region adjacent to the massive γ phase, but its quantity is less than that around the grain boundaries. Figure 14(e)-(f)). The second type is Ti-rich and Al-poor relative to the γ matrix, and is mainly distributed along lamellar boundaries. The third type is Cr-rich but lacks Ti, Al, and Zr relative to the γ matrix, and is also mainly distributed along lamellar boundaries. Compared with the first type, the latter two types have been significantly reduced in number.

[0093] Figure 15 The figures show the morphology of the B2 phase in the TZC-1C alloy after creep. (a) shows the B2 phase morphology, (b)-(f) show the elemental distribution of (a), (g) shows the morphology of Ti3AlC, (h) shows the diffraction results of the B2 phase, and (i) shows the diffraction results of Ti3AlC. As can be seen from the figures, (a)-(f) show the morphology and elemental distribution of the B2 phase at the interface between the blocky γ phase and the lamellar clusters. The B2 phase shows significant enrichment of Cr, and some lamellar clusters contain partial lamellar B2 phase along the interface. The presence of the B2 phase helps the lamellar clusters maintain their lamellar structure. The B2 phase contains some Ti3AlC phase, which shows significant enrichment of Ti and C elements while being depleted of Cr. (g)-(i) show the morphology of the B2 phase and Ti3AlC at the lamellar interface and the corresponding phase diffraction spots. It can be seen that B2 is oriented in the

[010] direction and Ti3AlC is oriented in the

[011] direction.

[0094] Figure 16 The images show the microstructure of the TZC-1.5C alloy after creep testing, where (a) represents the γ... b The morphology of the internal precipitates is shown in (b)-(d), which are the diffraction results of different phases in (a). (c) shows the morphology of Ti3AlC in the lamellar clusters. (e)-(i) are the elemental distribution diagrams of a. (j)-(i) show the morphology and diffraction results of the interface B2 phase. As can be seen from the figure, (a) shows the precipitate morphology of C14 phase and Ti3AlC in the bulk γ phase. C14 phase precipitates from the γ phase, while Ti3AlC is located inside the C14 phase. The diffraction spot results of each phase are shown in (b)-(d). The diffraction direction of the γ phase is

[110] , and the diffraction direction of Ti3AlC is consistent with that of γ. The C14-TiCr phase is along the

[211] direction. (e)-(i) are the element distribution diagrams corresponding to Figure (a). It can be seen that the C14 phase is mainly enriched with Zr and Cr elements, while Ti3AlC is enriched with Ti and C elements. (j) Morphology of the interface B2 phase and Ti3AlC. The phase structure is verified by diffraction spots as shown in (k) and (l). The further increase in C element content makes the B2 phase stable during the creep process.

[0095] Figure 17The figures show the internal morphology of the lamellar clusters of the TZC-1.5C alloy after creep testing. (a) shows the Ti2AlC morphology, (b)-(f) are the elemental distribution diagrams of (a), (g) shows the Ti3AlC morphology within the lamellar clusters, (h) shows the Ti3AlC diffraction results, and (i) shows the γ-phase diffraction results. As can be seen from the figures, (a)-(f) show the corresponding morphology and elemental distribution of Ti2AlC within the lamellar clusters. During creep, some TiCr2 phase precipitates along the Ti2AlC phase interface, and Zr and Cr elements enrich and transform along the interface to form TiCr2. (g) shows the morphology of the lamellar fracture region within the lamellar clusters. The lamellar clusters show significant coarsening, with some Ti3AlC precipitating inside the γ-phase of the lamellar clusters, while a large amount of Ti3AlC is distributed at the fracture interface. Excessive C content leads to a large precipitation of Ti3AlC within the lamellar clusters, which may cause uneven deformation within the lamellar clusters and result in fracture. The diffraction results show that the diffraction direction corresponding to Ti3AlC is... .

[0096] Finally, the morphology and performance characterization of TZF-0.5C and TZM-0.5C in Example 3 are as follows: Figure 18 and Figure 19 As shown.

[0097] Figure 18 The figures show the initial microstructure of the as-cast TiAl alloy in Example 3, where (a) and (b) are TZM-0.5C, and (c) and (d) are TZF-0.5C. As can be seen from the figures, the TZM-0.5C alloy has a near-lamellar structure, mainly composed of α2 / γ lamellar clusters and blocky γ phases. The content of the blocky γ phase is increased to 18.5% compared to the TZ-Mn alloy. A small amount of B2 phase is present at the interface between the lamellar clusters and the blocky γ phase, and its morphology and distribution are consistent with the B2 phase characteristics in TZM. The content of the blocky γ phase in the TZF-0.5C alloy remains basically unchanged compared to the TZ-Fe alloy, but the honeycomb τ3 phase increases to 6.2%.

[0098] Figure 19 The figure shows the room temperature compression curves of the as-cast TiAl alloy in Example 3. As can be seen from the figure, the fracture strength of the TZM-0.5C alloy is 1547 MPa, its yield strength is 511 MPa, and its fracture strain is 27.48%; the fracture strength of the TZF-0.5C alloy is 1521 MPa, its yield strength is 433 MPa, and its fracture strain is 23.44%. The TZC-0.5C alloy has the highest fracture strength (1669 MPa), followed by Mn (1547 MPa) and Fe (1521 MPa). The yield strength shows a similar trend to the fracture strength, with the order being TZC-0.5C (828 MPa) > TZM-0.5C (511 MPa) > TZF-0.5C (433 MPa).

[0099] The evolutionary patterns of microstructures are as follows: During creep in TiAl alloys, stress and dislocations accelerate Cr diffusion, causing TiCr2 to precipitate in a network pattern along the grain boundaries of the bulk γ phase. This brittle phase leads to stress concentration at the interface, and its interface is prone to inducing void nucleation, accelerating the accumulation of creep damage, significantly shortening the alloy's creep life, and adversely affecting its performance.

[0100] Ti₂AlC, as a MAX phase ceramic material, possesses a unique layered crystal structure with a semi-coherent interface between it and the γ-TiAl matrix. At this interface, the atomic arrangement is relatively loose, with numerous lattice-mismatched dislocations and vacancy clusters, providing low-barrier preferred sites for TiCr₂ nucleation. During creep, the interface exhibits significant dislocation pile-up, driving the formation of TiCr₂. Furthermore, Cr elements tend to accumulate along the interface during creep, with high Cr concentrations transforming into TiCr₂.

[0101] During TZC-1 / 1.5C creep, the α2 phase, due to its high C solid solution capacity and dislocation movement lowering the nucleation energy barrier, preferentially nucleates Ti3AlC within the phase or at the interface, gradually replacing the α2 lamellae. Simultaneously, the massive γ phase, with fewer defects, faster diffusion, and Ti / C enrichment, provides a stable and homogeneous nucleation environment for Ti3AlC. Both phases jointly promote carbide precipitation, leading to elemental redistribution and microstructure evolution in the α2 phase.

[0102] The high-temperature creep mechanism is as follows: High-temperature creep in TiAl alloys is controlled by atomic diffusion; the addition of carbon (C) can significantly suppress diffusion, thereby enhancing creep resistance. C atoms dissolve into the γ and α2 phases, occupying interstitial sites and inducing lattice distortion and localized strain fields. This hinders the diffusion of vacancies and alloying elements, raising the atomic transition energy barrier. This increases the diffusion activation energy of Ti atoms in the γ phase, reducing the diffusion coefficient by more than an order of magnitude, effectively delaying creep deformation and fracture processes.

[0103] The increase of the bulk γ phase deteriorates the creep properties of TiAl alloys. Its internal dislocations readily slip and climb, disrupting the continuous three-dimensional network of the lamellar structure and creating deformation-preferred pathways, leading to deformation concentration. Simultaneously, stress concentration occurs at the interface between the bulk γ phase and lamellar clusters due to the difference in creep behavior between the two phases, becoming a preferential nucleation site for creep cavities. These cavities propagate along the interface, eventually initiating intergranular creep fracture and significantly reducing the alloy's high-temperature creep life.

[0104] Excessive aggregation of coarse Ti2AlC in the TZC-1.5C alloy deteriorates creep properties. It induces the precipitation of hard and brittle TiCr2, leading to interfacial stress concentration and debonding due to modulus differences, resulting in voids. The dense carbides form a brittle framework, disrupting matrix continuity, exacerbating strain concentration, and providing rapid crack propagation paths. Furthermore, Ti2AlC consumes dissolved C atoms, weakening the solid solution strengthening effect, accelerating creep damage accumulation, and significantly shortening creep life.

[0105] In summary, carbon (C) addition refines lamellar clusters, slightly increases interlamellar spacing, and enhances the content of bulk γ-phase. High C content leads to the precipitation of Ti2AlC. The alloy's strength and ductility initially increase with increasing C, then decrease. TZC-1.5C achieves optimal room-temperature performance due to solid solution, grain refinement, and microstructure optimization. Excessive C causes carbide precipitation, thus reducing performance. Creep resistance is effectively improved with C addition, with TZC-0.5C showing the best performance. C enhances creep resistance by hindering diffusion, but TZC-1.5C exhibits metastable creep in later stages due to Ti2AlC precipitation. During creep, C helps maintain the lamellar structure. In TZC-1C, the α2 phase transforms into Ti3AlC, while TZC-1.5C contains more Ti3AlC and a small amount of TiCr2 precipitates along the Ti2AlC interface. A network of TiCr2 also precipitates in the bulk γ-phase. In conclusion, appropriate C content synergistically optimizes strength, ductility, and creep resistance, while excessive C has adverse effects due to the precipitation of brittle phases.

[0106] The various embodiments in this specification are described in a progressive manner, with each embodiment focusing on the differences from other embodiments. The same or similar parts between the various embodiments can be referred to each other.

[0107] The above description of the disclosed embodiments enables those skilled in the art to make or use the invention. Various modifications to these embodiments will be readily apparent to those skilled in the art, and the general principles defined herein may be implemented in other embodiments without departing from the spirit or scope of the invention. Therefore, the invention is not to be limited to the embodiments shown herein, but is to be accorded the widest scope consistent with the principles and novel features disclosed herein.

Claims

1. A cast TiAl alloy with optimized comprehensive mechanical properties through alloying, characterized in that, The composition, by atomic percentage, includes: 40-50% Al, 0-5% Zr, 0-3% β-stabilizing elements and 0-2% C, with the balance being Ti; The β-stabilizing element includes at least one of Cr, Mn and Fe; Among them, the amounts of Zr, β-stable elements, and C are not zero.

2. The as-cast TiAl alloy as described in claim 1, characterized in that, The composition of the as-cast TiAl alloy, by atomic percentage, comprises: 40-50% Al, 1-3% Zr, 1.5-2.5% β-stabilizing elements and 0.5-2% C, with the balance being Ti.

3. A method for preparing a cast TiAl alloy with optimized comprehensive mechanical properties through alloying, characterized in that the steps include... include: Prepare raw materials according to the composition of the as-cast TiAl alloy as described in claim 1 or 2; The raw material is placed in a vacuum electric arc furnace and then subjected to electric arc melting under argon protection. After cooling, the melting is repeated at least once to obtain the as-cast TiAl alloy.

4. The preparation method according to claim 3, characterized in that, The raw materials include titanium, aluminum, zirconium, and graphite powder, as well as at least one of chromium, manganese, and iron.

5. The preparation method according to claim 4, characterized in that, The graphite powder has a particle size of 200-400 mesh.

6. The preparation method according to claim 3, characterized in that, The current increase rate of the electric arc melting is 10-30 A / s, and the current is 350-450 A. After all the raw materials are melted, the current is kept constant for 1-3 minutes for heat preservation. After the heat preservation is completed, the furnace is cooled.

7. The preparation method according to claim 3, characterized in that, The repeated melting involves flipping the cooled ingot and melting it, and repeating this process 3-7 times.

8. The preparation method according to claim 3, characterized in that, The process before arc melting also includes a gas washing operation.

9. The preparation method according to claim 8, characterized in that, The steps of the gas washing operation include: evacuating the vacuum and then filling it with argon gas, and then evacuating the vacuum again and filling it with argon gas for protection.

10. The application of the as-cast TiAl alloy as described in claim 1 or 2 as a high-temperature structural material.

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