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Article

Microstructure and High-Temperature Compressive Properties of a Core-Shell Structure Dual-MAX-Phases-Reinforced TiAl Matrix Composite

by
Shiqiu Liu
* and
Huijun Guo
*
Advanced Materials Additive Manufacturing Innovation Research Center, School of Engineering, Hangzhou City University, Hangzhou 310015, China
*
Authors to whom correspondence should be addressed.
Crystals 2025, 15(4), 363; https://doi.org/10.3390/cryst15040363
Submission received: 17 March 2025 / Revised: 10 April 2025 / Accepted: 14 April 2025 / Published: 16 April 2025
(This article belongs to the Section Crystalline Metals and Alloys)
Browse Figures
Versions Notes

Abstract

:
As an advanced high-temperature structural material, TiAl alloy, is often used in the manufacturing of hot-end components of aviation and aerospace engines. However, it is difficult to increase the strength at high temperature, which limits its wider application. Adopting composite material technology is one of the effective ways to improve the comprehensive mechanical properties of TiAl alloy. In this work, by adding 3 wt.% SiC micro-particles to Ti-47.5Al-7Nb-0.4W-0.1B (at.%) pre-alloyed powder, a core-shell structure dual-MAX-phase high-temperature strengthened TiAl matrix composite (also known as TiAl-SiC composite) was prepared by combining powder metallurgy and hot forging. The microstructure and high-temperature compressive properties of the prepared TiAl-SiC composites were studied and compared with TiAl alloy prepared by the same process, and the microstructural characteristics of the TiAl-SiC composite and its microstructure evolution during processing were revealed. The results show that the matrix of as-sintered TiAl-SiC composites was mainly composed of γ phase and a small amount of Ti2AlC particles, while the reinforcement phase was a dual-MAX-phase core-shell structure, which was mainly composed of core Ti2AlC phase, shell Ti3SiC2 phase, and small Ti2AlC particles distributed in the outer layer. After hot forging, the microstructure of TiAl-SiC composite became more compact, finer, and more uniform; the phase composition was almost not changed, but the content of Ti2AlC, Ti3SiC2, and TiB2 phases increased significantly; the content of C in each constituent phase decreased obviously, and a granular Si-rich phase was generated in the core of the reinforcement phase. The yield strength of the as-forged TiAl-SiC composite was significantly higher than that of the as-forged TiAl alloy at temperature higher than 859 °C. This is because the core-shell structure dual MAX phases can effectively reduce the softening rate of TiAl alloy in the range of 800–900 °C, thus playing a strengthening role and increasing the service temperature of TiAl alloy.

1. Introduction

Due to its low density (~4.0 g/cm3), high specific strength, high specific modulus, and excellent high-temperature oxidation resistance and creep resistance [1,2,3,4,5,6,7,8], TiAl alloy is considered to be a new generation of high-temperature structural material that is most promising to partly replace Ni-based alloys. After years of research and development, TiAl alloy has achieved initial industrial production and application in aviation engine and other manufacturing fields. In 2006, GE announced that Ti-48Al-2Cr-2Nb alloy (48-2-2) blades processed by investment casting method were successfully applied to 6–7 low-pressure turbine blades of GEnxTM aero-engines, which is the first time that TiAl alloy has been applied to the rotating parts of a commercial aero-engine [9]. In addition, SNECMA from France has also applied cast 48-2-2 alloy to the low-pressure turbine blades of the LEAPTM series aero-engines, which will replace CFM56 engines and serve in advanced aircrafts, including Boeing 737, Airbus A320neo, and COMAC C919. Recently, GE successfully applied TiAl alloy low-pressure turbine blades prepared by additive manufacturing technology to the globally largest commercial aviation engine, the GE9X, which powers the Boeing 777X airliner [10].
However, the insufficiency of its high-temperature strength has limited the large-scale application of TiAl alloys [11]. According to previous studies, fabricating a composite is one of the effective ways to improve the comprehensive mechanical properties of TiAl alloys [12,13]. TiB2 [14,15], TiC [16,17], SiC [18], Al2O3 [19,20], Ti5Si3 [21,22], and other ceramic particles are common binary hard-ceramic phases with good high-temperature chemical stability and compatibility with the TiAl matrix. However, the intrinsic brittleness of the binary ceramic phase sacrifices the plasticity of the composite to a certain extent, thus limiting the improvement of its overall performance. It is worth noting that the ternary layered compound MAX phase can be expressed by Mn+1AXn (where M represents the transition metal element, A represents the IIIA or IVA group element, and X is the C or N element). Due to its special crystal structure, MAX has the dual characteristics of metal and ceramic: it has thermal conductivity, electrical conductivity, and machinability similar to metal and also has a high melting point, high hardness, high modulus, high oxidation resistance, and thermal stability similar to ceramics [23,24,25,26]. Introducing a MAX phase into TiAl alloy by an in situ synthesis method can significantly refine the solidified or sintered microstructure of the alloy, and the secondary precipitation of a MAX phase generated by solid-state phase transformation can greatly improve the thermal stability of the α2/γ layer in the TiAl alloy matrix, so it has received wide attention [27].
In recent years, studies on the high-temperature mechanical properties of MAX-reinforced TiAl matrix composites have increases, and it has been found that MAX-reinforced TiAl matrix composites have excellent high-temperature mechanical properties [28,29,30,31,32,33,34,35,36]. Zhou et al. [28] studied a Ti2AlN/TiAl composite (a mixture of Al and nitrided Ti powders with a mass ratio of 48:27) and found that Ti2AlN distributed in a network at the grain boundary can significantly improve the high-temperature mechanical properties of the TiAl composite. The compressive strength of the composite at 800 °C and 1000 °C was as high as 1112.1 MPa and 687.7 MPa, respectively, and the increase was 14.1% and 61.6%, respectively, in comparison to the TiAl matrix alloy (Ti-47Al). Tao Sun [29] synthesized Ti2AlN/TiAl composites with the volume fraction of reinforcement phase being 20% and 50%, respectively, by using Ti, Al, and TiN particles through a hot-press sintering process and characterized their high-temperature compressive properties. The results showed that the strength of Ti2AlN/TiAl was 40–50% higher than that of the TiAl alloy under conditions of high-temperature high-strain rate and low-temperature low-strain rate. Wang et al. [30] prepared dual-scale Ti2AlC-reinforced TiAl matrix composite by adding multi-wall carbon nanotubes to Ti-48Al-2Cr-2Nb (at.%) pre-alloyed powder and using spark plasma sintering method, and found that when the addition of carbon nanotubes is 0.5 wt.%, the tensile strength of the prepared composite at 800 °C is 599.6 MPa, which is 28.3% higher than that of Ti-48Al-2Cr-2Nb alloy. Liu et al. [31] prepared a hierarchically heterogeneous Ti2AlC-reinforced TiAl matrix composite (the fraction of Ti2AlC is 12.4%) by adding graphene oxide nanosheets to Ti-48Al-2Cr-2Nb (at.%) pre-alloyed powder and also using spark plasma sintering method, and found that it not only has good strength and plasticity at room temperature, but also has excellent high-temperature strength, which makes the service temperature of TiAl alloy increased by about 100 °C. Li et al. [32] prepared TiB and Ti2AlC dual-phase-reinforced TiAl matrix composite by adding B4C powder to Ti-45Al-8Nb (at%) pre-alloyed powder and using a hot-pressing sintering method. It was found that when the addition of B4C was 5 wt.%, the compressive strength of the composite at 800 °C reached 799.19 MPa, which was significantly higher than that without B4C addition (723.10 MPa). In addition to studying the high-temperature compressive and tensile properties of MAX-reinforced TiAl matrix composites, Lapin et al. [33] prepared in situ (Ti2AlC + TiB)/TiAl composites by casting method and conducted a preliminary study on their creep properties. It was found that the high-temperature creep properties of the composite with coarse-grained fully lamellar structure were significantly higher than those of the Ti-47Al-5.2Nb-0.2C-0.2B alloy. In summary, the high-temperature mechanical properties of MAX-reinforced TiAl matrix composites are significantly higher than those of TiAl matrix alloys, which gives them great potential for high-temperature strengthening and toughening of TiAl alloys. However, studies on the high-temperature mechanical properties of MAX-reinforced TiAl matrix composites are not comprehensive yet, and most of them focus on single-MAX-reinforced TiAl matrix composites. Considering that dual MAX phases are expected to achieve compound strengthening of TiAl alloy, it is necessary to study the high-temperature mechanical properties of dual-MAX-phases-reinforced TiAl matrix composite.
Powder metallurgy method can obtain uniform and fine microstructures [37,38,39,40,41] and is one of the common methods for preparing composite materials. In this research, by adding SiC powder to Ti-47.5Al-7Nb-0.4W-0.1B (at.%) pre-alloyed powder, a core-shell structure dual-MAX-phases-reinforced TiAl matrix composite was prepared by powder metallurgy method of mechanical mixing, hot-pressing sintering, and hot forging. The microstructure of the prepared dual-MAX-reinforced TiAl matrix composites in different processing states was characterized, and the high-temperature compressive properties of the as-forged composite were studied. The microstructure evolution of the dual-MAX-reinforced TiAl matrix composite during hot processing was emphatically revealed, which lays a foundation for the research and development of novel TiAl matrix composites with excellent high-temperature strength.

2. Experimental

2.1. Materials Preparation

In this research, Ti-47.5Al-7Nb-0.4W-0.1B (at.%) pre-alloyed powder and SiC powder were used as raw materials to prepare dual-MAX-phases-reinforced TiAl matrix composites. To effectively improve the high-temperature strength of TiAl alloy while not obviously deteriorating its plasticity, on the basis of previous research [42,43], 3 wt.% SiC was added to TiAl alloy in this research. The preparation method was divided into three steps: firstly, 3 wt.% SiC powder and TiAl powder were evenly mixed by ball milling. The detailed mixing process was as follows: A planetary ball mill was used to mix powder in an argon atmosphere; stainless steel grinding balls with different diameters (with diameters of 10 mm, 5 mm, and 3 mm and a corresponding quantity ratio of 1:1:4) were added according to the ball material ratio of 1:5. In order to prevent the powder temperature in the mixing process from being too high, an intermittent ball-milling mixing method was adopted, in which the ball mill was stopped for 5 min every 10 min. The ball mill speed was set at 150 r/min, and the powder was mixed for 7.5 h. Secondly, hot-press sintering was utilized, in which the composite green body was prepared in argon atmosphere by hot-pressing sintering. The mixed TiAl-SiC composite powder was loaded into a graphite mold and vibrated to prevent the powder from ejecting from the mold during vacuum pumping. The hot-pressing sintering process was as follows: the heating rate was 10 °C/min, the hot-pressing sintering temperature was 1250 °C, the pressure was 20 MPa, and the heat and pressure holding time was 2 h, after which the furnace was cooled. The size of the as-sintered TiAl-SiC composite ingot was Ф50 mm × 56.7 mm. Finally, in order to further improve the densification and mechanical properties of the as-sintered TiAl-SiC composite, the canned hot-forging process was carried out, in which 304 stainless steel was used as the canned material, and the wall thickness of the can was 8~12 mm. The canned body was heated to 1250 °C at a heating rate of 10 °C/min by muffle furnace and held at this temperature for 1 h. A 300 ton hydraulic press was used for hot forging of the canned body. Thermal insulation cotton of aluminum silicate (10 mm in thickness) and graphite paper (0.5 mm in thickness) were placed between the canned body and the indenter or anvil of the hydraulic press to provide insulation and lubrication. The reduction due to hot forging was 50%, which was divided into two steps of hot forging, in which each reduction was 25%; the canned body was held at 1250 °C for 0.5 h between the two steps of hot forging. After the hot forging, the canned body was annealed at 850 °C for 2 h, and then, the furnace was cooled. The preparation process of TiAl alloy was the same as that of TiAl-SiC composite, excluding the powder mixing process.

2.2. Examination of Powder Size and Material Density

A laser particle size analyzer (Mastersizer 3000, Malvern Instruments Ltd., Malvern, UK) was used to analyze the particle size distribution of two adopted raw material powders; during the examination, the powder samples were mixed with a dispersing medium of pure water and ultrasonically dispersed to prevent agglomeration. The density of TiAl alloy and its composites was measured according to the Archimedes principle. The density test specimen was a small cube with a size of 10 mm × 10 mm × 10 mm; before the test, it was necessary to use anhydrous ethanol for ultrasonic cleaning of the test specimen to reduce the influence of impurities such as oil on the testing results. In order to reduce the error, each specimen was measured three times, and the average value was taken.

2.3. Microstructure Analysis

The phase composition of the raw powders, TiAl alloy, and its composites was analyzed by X-ray diffractometry (XRD, D/Max 2500). The test parameters used for common phase identification were as follows: continuous scanning with a 2θ range of 10~90° and a scanning speed of 6 °/min; the test specimen was ground flat until there was no obvious scratch. MDI Jade 6 software was used for phase identification analysis of XRD patterns. The microstructure and powder morphology were observed by a Quanta 250 FEG field emission scanning electron microscope (SEM). The chemical composition of each constituent phase of TiAl matrix composites was analyzed by an EDAX energy dispersive spectrometer (EDS). The EDS examination of chemical element content complied with the ISO 22309:2012 standard in this research. According to this standard, the C content was directly quantified according to the X-ray intensity, but the examination error was relatively large than other elements whose atomic number is not less than 11. Therefore, the EDS examination results of C were semi-quantified in this research. For the accuracy of measurement, at least three EDS examinations were performed for each phase to calculate the average chemical composition; the measurement location was randomly selected. The chemical element mapping analysis of the prepared composites was conducted by a JAX-8030 electron probe microanalyzer (EPMA). All microstructural dimensions were measured using the Gatan Microscopy Suite software (version 2. 32. 888. 0).

2.4. High-Temperature Compressive Properties Test

A thermal mechanical simulator (Gleeble 3500) was used to measure the high-temperature compressive properties of the as-forged TiAl-SiC composite and TiAl alloy. The high-temperature compression specimens (Φ6 mm × 9 mm in size) were cut along the hot-forging direction of the forging blank by wire cutting, and their surface was polished as necessary. The test temperature was 700 °C, 800 °C, 900 °C, and 1000 °C, respectively, which was held for 2 min after the temperature was reached; the heating rate was 5 °C/s, and then, the test was carried out. The strain rate was 10−3 s−1, and the reduction (true strain) was 0.3. After reaching the required deformation, water quenching was carried out to retain the high-temperature microstructure.

3. Results and Discussion

3.1. Raw Material Characterization

TiAl-SiC composites were prepared by powder metallurgy process. The raw materials used to prepare the composite included Ti-47.5Al-7Nb-0.4W-0.1B (at.%) pre-alloyed powder and SiC powder. The sphericity degree of the TiAl pre-alloyed powder was excellent, the powder surface was smooth, and the difference in powder size was relatively large; its morphology is shown in Figure 1a. SiC powder particles had irregular morphology, and the difference in powder size was relatively small, as shown in Figure 2a. The particle size distribution analysis results of TiAl pre-alloyed powder and SiC powder are shown in Figure 1b and Figure 2b, respectively. The results show that the particle size of TiAl powder was mainly distributed in the range of 50~125 μm, with a median particle size of 90.7 μm. The particle size of SiC powder was mainly distributed in the range of 3.5–11 μm, and the median particle size was 7.35 μm. As seen from Figure 1c, the main constituent phase of TiAl powder was α2, followed by γ and a small amount of B2 phase. It can be seen from Figure 2d that the applied SiC powder was not pure SiC powder and mainly consisted of the three phases of SiC, C, and Si. Among them, SiC included hexagonal SiC and cubic SiC, and hexagonal SiC accounted for the vast majority. According to the intensity and quantity of X-ray diffraction peaks, it can be inferred that the SiC phase proportion of the applied SiC powder can reach more than 80%. The EDS analysis results of SiC powder indicated (see Figure 2c) that the content of C element in SiC powder was higher than that of Si element.

3.2. Microstructure of As-Sintered TiAl-SiC Composite

It is indicated by Figure 3 that the as-sintered TiAl-SiC composite was mainly composed of the three phases of γ-TiAl, Ti2AlC, and Ti3SiC2. According to the number and intensity of diffraction peaks, the volume fraction of γ phase in the as-sintered TiAl-SiC composite was the largest, followed by Ti2AlC and Ti3SiC2 phases. Among of them, the Ti3SiC2 MAX phase is extremely necessary in the future as the most promising material for protecting containers of fuel elements in nuclear power, and materials based on the MAX-phase Ti3SiC2 can also become indispensable when operating in high-strength friction units up to temperatures up to 1000 °C [44,45,46]. In comparison, the as-sintered TiAl alloy was mainly composed of γ, α2-Ti3Al, and B2 phases. Obviously, the phase composition of TiAl alloy changed significantly with the addition of SiC particles, and the volume fraction of α2 and B2 phases was significantly reduced or even disappeared, which may be related to the increase in C and Si content of the TiAl matrix.
As shown in Figure 4, the as-sintered TiAl-SiC composite was basically densified, the contours of the TiAl pre-alloyed powder were still clearly visible, and most of them were of a regular circle shape. In addition, the dark reinforcement phase was irregular and mainly distributed in the gap position formed by the accumulation of different TiAl pre-alloyed particles (see Figure 4a). As seen from Figure 4d, the as-sintered TiAl alloy was fully densified, and the outline of the TiAl pre-alloyed powder could be roughly distinguished because a discontinuous gray phase formed on the outermost layer of the TiAl pre-alloyed powder particle. These gray phases were α2 phase, which resulted from oxidation and Al deficiency on the surface of TiAl pre-alloyed powder, and the high oxygen content and low Al content are conducive to the formation of α2 phase. In the as-sintered TiAl alloy, some of the original powder particle contours were of a regular circle shape, but many of them were of irregular round shape and polygonal. This indicates that during the hot-pressing sintering process of TiAl alloy, a considerable part of the TiAl pre-alloyed powders underwent obvious plastic deformation, which filled the gap between the powders and densified the prepared TiAl alloy. Compared with the as-sintered TiAl alloy, the densification mechanisms of TiAl-SiC composite are more complicated due to the addition of SiC powder. In addition to the deformation of TiAl pre-alloyed powders and their interdiffusion of chemical elements, it also includes the filling of the gap as well as the diffusion and reaction of elements between the SiC powders and the TiAl pre-alloyed powders. The above analysis shows that the hot-pressing sintering process at 1250 °C/20 MPa/2 h can basically meet the densification requirement of powder metallurgical Ti-47.5Al-7Nb-0.4W-0.1B alloy and Ti-47.5 Al-7Nb-0.4W-0.1B-3 wt.% SiC composite.
The matrix microstructure of the as-sintered TiAl-SiC composite was mainly composed of three parts (see Figure 4b), namely the dark matrix phase, the dispersed black particle phase, and the finer white granular phase. In comparison, the as-sintered TiAl alloy exhibited near-γ microstructure (see Figure 4e), in which the majority of the dark phase was γ phase, the irregular gray phase was α2, and the dispersed fine white particles were TiB2 phase. It was revealed that after the addition of SiC particles, the gray α2 phase in the TiAl alloy disappeared, and a new black particle phase was formed. During the hot-pressing sintering process, the added SiC particles reacted with the TiAl pre-alloyed powders to form a multi-phase microstructure (see Figure 4c). The multi-phase microstructure had a core-shell structure and was mainly composed of three parts, namely the white phase on the outermost layer, the dispersed black particles in the white phase, and the dark phase in the interior. Based on the XRD results, it can be inferred that the matrix phase of the TiAl-SiC composite was mainly γ-TiAl phase, while the reinforcement phase was a core-shell structure composed of Ti2AlC and Ti3SiC2 dual MAX phases.
To further determine the phase composition and chemical element distribution of the as-sintered TiAl-SiC composite and TiAl alloy, both EDS and EPMA analyses were carried out. By EDS analysis, it was found the chemical composition of γ phase in the matrix of the as-sintered TiAl-SiC composite was significantly different from that of the as-sintered TiAl alloy (see Table 1). The difference in chemical composition between them was mainly in the content of Ti, Al, and C: compared with γ phase in the TiAl alloy, the γ phase in the TiAl-SiC composite contained C element and less Al and Ti elements. The C element mainly came from the diffusion of C element from SiC particles to TiAl pre-alloyed powder during hot-pressing sintering. A study indicated [47] that the solubility of C in the γ phase at room temperature is very low, about 300 ppm. However, EDS results showed that the content of C in the γ phase of the as-sintered TiAl-SiC composite matrix was as high as 13.43 at.%, which was much higher than the solubility of C in the γ phase. Consequently, it can be inferred that there are a large number of very small C-rich phases in the γ phase of the as-sintered TiAl-SiC composite matrix. Because these C-rich phases are very small (their size maybe in submicron scale or even nanoscale) and have a similar backscattered contrast to the γ phase, it is difficult to distinguish them in low-magnification SEM pictures. By using the method of Al equivalent, the content of Nb, W, and Si was converted into Al content so as to obtain a ternary composition of Ti-Al-C, and then, the ternary phase diagram of Ti-Al-C was used to further speculate the type of the C-rich phase in the matrix. According to a reference [48], the ternary equivalent composition of Ti-Al-C corresponding to label A can be calculated as Ti-39.55Al-13.43C. Then, according to the isothermal section for the Ti-Al-C system at room temperature, provided in the reference [49], the phase composition of the equivalent chemical composition is γ-TiAl + Ti2AlC + a small amount of Ti3Al5. Therefore, it can be inferred that the C-rich phase in the as-sintered TiAl-SiC matrix is Ti2AlC phase. In addition, there were a small number of dispersed black particles in the matrix of the composite (see Figure 4b). Through EDS analysis (see the corresponding results of label B in Table 1), it was found that they were a kind of C-rich phase (C content up to 35.64 at.%), and the content of C and Al was not much different. Combined with the XRD result and its contrasting characteristics in BSE mode, it can be inferred as Ti2AlC phase. This is consistent with the above-stated result obtained by using the Ti-Al-C ternary phase diagram, indicating that the matrix of the as-sintered TiAl-SiC composites is mainly composed of γ and Ti2AlC phases. The reason why the phase labeled as B contained a high amount of Si may be that these Si-rich Ti2AlC phases are derived from extremely fine SiC particle (~1 μm in size). The fine particle size of SiC may effectively inhibit the formation of Ti3SiC2 phase, although the relevant mechanism is still unclear and requires further study in the future.
The internal dark part of the reinforcement phase of the as-sintered TiAl-SiC was analyzed by EDS (see the corresponding result of label C in Table 1). It was found that the phase was a C-rich and Si-poor phase with C content up to 35.57 at.% and Si content almost negligible (0.17 at.%). In addition to the relatively low content of Nb, W, and Si, its chemical composition was basically the same as that of the Ti2AlC phase in the matrix, so it can be determined that the internal phase is exactly Ti2AlC phase. The relatively low content of Nb and W may be related to the difficulty of diffusion of the two elements with their large atomic number into the interior of the reinforcement phase, while the relatively low content of Si may be related to the enrichment of Si element in the outer layer of the reinforcement phase during hot-pressing sintering. Through EDS analysis of the “shell” (white phase) of the reinforcement phase (see the corresponding results of label D in Table 1), it was found that this phase was a kind of Si-rich and C-rich phase, and its Si content was the highest among all the constituent phases (16.70 at.%), while its C content (21.41 at.%) was higher than that of Si content. According to its composition characteristics and the above-stated analysis results of microstructure and phase composition, it can be determined that this phase is Ti3SiC2 phase. There was also a fine black particle phase distributed in the outer white phase of the reinforcement phase; its contrast was similar to the interior dark phase of the reinforcement phase. Through EDS analysis of this phase (see the corresponding result of label E in Table 1), it was found that the C content of this phase was very high (31.6 at.%), which is similar to the C content of the above-mentioned Ti2AlC phase. Therefore, it can be determined that this black particle phase is also Ti2AlC phase. In conclusion, the matrix of the as-sintered TiAl-SiC composite is primarily composed of γ phase and multi-scale Ti2AlC precipitates, while the reinforcement phase is mainly composed of the interior Ti2AlC phase, the outer Ti3SiC2 phase, and fine Ti2AlC particles distributed in the outer layer.
The EPMA analysis results of the reinforcement phase of the as-sintered TiAl-SiC composite are shown in Figure 5. It is indicated by this figure that the Ti element in the composite was mainly enriched in the inner region of the reinforcement phase (see Figure 5b). The distribution of Al element is much different than that of Ti element: Al element was primarily distributed in the matrix of the composite, while in the reinforcement phase region, Al element was mainly distributed in the inner Ti2AlC phase and was rarely distributed in the outer Ti3SiC2 phase (see Figure 5c). This indicates that the outer white phase of the reinforcement phase is an Al-poor phase, while the inner phase is a relatively Al-rich phase, which is consistent with the chemical composition characteristics of Ti3SiC2 and Ti2AlC. Both Nb and W elements in the composite were mainly distributed in the outer phase of the reinforcement phase (see Figure 5d,e). This phenomenon may be caused by two reasons: On one hand, due to the small atomic diffusion coefficients of Nb and W atoms, they have difficulty diffusing into the interior of the reinforced phase during hot-pressing sintering so as to enrich in the outer layer of the reinforcement phase. On the other hand, the solubility of Nb and W elements is relatively large in the Ti3SiC2 phase and small in the Ti2AlC phase, so they tended to be distributed in the shell of the reinforcement phase. In addition, the B element was also mainly enriched in the outer phase of the reinforcement phase (see Figure 5f), which may be related to the higher solubility of B element in the outer phase than in the interior one.
In the as-sintered TiAl-SiC composite, Si was mainly distributed in the outer phase of the reinforcement phase, while C was mainly distributed in the inner phase as well as the fine particles in the outer phase of the reinforcement phase, i.e., C-rich phase in the matrix (see Figure 5g,h), which is consistent with the results of EDS analysis. This indicates that Si and C elements in SiC particles separated during hot-pressing sintering. C element generally presented a high-to-low gradient distribution from the inside of the reinforcement phase to the matrix, while Si element was mainly enriched in the outer layer of the reinforcement phase. This phenomenon may be due to the following two reasons: On one hand, because Si atoms have low solubility in the TiAl matrix, they cannot diffuse into the matrix on a large scale during hot-pressing sintering and thus accumulate in the outer layer of the reinforcement phase. On the other hand, Si atoms have a large solubility in the Ti3SiC2 phase and a low solubility in the Ti2AlC phase, so they tend to be distributed outside the reinforcement phase. According to the above EPMA analysis, it was revealed that Ti, Al, and C are primarily distributed in the Ti2AlC phase, while the other elements are mainly enriched in the outer Si-rich phase of the reinforcement phase.

3.3. Microstructure of As-Forged TiAl-SiC Composite

The microstructures of the as-forged TiAl-SiC composite and TiAl alloy are shown in Figure 6. By comparing the low-magnification microstructure of the as-sintered and as-forged TiAl-SiC composites (as shown in Figure 4a and Figure 6a, respectively), it was found that after hot forging, the composite became more densified, the contour of the original TiAl pre-alloyed powder disappeared, and the coarse reinforcement phase was crushed, becoming finer and more uniform. However, the phase composition of the matrix microstructure of TiAl-SiC composite did not change significantly after hot forging and was still mainly composed of dark matrix phase, black particle phase, and fine white granular phase (see Figure 6b). The reinforcement phase of TiAl-SiC composite was obviously broken after hot forging, but it was still not completely disintegrated and was still composed of three parts, namely the interior dark phase, the outer white phase, and the dark particle phase distributed in the outer layer.
The low-magnification microstructure of the as-forged TiAl alloy is shown in Figure 6d. Comparison with that in hot-pressing state (see Figure 4d) indicated that after hot forging, the microstructure of TiAl alloy became more uniform, and the size and volume fraction of the gray phase increased significantly. Through observing the high-magnification microstructure of the as-forged TiAl alloy (see Figure 6e), it was found that the gray phase in the low-magnification microstructure picture was α2/γ lamellar microstructure or α2 phase without eutectoid transformation. Therefore, it was revealed that the microstructure type of the as-sintered TiAl alloy changed from near γ to duplex after hot forging. In addition, there were many white rod-like TiB2 phases in the as-forged TiAl alloy, which were obviously coarsened compared with the TiB2 in the as-sintered TiAl alloy. The microstructure evolution of the as-sintered TiAl alloy after hot forging, including microstructure homogenization, change of microstructure type, and coarsening of constituent phases, is closely related to the plastic deformation and high-temperature heat treatment during hot forging. Especially, the coarsening of TiB2 particles during forging is mainly due to two reasons: On one hand, there are two heat preservation processes at 1250 °C, 1.5 h in total, during hot forging, and the long-term heat treatment at high temperature provides beneficial thermodynamic conditions for the growth of TiB2. On the other hand, deformation during hot forging can introduce large numbers of lattice defects, which can improve the internal energy of the as-forged TiAl-SiC composite or TiAl alloy as well as provide favorable diffusion paths for solute atoms, and both of these factors are beneficial to the coarsening of TiB2. The comparison in microstructure between the TiAl-SiC composite and TiAl alloy before and after hot forging shows that the matrix of TiAl-SiC composite with relatively high C content exhibits more excellent microstructure thermal stability than TiAl alloy at hot-forging temperature (1250 °C). Lots of previous research has indicated that proper C alloying exhibits beneficial effects on the microstructure thermal stability of TiAl alloys [50,51,52], which is primarily because C is not only an efficient solid solution strengthener in TiAl but also an efficient precipitation strengthener due to its fine dispersion of carbide. Both the solid–solution strengthening effect of C atoms and the precipitation strengthening effect of fine carbides make the interface migration difficult at high temperature, which improves the microstructure thermal stability of TiAl alloy. Consequently, the matrix of TiAl-SiC composite with relatively high C content exhibited more excellent microstructure thermal stability than TiAl alloy at hot-forging temperature.
In addition, by comparing the density of TiAl-SiC composite and TiAl alloy in different states (see Figure 7), it was shown that the density of the two materials did not change much before and after hot forging, indicating that the as-sintered TiAl-SiC composite and TiAl alloy prepared in this research basically reached densification. Secondly, the density of TiAl alloys in different states was about 4.2 g/cm3, while the density of TiAl-SiC composites prepared by the same process was about 4.05 g/cm3, indicating that the addition of SiC significantly reduces the density of TiAl alloy.
The XRD patterns of the as-forged TiAl-SiC composite and TiAl alloy are shown in Figure 8. It is indicated that the as-forged TiAl-SiC composite was mainly composed of γ, Ti2AlC, Ti3SiC2, TiB2, and Ti5Si3 phases. According to the intensity of diffraction peaks, the γ phase content of the as-forged TiAl-SiC composite still accounted for the majority, followed by Ti2AlC and Ti3SiC2 phases, and a small amount of TiB2 and Ti5Si3 phases existed. By comparing the XRD patterns of the two TiAl-SiC composites, it was found that there were obvious differences between them. Firstly, the diffraction peaks corresponding to Ti2AlC phase on the XRD pattern of the as-forged TiAl-SiC composite increased significantly, and the relative peak intensity also increased obviously. The increase in the number of diffraction peaks may be due to the fact that the phase was crushed in the process of hot forging; therefore, diffraction peaks from different crystal planes appeared. The increase in the relative peak intensity indicates the improvement in the content of Ti2AlC phase after hot forging. Secondly, the relative peak intensity corresponding to the Ti3SiC2 phase on the XRD pattern of the as-forged TiAl-SiC composite also increased significantly, which indicates that the content of the Ti3SiC2 phase also increased after hot forging. In addition, the diffraction peaks of TiB2 and Ti5Si3 also appeared on the XRD pattern of the as-forged composite, indicating the increase or precipitation of the two phases after hot forging. The XRD analysis basically revealed the phase evolution of the as-sintered TiAl-SiC composite after hot forging: The phase content of Ti2AlC, Ti3SiC2, and TiB2 phases in the as-sintered TiAl-SiC composite increased significantly. In addition, the Ti5Si3 phase was precipitated. However, it is difficult to distinguish Ti5Si3 phase in the BSE image of the as-forged TiAl-SiC composite due to its small phase content or very small size. The as-forged TiAl alloy was still mainly composed of the three γ, α2, and B2 phases (see Figure 8). In addition, compared with the XRD pattern of the as-sintered TiAl alloy, the diffraction peaks of B2 phase on the XRD pattern of the as-forged TiAl alloy increased significantly, and the diffraction peak of TiB2 phase appeared. This indicates that both the B2 and TiB2 phases in the as-sintered TiAl alloy were increased significantly after hot forging. The XRD analysis results show that hot-forging process is conducive to obtaining more stable phase composition for the as-sintered TiAl-SiC composite and TiAl alloy.
To further determine the constituent phase and chemical element distribution of the as-forged TiAl-SiC composite, EDS and EPMA analyses were carried out. By comparing the EDS analysis results of each phase listed in Table 1 and Table 2, it was found that the difference in chemical composition of the constituent phases between the as-forged and as-sintered materials was not much, and the most obvious difference was the C content of each constituent phase in the TiAl-SiC composites. Through detailed comparison, it was found that the C content of each constituent phase in the as-forged TiAl-SiC composite was significantly lower than that in the as-sintered composite. Specifically, the average C content of the matrix, Ti2AlC, and Ti3SiC2 phases in the as-forged TiAl-SiC composite was 11.63 at.%, 30.70 at.%, and 18.52 at.%, respectively (see Table 2). The average C content of the corresponding phases in the as-sintered composite was 13.43 at.%, 35.57 at.%, and 21.41 at.%, respectively (see Table 1). After hot forging, the C content of each constituent phase in the TiAl-SiC composite decreased mainly due to the following two reasons. Firstly, the heat treatment process included in the hot-forging process further diffused the C element and made it more uniformly distributed, and it also promoted the precipitation of C-rich phase such as Ti2AlC, thus effectively reducing the enrichment of C element in each constituent phase. Secondly, the hot-forging process effectively broke the reinforcement phase of the TiAl-SiC composite, which significantly refined the reinforcement phase and obviously increased the contact area between the reinforcement phase and the matrix, thus promoting the diffusion of C element from the reinforcement phase to the matrix. This is conducive to the uniform distribution of C element and the precipitation of C-rich phase and thus facilitates the reduction in the C content of each constituent phase. In summary, the hot-forging process can effectively break the reinforcement phase of the as-sintered TiAl-SiC composite, promote the diffusion of C and other elements, and facilitate the precipitation of C-rich phase so as to effectively reduce the C element in each constituent phase of TiAl-SiC composite and make the chemical composition more uniformly distributed.
The EPMA mapping results of the reinforcement phase in the as-forged TiAl-SiC composite are shown in Figure 9. Compared with the as-sintered TiAl-SiC composite, it was found that the distribution of elements in the reinforcement phase had not changed obviously after hot forging; namely, Ti, Al, and C were still mainly distributed in the Ti2AlC phase, while the other elements were primarily enriched in the outer Si-rich phase of the reinforcement phase. However, some elements had slight change in distribution location after hot forging; the most obvious were Si and B elements. The two elements were not only distributed in the outermost white phase of the reinforcement phase but also in its interior in granular form (see Figure 9f,g). This indicates that the Si-rich particles are formed or precipitated in the Ti2AlC phase inside the reinforcement phase after hot forging. This phenomenon is difficult to observe in the as-sintered TiAl-SiC composite because the Si element is almost completely enriched in the outermost layer of the reinforcement phase during the hot-pressing sintering process, and the solubility of Si element in the Ti2AlC phase is limited. The formation of granular Si-rich phase in the reinforcement phase after hot forging may be mainly related to the following two factors: On one hand, the hot-forging process promoted the diffusion of Si element and increased the Si content of Ti2AlC phase in the reinforcement phase (see Table 1 and Table 2), thus facilitating the precipitation of Si-rich phase. On the other hand, the hot-forging process caused the Si-rich phase in the shell of the reinforcement phase to be broken up into granular fragments; these fragments moved during the hot-forging process and entered the interior of the reinforcement phase. This indicates that hot forging can not only effectively break the coarse reinforcement phase but also promote the uniform distribution of different constituent phases.

3.4. High-Temperature Compressive Properties of TiAl-SiC Composite

Figure 10a shows the high-temperature compression true stress–strain curves of the as-forged TiAl-SiC composite at different temperatures (700~1000 °C). As shown in the figure, the elastic modulus of the as-forged TiAl-SiC composite gradually decreased with the increase in temperature. It is worth noting that the elastic modulus of the as-forged TiAl-SiC composite was almost the same at 700 °C and 800 °C. Since the elastic modulus represents the difficulty of making atoms leave their equilibrium position and is a physical quantity that measures the strength of interatomic binding force in crystal, the elastic modulus of a material mainly depends on the strength of the interatomic binding force in the constituent phase and has little relationship with the microstructure characteristics. Therefore, it can be inferred that the phase composition and phase elastic modulus of the as-forged TiAl-SiC composite can remain relatively stable below 800 °C. The compressive yield strength of the as-forged TiAl-SiC composite decreased with increasing temperature (see Figure 10a). With the increase in temperature, the plastic deformation stage of the as-forged TiAl-SiC composite also exhibited a different variation trend with true strain. When the temperature was 700 °C, the stress in the plastic deformation stage of the as-forged composite increased with the increase in strain, showing an obvious work-hardening phenomenon. When the temperature was 800 °C, the stress in the plastic deformation stage of the as-forged composite almost remained unchanged with the increase in strain, and there was no obvious work-hardening phenomenon. When the temperature rose to 900 °C or above, the stress in the plastic deformation stage of the as-forged composite decreased with the increase in strain, and the rheological softening phenomenon was obvious. This shows that the as-forged TiAl-SiC composite exhibits obvious softening started from 900 °C on the condition of uniaxial compression.
Figure 10b shows the high-temperature compression true stress–strain curves of the as-forged TiAl alloy at different temperatures (700~1000 °C). The elastic modulus of the as-forged TiAl alloy constantly decreased with the increase in temperature, which indicates that the phase composition and elastic modulus of the as-forged TiAl alloy are less stable than those of the as-forged TiAl-SiC composite at 700~800 °C. Like the as-forged TiAl-SiC composite, the yield strength of the as-forged TiAl alloy decreased with the increase in temperature. With the increase in temperature, the variation trend of the plastic deformation stage of the as-forged TiAl alloy with true strain was different from that of the as-forged TiAl-SiC composite. When the temperature was lower than 800 °C, the plastic deformation stage of the as-forged TiAl alloy presented an obvious work-hardening phenomenon, and the rheological softening phenomenon was obvious when the temperature was higher than 900 °C. This indicates that the softening temperature of the as-forged TiAl alloy under uniaxial compressive stress may be slightly higher than that of the as-forged TiAl-SiC composite. In conclusion, both the phase composition and phase modulus of the as-forged TiAl-SiC composite are more stable than those of the as-forged TiAl alloy below 800 °C, and its softening temperature under uniaxial compressive stress is slightly lower than that of the as-forged TiAl alloy. These high-temperature compressive mechanical properties of the as-forged TiAl-SiC composites are closely related to their phase composition and other microstructure characteristics. The excellent thermal stability in the structure and modulus of the reinforcement phase makes the TiAl-SiC composite maintain a relatively stable elastic modulus below 800 °C. However, the large deformation mismatch between the reinforcement phase and the γ phase leads to stress concentration during high-temperature compression deformation, which promotes dynamic recovery and recrystallization and reduces the softening temperature to some extent.
Figure 10c shows the variation curves of high-temperature compressive yield strength of the as-forged TiAl-SiC composite and TiAl alloy with temperature. It indicates that the compressive yield strength of the as-forged TiAl-SiC composite and TiAl alloy decreased with the increase in temperature. For the compressive yield strength, the as-forged TiAl-SiC composite decreased from 485.34 MPa at 700 °C to 227.58 MPa at 1000 °C, and correspondingly, the as-forged TiAl alloy decreased from 512.77 MPa at 700 °C to 206.95 MPa at 1000 °C. There was an intersection point between the variation curves of high-temperature compressive yield strength of the two materials with temperature at about 850 °C. When the temperature was higher than this temperature, the yield strength of the as-forged TiAl composite was greater than that of the as-forged TiAl alloy; when the temperature was lower than this temperature, the comparative result was opposite. The compressive yield strength of the two materials at different temperatures is shown in Table 3. It indicates that, compared with the as-forged TiAl alloy, the as-forged TiAl-SiC composite is more suitable for application in a temperature range above 850 °C. The main reason for this result is that the as-forged TiAl-SiC composite has a lower softening rate than the as-forged TiAl alloy in the range of 800~900 °C. It is thus shown that the core-shell structure dual-MAX-reinforcement phase can effectively reduce the softening rate of TiAl alloy in the range of 800~900 °C, thus strengthening the TiAl alloy.

4. Conclusions

In this research, by using TiAl pre-alloyed powder and SiC powder as raw materials, a core-shell structure dual-MAX-phases-reinforced TiAl matrix composite was prepared by a combined method of mechanical mixing, hot-pressing sintering, and hot forging. The microstructure and high-temperature compressive properties of dual-MAX-phases-reinforced TiAl matrix composites were emphatically investigated. The conclusions are as follows:
(1)
The matrix of the as-sintered TiAl-SiC composite was mainly composed of γ phase and Ti2AlC precipitates with different scales, while the reinforcement phase exhibited a core-shell structure, which was mainly composed of core Ti2AlC phase, shell Ti3SiC2 phase, and fine Ti2AlC particles distributed in the outer layer. In the as-sintered TiAl-SiC composite, Ti, Al, and C were mainly distributed in the Ti2AlC phase, while the other elements were mainly enriched in the outer Ti3SiC2 phase of the reinforcement phase;
(2)
After hot forging, the microstructure of TiAl-SiC composite became finer and more uniform, and the phase composition changed little. The phase contents of Ti2AlC, Ti3SiC2, and TiB2 increased significantly;
(3)
After hot forging, the C content of each constituent phase of TiAl-SiC composite was significantly reduced mainly because the hot-forging process can effectively promote the diffusion of C element and the precipitation of C-rich phase. In addition, a granular Si-rich phase was formed inside the reinforcement phase, which was mainly because the hot-forging process can promote the diffusion of Si element, breaking the Si-rich phase and moving it;
(4)
The compressive yield strength of the as-forged TiAl-SiC composite decreased with the increase in temperature. When the temperature was higher than 859 °C, its yield strength was greater than that of the as-forged TiAl alloy, and the comparative result was opposite when the temperature was lower than 859 °C. This is because the core-shell structure dual MAX phases can effectively reduce the softening rate of TiAl alloy in the range of 800–900 °C, thus playing a strengthening role and increasing the service temperature of TiAl alloy.
The microstructure and high-temperature compressive properties of a core-shell structure dual-MAX-phases-reinforced TiAl matrix composite were investigated preliminarily in this research. There are several aspects that require further study in the future, mainly presenting as the following two aspects: On one hand is the effect of an additive amount of SiC particles on the microstructure and mechanical properties of TiAl-SiC composite, on the basis of which the optimum additive amount of SiC can be revealed, thereby facilitating the development of a TiAl matrix composite with superior high-temperature mechanical properties. On the other hand, the high-temperature strengthening mechanism of the core-shell structure dual MAX phases for the TiAl-SiC composite also necessitates in-depth investigation, which is critical for further tailoring the core-shell structure and making full use of its high-temperature strengthening effect.

Author Contributions

Conceptualization, S.L.; methodology, S.L.; software, S.L.; validation, S.L., and H.G.; formal analysis, H.G.; investigation, S.L.; resources, S.L.; data curation, S.L.; writing—original draft preparation, S.L. and H.G.; writing—review and editing, S.L.; supervision, S.L.; project administration, S.L.; funding acquisition, S.L. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by A Project Supported by Scientific Research Fund of Zhejiang Provincial Education Department (Y202454336), A Project Supported by State Key Laboratory of Powder Metallurgy, Central South University, Changsha China (Sklpm-KF-015) and a project from Hangzhou City University (204000-581870).

Data Availability Statement

The data that support the findings of this study are available on request from the corresponding author upon reasonable request.

Acknowledgments

The authors would like to thank A Project Supported by Scientific Research Fund of Zhejiang Provincial Education Department (Y202454336) and A Project Supported by State Key Laboratory of Powder Metallurgy, Central South University, Changsha China (Sklpm-KF-015) for providing financial support.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. Characteristics of TiAl pre-alloyed powder used for fabrication of TiAl-SiC composite: (a) morphology, (b) particle size distribution of powder, and (c) XRD pattern.
Figure 1. Characteristics of TiAl pre-alloyed powder used for fabrication of TiAl-SiC composite: (a) morphology, (b) particle size distribution of powder, and (c) XRD pattern.
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Figure 2. Characteristics of SiC powder used for fabrication of TiAl-SiC composite: (a) morphology, (b) particle size distribution of powder, (c) EDS analysis result of SiC powder, and (d) XRD pattern, where h-SiC represents hexagonal SiC phase, while c-SiC represents cubic SiC phase.
Figure 2. Characteristics of SiC powder used for fabrication of TiAl-SiC composite: (a) morphology, (b) particle size distribution of powder, (c) EDS analysis result of SiC powder, and (d) XRD pattern, where h-SiC represents hexagonal SiC phase, while c-SiC represents cubic SiC phase.
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Figure 3. XRD patterns of as-sintered TiAl-SiC composite and TiAl alloy.
Figure 3. XRD patterns of as-sintered TiAl-SiC composite and TiAl alloy.
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Figure 4. Backscattered electron (BSE) images of microstructures of as-sintered TiAl-SiC composite (a–c) and TiAl alloy (d,e), in which the images (b,c) are magnified views of matrix and reinforcement phase of TiAl-SiC composite, respectively; labels A–G are EDS testing points of different phases, and the corresponding testing results are shown in Table 1.
Figure 4. Backscattered electron (BSE) images of microstructures of as-sintered TiAl-SiC composite (a–c) and TiAl alloy (d,e), in which the images (b,c) are magnified views of matrix and reinforcement phase of TiAl-SiC composite, respectively; labels A–G are EDS testing points of different phases, and the corresponding testing results are shown in Table 1.
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Figure 5. EPMA mapping analysis of the reinforcement phase of as-sintered TiAl-SiC composite: (a) BSE microstructure, (b) Ti element distribution, (c) Al element distribution, (d) Nb element distribution, (e) W element distribution, (f) B element distribution, (g) Si element distribution, and (h) C element distribution.
Figure 5. EPMA mapping analysis of the reinforcement phase of as-sintered TiAl-SiC composite: (a) BSE microstructure, (b) Ti element distribution, (c) Al element distribution, (d) Nb element distribution, (e) W element distribution, (f) B element distribution, (g) Si element distribution, and (h) C element distribution.
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Figure 6. Backscattered electron (BSE) images of microstructures of as-forged TiAl-SiC composite (a–c) and TiAl alloy (d,e), in which the images (b,c) are magnified views of matrix and reinforcement phase of TiAl-SiC composite, respectively; labels H–N are EDS testing points of different phases; the corresponding testing results are shown in Table 2.
Figure 6. Backscattered electron (BSE) images of microstructures of as-forged TiAl-SiC composite (a–c) and TiAl alloy (d,e), in which the images (b,c) are magnified views of matrix and reinforcement phase of TiAl-SiC composite, respectively; labels H–N are EDS testing points of different phases; the corresponding testing results are shown in Table 2.
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Figure 7. Density of TiAl-SiC composite and TiAl alloy with different states.
Figure 7. Density of TiAl-SiC composite and TiAl alloy with different states.
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Figure 8. XRD patterns of as-forged TiAl-SiC composite and TiAl alloy.
Figure 8. XRD patterns of as-forged TiAl-SiC composite and TiAl alloy.
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Figure 9. EPMA mapping analysis of the reinforcement phase of as-forged TiAl-SiC composite: (a) BSE microstructure, (b) Ti element distribution, (c) Al element distribution, (d) Nb element distribution, (e) W element distribution, (f) B element distribution, (g) Si element distribution, and (h) C element distribution.
Figure 9. EPMA mapping analysis of the reinforcement phase of as-forged TiAl-SiC composite: (a) BSE microstructure, (b) Ti element distribution, (c) Al element distribution, (d) Nb element distribution, (e) W element distribution, (f) B element distribution, (g) Si element distribution, and (h) C element distribution.
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Figure 10. High-temperature compressive properties of as-forged TiAl-SiC composite and TiAl alloy: (a) high-temperature compressive true stress–true strain curves of as-forged TiAl-SiC composite at different temperatures, (b) high-temperature compressive true stress–true strain curves of as-forged TiAl alloy at different temperatures, and (c) dependence of high-temperature compressive yield strength of as-forged TiAl-SiC composite and TiAl alloy on temperature.
Figure 10. High-temperature compressive properties of as-forged TiAl-SiC composite and TiAl alloy: (a) high-temperature compressive true stress–true strain curves of as-forged TiAl-SiC composite at different temperatures, (b) high-temperature compressive true stress–true strain curves of as-forged TiAl alloy at different temperatures, and (c) dependence of high-temperature compressive yield strength of as-forged TiAl-SiC composite and TiAl alloy on temperature.
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Table 1. Chemical composition analysis of constituent phases of as-sintered TiAl-SiC composite and TiAl alloy.
Table 1. Chemical composition analysis of constituent phases of as-sintered TiAl-SiC composite and TiAl alloy.
LabelsChemical Composition (at.%)Identified Phases
TiAlNbWSiC
A37.71 ± 1.2441.91 ± 1.276.50 ± 0.180.44 ± 0.050.01 ± 0.0213.43 ± 2.63γ + Ti2AlC
B23.78 ± 2.1924.81 ± 1.864.06 ± 0.210.29 ± 0.0611.44 ± 2.1435.64 ± 3.27Ti2AlC
C40.96 ± 2.2621.31 ± 0.551.86 ± 0.460.13 ± 0.040.17 ± 0.1435.57 ± 2.78Ti2AlC
D40.04 ± 0.5013.15 ± 2.877.90 ± 0.240.80 ± 0.0516.70 ± 2.8021.41 ± 0.67Ti3SiC2
E37.28 ± 1.8214.87 ± 1.135.67 ± 0.530.47 ± 0.0610.11 ± 1.2531.60 ± 3.51Ti2AlC
F46.39 ± 0.2446.28 ± 0.126.83 ± 0.090.50 ± 0.10----γ
G53.43 ± 1.5839.66 ± 1.496.28 ± 0.090.63 ± 0.04----α2
Note: The locations corresponding to labels A–G are shown in Figure 4; the format of EDS testing results is mean ± standard deviation.
Table 2. Chemical composition analysis of constituent phases of as-forged TiAl-SiC composite and TiAl alloy.
Table 2. Chemical composition analysis of constituent phases of as-forged TiAl-SiC composite and TiAl alloy.
LabelsChemical Composition (at.%)Identified Phases
TiAlNbWSiC
H38.55 ± 0.3543.26 ± 0.726.05 ± 0.050.46 ± 0.050.05 ± 0.0611.63 ± 1.00γ + Ti2AlC
I43.10 ± 0.8322.19 ± 0.023.42 ± 0.050.24 ± 0.070.35 ± 0.2230.70 ± 1.01Ti2AlC
J42.39 ± 1.8912.27 ± 0.318.44 ± 0.310.81 ± 0.1117.56 ± 1.2418.52 ± 2.96Ti3SiC2
K40.10 ± 0.6915.40 ± 0.676.56 ± 0.380.61 ± 0.0911.11 ± 0.7626.22 ± 1.98Ti2AlC
L51.80 ± 0.3041.13 ± 0.336.45 ± 0.060.62 ± 0.02----α2/γ
M46.17 ± 0.1646.74 ± 0.076.59 ± 0.100.49 ± 0.04----γ
N52.40 ± 1.2040.51 ± 1.216.44 ± 0.090.65 ± 0.03----α2
Note: The locations corresponding to labels H–N are shown as the Figure 6; the format of EDS testing results is mean ± standard deviation.
Table 3. High-temperature compressive properties of as-forged TiAl-SiC composite and TiAl alloy.
Table 3. High-temperature compressive properties of as-forged TiAl-SiC composite and TiAl alloy.
Temperature (°C)As-Forged TiAl AlloyAs-Forged TiAl-SiC Composite
Yield Strength σ0.2 (MPa)Maximum Flow Stress σm (MPa)Yield Strength σ0.2 (MPa)Maximum Flow Stress σm (MPa)
700512.77--485.34--
800473.14--444.49606.00
900377.63493.85396.83459.54
1000206.95239.73227.58277.75
Note: Considering that the thermal simulation compression examinations of as-forged TiAl-SiC composite and TiAl alloy stopped at the true strain of 30%, the specimens did not fail at this moment. “--” represents that the maximum flow stress was not reached during the process of thermal simulation compression.
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Liu, S.; Guo, H. Microstructure and High-Temperature Compressive Properties of a Core-Shell Structure Dual-MAX-Phases-Reinforced TiAl Matrix Composite. Crystals 2025, 15, 363. https://doi.org/10.3390/cryst15040363

AMA Style

Liu S, Guo H. Microstructure and High-Temperature Compressive Properties of a Core-Shell Structure Dual-MAX-Phases-Reinforced TiAl Matrix Composite. Crystals. 2025; 15(4):363. https://doi.org/10.3390/cryst15040363

Chicago/Turabian Style

Liu, Shiqiu, and Huijun Guo. 2025. "Microstructure and High-Temperature Compressive Properties of a Core-Shell Structure Dual-MAX-Phases-Reinforced TiAl Matrix Composite" Crystals 15, no. 4: 363. https://doi.org/10.3390/cryst15040363

APA Style

Liu, S., & Guo, H. (2025). Microstructure and High-Temperature Compressive Properties of a Core-Shell Structure Dual-MAX-Phases-Reinforced TiAl Matrix Composite. Crystals, 15(4), 363. https://doi.org/10.3390/cryst15040363

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