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Review

Microstructure and Mechanical Properties of TiAl Matrix Composites Reinforced by Carbides

by
Ying Yang
1,
Yongfeng Liang
1,*,
Chan Li
2,* and
Junpin Lin
1
1
State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China
2
Nuclear and Radiation Safety Center, Ministry of Environmental Protection, Beijing 100082, China
*
Authors to whom correspondence should be addressed.
Metals 2022, 12(5), 790; https://doi.org/10.3390/met12050790
Submission received: 2 April 2022 / Revised: 26 April 2022 / Accepted: 28 April 2022 / Published: 4 May 2022
(This article belongs to the Special Issue Microstructure and Properties of High Temperature Intermetallic)
Browse Figures
Versions Notes

Abstract

:
TiAl alloys have the potential to become a new generation of high-temperature materials due to their lightweight and high-strength properties, while the brittleness at room temperature and microstructure stability at elevated temperature are the key problems. The preparation of composite materials is an effective way to solve these problems, because the mechanical properties of TiAl matrix composites can be improved by the close combination of the reinforced phase and matrix. The preparation methods, microstructure, and mechanical properties of TiAl matrix composites reinforced by carbides are reviewed from the literature in this paper. A comprehensive summary of the effect of C on TiAl alloys can reveal the relationship between the microstructure and mechanical properties and provide guidance for subsequent experimental works. Two forms of C in TiAl matrix composites are reviewed: solid solutions in matrix and carbide precipitations. For TiAl alloys, the minimum carbon content for the carbide precipitation is about 0.5 at.% for low-Nb-containing TiAl alloys and about 0.8 at.% for high-Nb-TiAl alloys. An appropriate amount of C can improve the tensile properties and flexural strength of TiAl alloys. The hardness of the composites is higher than that of pure TiAl due to solution strengthening when the carbon content is low. The minimum creep rate of TiAl alloys can be reduced by one order of magnitude by adding C at the amount near the solubility limit.

1. Introduction

In recent years, TiAl alloys have gradually developed into a new generation of alloys in the field of high-temperature materials. Its lightweight and high-strength properties, as well as attractive antioxidant properties, make it possible to replace nickel-based superalloys in the temperature range of 650~800 °C. TiAl alloys have been developed to the third generation after decades of involution, including high-Nb-TiAl alloys, block-transformed alloys, and β-solidified alloys [1]. These alloys can be used at temperatures of around 760 °C, such as turbine blades of aeroengines, automobile engines, and exhaust valves of racing cars [2].
The stable phases in Ti-Al intermetallic compounds mainly include Ti3Al, TiAl, and TiAl3, the crystal structures of which are shown in Figure 1. Ti3Al is called an α2 phase with good plasticity at room temperature, which can obviously improve the room-temperature brittleness of Ti-Al alloys; however, its wide production and application are limited by its easy oxidation and greatly reduced strength [3,4]. γ-TiAl has good oxidation resistance and mechanical properties, but the plasticity and toughness at room temperature is relatively poor [5]. TiAl3 has great potential as a high-temperature lightweight material for its lowest density and the best oxidation resistance due to its high Al content. However, its development has been relatively slow because of the room-temperature brittleness [6,7]. At present, the main challenges for Ti-Al alloys are the low-room-temperature ductility (0~2%) and insufficient high-temperature microstructure stability. In view of these shortcomings, the prime ways to improve the microstructure and mechanical properties of Ti-Al alloys can be divided into two categories:
(1) Adding alloying elements to control microstructure, and consequently improving mechanical properties at room temperature and elevated temperature, e.g., Mo, Nb, Cr. The presence of Mo can increase the composition of the β/B2 phase in TiAl alloys, which can improve the high-temperature plasticity of the alloys [8]. The β/B2 phase can prevent the growth of α grains during the cooling process, for it is formed along the grain boundary, and thus refines the grain. However, the resulting reduction in the α phase will reduce the strength of the alloys [9]. Because of the low diffusion rates of Nb and Mo elements, they can delay the climbing process of diffusion-assisted dislocation, improve the creep resistance, and significantly reduce the steady creep rate. At the same time, Mo can also improve the oxidation resistance of TiAl alloys, but excessive doping will form oxides that are difficult to degrade, such as MoO3, which is not conducive to improving the oxidation resistance [10,11]. The addition of Cr (>8 at.%) can promote the formation of an Al2O3 oxide layer and reduce the thermal stress, but such a large amount of Cr will affect the plasticity and fracture toughness [12].
(2) Preparation of the composites in which the TiAl matrix combined with reinforced phases can improve the high-temperature microstructure stability. According to the morphology of the reinforced phases, the composites can be divided into two main types: continuous fiber-reinforced type and noncontinuous particle-reinforced type.
For TiAl matrix composites, noncontinuous particle-reinforced TiAl matrix composites have become an important research direction with the advantages of a simple preparation process, low cost, and controlled performance. The reinforced phases include Y2O3 [13], TiB2 [14], TiC [15], Ti2AlC [16], and B4C [17]. The addition of majority-reinforced phases will inevitably reduce plasticity while improving strength, which has an influence on the practical application of TiAl alloys at elevated temperature. Therefore, it is of great practical significance to understand the existence forms of reinforced phases in TiAl matrix composites and their influence on mechanical properties. In this work, carbide-reinforced TiAl matrix composites are selected as the research object, as the previous studies have shown that an appropriate addition of carbide can increase the strength without seriously deteriorating its plasticity [18]. In addition, the preparation methods are summarized and analyzed from the literature in this field in recent years, in order to provide theoretical guidance for subsequent experiments.

2. The Preparation Methods

2.1. Mechanical Alloying

Mechanical alloying (MA) was proposed by Benjamin [19] in 1970 as a high-energy ball milling technique for the preparation of alloy powders. The MA process is simple and does not go through the melting process. It is suitable for the alloying of refractory metals and the generation of the nonequilibrium phase [20]. Gu et al. [21] prepared nano-composite powders of a Ti(Al) solid solution matrix reinforced by in situ TiC by high-energy MA technology. They found that with the increase in grinding time, the morphology of powder with an irregular initial shape underwent continuous change, and such a microstructure change was determined by the competition between cold welding and particle fracture. Karimi et al. [22] prepared TiAl(Nb)/Ti2AlC composites by MA and hot pressing sintering, and investigated the effect of the Nb element on the microstructure and oxidation resistance of the composites.
At present, the quantitative description and phase transformation process analysis of the MA technology are not clear. Future research could focus on powder deformation mechanics to illustrate the strain–stress process during the ball milling and the changes in microstructure and grain size.

2.2. Spark Plasma Sintering

Spark plasma sintering (SPS) can produce high temperature through the instantaneous discharge between powder particles for material sintering [23], which can greatly shorten the sintering time and reduce the sintering temperature. The prepared materials have fine grains, low pollution, and excellent performance [24]. Yang et al. [25] prepared TiAl matrix composites using SPS technology at different sintering temperatures and studied the phase composition. Mei et al. prepared Ti2AlC/TiAl composites by sintering mixed powders of Ti, A1, and TiC [26]. When the mixed powder contained 7 vol.% TiC, the fracture toughness KIC of the composite sintered at 1200 °C was up to 43 MPa·m1/2, which was 15% higher than that of TiAl alloys sintered by SPS.
At present, SPS technology is widely used and its technological maturity is constantly improving [27]. In the future, it can be studied to improve the versatility of equipment and the capacity of pulse current, so as to prepare large-size products. The development of mold materials with a higher strength than the graphite mold currently used can improve the bearing capacity of the mold and reduce the mold cost.

2.3. Hot Pressing and Hot Isostatic Pressing

Reactive sintering by hot pressing (HP) and hot isostatic pressing (HIP) can effectively eliminate the residual pores of products and obtain materials that are close to complete density [28]. Li et al. [29] prepared TiAl/Ti2AlNb composites by the HP method and revealed the relationship between the deformation behavior and fracture process of composites. The HP method may have thermal stress due to the uneven internal and external temperature distribution in the preparation of large-size samples. The HIP method usually has a higher tensile plasticity and better compactness than the HP method.

2.4. Combustion Synthesis

Combustion synthesis is also known as self-propagating high-temperature synthesis (SHS). The combustion temperature of the SHS method can reach 5000 K, and the reaction time is short the synthesis speed is very fast [30]. The composite materials synthesized by SHS have the advantages of a clean interface, and fast and simple process. Ramaseshana et al. [31] successfully prepared Ti2AlC/γ-TiAl composites by combustion synthesis and casting, the strength of the composite was 800 MPa at room temperature and 400 MPa at 900 °C, the fracture toughness was 17.8 MPa·m1/2, and it had excellent oxidation resistance. Andreev et al. [32] prepared TiAl matrix composites using the spin-casting SHS method and optimized the process parameters. XD technology is a material preparation technology developed on the basis of combustion reaction [33]. It is used to evenly mix the elemental powder of the matrix and reinforced phase, form a blank by cold pressing or hot pressing, and heat it quickly to the temperature above the melting point of the matrix but below the synthesis temperature of reinforcement, and the composite material with a uniform distribution of the reinforced phase is generated by in situ reaction.
It is worth noting that in the process of preparing TiAl matrix composites by powder metallurgy using elemental powders, because Ti is easy to react with O and N, the sintering of TiAl matrix composites generally needs to be carried out in a protective atmosphere, which greatly increases the production cost, and TiH2 may be an effective solution to this problem [34].
Ivasishin et al. [35] found in 1999 that using TiH2 instead of Ti powder to prepare TiAl alloys can reduce the sintering temperature and obtain high-density samples without applying high pressure. Robertson and Schaffer [36] concluded that enhanced densification using TiH2 powders may result from the volume reduction during TiH2 decomposition and finer original particle size, reduced oxygen content in hydride powders, and the enhanced chemical activity of the Ti surface with the release of atomic hydrogen. Wang et al. [37,38] proposed a pretreatment method that can be used in TiH2 and Al powders to obtain a TiAl matrix composite with high-density and satisfactory mechanical properties. Yan et al. [39] prepared a nonspherical pre-alloyed TiAl powder by using TiH2 and Al powder, and prepared Ti-48Al alloys with high density and low oxygen content.
In addition, the casting method is also commonly used in the preparation of composite materials. Lapin et al. [40] studied the effects of centrifugal casting and tilt casting on the microstructure and properties of TiAl matrix composites. It was found that the application of the casting process affects the size of carbide particles and the structure of the intermetallic compound matrix. The alloy made by centrifugal casting has a finer size of carbide particles. The Vickers hardness and matrix Vickers microhardness of the alloy made by centrifugal casting are higher than those made by tilt casting ones.
Liu et al. [41] fabricated Ti-Al-C composites by infiltration in situ reaction. The compound mechanism of the Ti-Al-C system can be described as formation-decomposition-precipitation-oxidation. Zhang et al. [42] prepared cast rods of TiAl alloys with a lamellar microstructure by cold crucible levitation melting, and the effects of carbon content on creep properties and tensile properties at room temperature were investigated. Li et al. [43] prepared a carbide and boride-reinforced TiAl matrix composite by induction skull melting (ISM) with a water-cooled copper crucible, and studied the effect of the addition of B/C on the mechanical properties of the alloy. The results showed that the alloy with B4C addition had finer grains.

3. The effect of C on Microstructure of Composites

3.1. Solubility of C in TiAl Alloys

C has a certain solid solubility in Ti-Al alloys, which is different in α2 and γ phases. From the perspective of crystal structures, it is explained that the appropriate environment occupied by C atoms in the sublattice is related to the surrounding atoms [44]. The preferred octahedral coordination of C atoms in α2 and γ phases is shown in Figure 2, where one C atom is surrounded by six Ti(Al) atoms. Different Ti-Al-C ternary systems are formed in α2 and γ phases: hexagonal-type Ti2AlC for the α2 phase and perovskite-type Ti3AlC for the γ phase.
In the α2 phase with an HCP crystal structure, C atoms are solidly dissolved in the octahedral interstice of the crystal structure and form a coordination octahedron with six Ti atoms [45]. In the L10-type FCC crystal structure of the γ phase, the corner positions of cells, as well as the bottom and top surface center positions, are occupied by Al atoms, and the remaining surface center positions are occupied by Ti atoms. The octahedron has Ti4Al2 and Ti2Al4 atomic occupancies in the L10-lattice, which is not suitable for accommodating an interstice C atom. Therefore, the solubility of C in the γ phase is lower than that in the α2 phase [46], the value of which is about 0.3 at.% in the γ phase, and about 1.2 at.% in the α2 phase [47].
Scheu et al. [48] found that the solid solubility of C in the γ phase could be improved by adding an Nb element. The addition of Nb leads to the formation of Ti anticrystal defects in the crystal. Nb atoms preferentially occupy Ti sites in the L10-lattice, forcing Al atoms to be replaced by Ti atoms in the Al sublattice, thus forming an octahedron surrounded by six Ti(Nb) atoms, which is more suitable for accommodating interstice C atoms. The solubility limit of C in the TiAl alloy is also related to the volume fraction of γ and α2 phases. Perdrix et al. [49] found that for Ti-48Al-based alloys, the solubility limit of C was 3000 wt.pmm, equivalent to 0.86 at.%. Dong et al. [50] found that for the Ti-46Al-2Cr-1.5Nb-0.2Si alloy, carbides began to precipitate when the carbon content reached 0.5 at.%, while no carbides were formed in the Ti-45Al-5Nb-0.5C alloy with the same carbon content of 0.5 at.% [18]. This proves that Nb can improve the solubility of C in TiAl alloys. In the Ti-46Al-8Nb-0.7C alloy with a higher Nb content, no carbides were precipitated, and carbides appeared when the carbon content increased to 1.4 at.% [51].

3.2. Effect of Carbide Precipitation on Microstructure

Cabibbo found that the carbides precipitate when the carbon content is higher than 0.9 at.% in the Ti-46Al-4Nb system [52], while Gabrisch et al. found that the carbides precipitate when the carbon content is higher than 0.75 at.% in the Ti-45Al-5Nb alloy [46]. The minimum carbon content for carbide precipitation is different for TiAl alloys with different addition elements, which is generally in the range of 0.75~1.0 at.% for Nb-containing TiAl alloys. Witusiewicz et al. [53] gave a thermodynamic description of the ternary Al–C–Ti system that used the CALPHAD approach (Thermo-Calc/PARROT). The results showed that H (Ti2AlC), P (Ti3AlC), and N (Ti3AlC2) in the system are thermodynamically stable over a wide range of temperatures. Therefore, the main reinforced phases are H (Ti2AlC), P (Ti3AlC), and N (Ti3AlC2). In addition, TiC usually occurs at the initial stage of the composite preparation reaction. These carbide crystal structures are shown in Figure 3.

3.2.1. Ti2AlC

As shown in Figure 3a, Ti2AlC is a typical compound of the MAX phase [54]. Each cell contains 2 Ti2AlC molecules. According to the literature [55], the crystal structure of Ti2AlC contains mixed bonding types of metallic, covalent, and ionic, which makes it exhibit a MAX phase. Due to the existence of ionic bonds, Ti2AlC has excellent electrical conductivity. The bond between the Ti atom and C atom is a strong covalent bond, which gives the material a high strength and a high elastic modulus. The layered structure results in a very weak bonding between the Ti and Al atomic plane, making the material self-lubricating and anisotropic.
The density and thermal expansion coefficient of Ti2AlC (4.11 g/cm3 and 8.8 × 10−6/K, respectively) are close to those of the TiAl-based alloy (3.8 g/cm3 and 12 × 10−6/K, respectively) [56,57], which can avoid the segregation phenomenon of reinforced particles to a large extent. In addition, the internal stress between the reinforced particles and matrix during the manufacturing process of the composite material is greatly reduced. Research by Yue et al. [58] also verified this statement. They found that the interface between the TiAl and Ti2AlC prepared by SPS technology was neatly bonded and there was no common residual stress layer on the interface, which was beneficial to the preparation of high-strength composites.
Chen et al. [59] carried out phase transformation analysis and a synthesis mechanism study on samples of hot-pressing sintered powders of Ti, Al, TiC and Ti, Al, C elements at 600~1300 °C. The results showed that below 900 °C, Ti reacted with Al to form Ti-Al intermetallic compounds, and above 900 °C, TiAl reacted with TiC to form dense TiAl/Ti2AlC composites. The Ti2AlC particles were evenly distributed in the TiAl matrix after sintering at 1200 °C. Lapin et al. [60] prepared (Ti,Nb)2AlC particle-reinforced TiAl matrix composites using centrifugal casting. As shown in Figure 4, the microstructure of as-cast composites was composed of (Ti,Nb)2AlC particles (1 and 2) distributed in the matrix composed of layered γ + α2 (3), a single γ (4) phase region, and β/B2 (5) particles.
Heat treatment has a significant effect on the microstructure of TiAl matrix composites. Lapin et al. [60] found that after heat treatment, the matrix of the composites containing 1.4 at.% C maintained the microstructure of γ + α2 + β/B2 type, while the matrix of the composites with a carbon content of 3.6 at.% transformed into nearly a γ phase, which is shown in Figure 5. As shown in Figure 6, Ramaseshan [31] found that after heat treatment of Ti2AlC/TiAl matrix composites prepared by SHS technology, the matrix structure changed from layered α2 + γ to near γ, and the precipitation of fine secondary carbide particles Ti2AlC was found along the lamellar direction. This indicates that the α2 phase containing carbon is decomposed into γ-TiAl and Ti2AlC particles during homogenization. Yue et al. [61] also carried out multi-step heat treatment studies on Ti2AlC/TiAl matrix composites, and found that Ti2AlC dispersed in the matrix can pin α and γ grain boundaries, thus hindering grain boundary migration and preventing grain growth. At the same time, γ grains can nucleate out from the Ti2AlC/matrix interface, which increases the nucleation location of γ grains and improves the nucleation rate.

3.2.2. Ti3AlC

As shown in Figure 3d, Ti3AlC has a cubic perovskite structure [62]. Studies on Ti3AlC have mainly focused on the formation mechanism and carbide evolution. Lapin et al. studied the evolution of carbides in TiAl matrix composites with the alloy system of Ti-44.6Al-7.9Nb-3.6C-0.7Mo-0.1B [63]. The carbides in the as-cast samples were mainly coarse H-Ti2AlC phases. A fine secondary P-Ti3AlC phase and H-Ti2AlC phase were precipitated after heat treatment. Zhou et al. [51] found that when the carbon content increased to 1.4 at.%, elongated H-Ti2AlC precipitated with a random distribution in Ti-46A1-8Nb alloys after solution treatment at 1380 °C. After aging at 900 °C for 6 h, acicular P-Ti3AlC precipitated, and the amount of P-Ti3AlC increased with the carbon content. The length of Ti3AlC was about 130 nm, and it precipitated from the γ phase with an orientation relationship with the γ phase: (100)Ti3AlC//(100)γ, <001>Ti3AlC//<001>γ. In addition to the distribution in the γ phase, Ti3AlC also distributed at the grain boundary and lamellar interface.
Mei et al. [64] found that a small amount of Ti3AlC was generated during hot pressing at a temperature above 1200 °C. At the initial stage of reaction, there was a large amount of Ti, Al, and TiC in the green body, and the contact between them was sufficient; therefore, the ratio of reaction products and raw materials was consistent, that is, Ti2AlC was generated. At the later stage of the reaction, the reaction was fully carried out, a large amount of Ti2AlC was formed in the green body, and only a small amount of reactants Ti-Al and TiC were left. In this way, an Al-deficient environment was easily created, and Ti3AlC with a similar layered structure to Ti2AlC but lacking more Al was generated.

4. Effect of C on Mechanical Properties and Its Mechanism

4.1. Tensile Properties

As for the tensile properties of composites, many researchers have conducted studies, as shown in Table 1, which shows a summary of the data of elongation and ultimate tensile strength (UTS) of TiAl matrix composites with different carbon contents. It can be seen that the addition of C can increase the UTS, but generally has little effect or a slight reduction in plasticity. Zhang et al. [42] found that an alloy with 0.05 at.% C, 0.1 at.%C, and 0.2 at.%C could improve the UTS by 7.7%, 9.3%, and 18.0% but reduce the elongation by 8%, 24%, and 52%, respectively. The addition of 0.2 at.% C had the highest degree of improvement on UTS, while the addition of 0.1 at.% C resulted in a good match between strength and plasticity of the alloy. Chlupov´A et al. [65] found that the UTS increased with the carbon content, while the plastic elongation showed a slight decline.
Wang et al. [44] found that the improvement of the tensile strength of TiAl alloys by the addition of C was mainly due to the refinement of interlayer spacing. As shown in Figure 7, the addition of C can reduce the lamellar spacing, and there were more semi-coherent interfaces in the thin lamellar TiAl alloy, which can hinder the dislocation movement. During the tensile process, the dislocations accumulated at the semi-coherent interface, which requires higher stress to drive the dislocations, so the carbide-reinforced TiAl matrix composites had a higher yield tensile strength than pure TiAl. In addition, the volume fraction of the B2 phase is another factor affecting the tensile strength of TiAl-based alloys at room temperature. The researchers noted that the presence of the B2 phase is a common source of cracking, damaging the tensile properties of TiAl matrix composites [67,68,69]. As mentioned above, the addition of C promoted the β→α phase transition, resulting in the decrease in the residual high-temperature β phase. When the temperature dropped to room temperature, the amount of B2 phase in the alloy decreased, and the tensile properties improved.

4.2. Flexural Properties and Fracture Toughness

Flexural strength is a comprehensive reflection of strength and toughness, which has reference value for evaluating the mechanical properties of alloys. Table 2 shows the data of flexural strength and fracture toughness in TiAl matrix composites with the change in carbon content. It can be seen that the flexural strength and fracture toughness have a similar trend with carbon content. Generally, there is a medium carbon content corresponding to the best properties. A higher or lower carbon content will have adverse effects on the flexural properties.
Gong et al. [70] prepared TiAl and TiAl/TiC composites by SPS. The experimental results showed that the flexural strength of TiAl/TiC composites greatly improved compared with that of pure TiAl. The maximum flexural strength was obtained when the amount of TiC was 10 wt.%, equivalent to the carbon content of 7.1 at.%. The flexural strength increased by 22%, from 643 MPa to 784.5 MPa. The fracture toughness of 3.6 at.% C containing TiAl/TiC composites was 16.8 MPa·m1/2, which was 39% higher than that of pure TiAl, and then the fracture toughness of composites reduced with a higher carbon content. The fracture toughness of the composites was lower than that of pure TiAl when the carbon content was over 10.6 at.%. Li et al. [71] found that the flexural strength of Ti2AlC/TiAl matrix composites prepared by vacuum hot pressing sintering could reach 743.84 MPa, during which the formation of Ti2AlC inhibited the abnormal growth of TiAl crystals. It is worth noting that Ruan et al. [72] found that the flexural strength and fracture toughness of the composite decreased with the carbon content in the Ti-(43-48)Al-(0-6.5)C alloy. Agglomeration of the reinforced phases was obvious, and the material structure was loose, as the relative density decreased from 92.7% to 86.8% with the increase in carbon content, which was the main reason for the deterioration in its mechanical properties.
Compared with pure TiAl, the crack propagation is retarded and the fracture energy is increased due to the presence of embedded carbides in composites. The improvement of fracture toughness can be achieved by increasing the strength of the reinforcement [73]. Lapin et al. [74] suggested that the brittle fracture of composites was mainly characterized by crack deviation, carbide fracture, stratification at the interface between the carbide particles and matrix, and carbide particles pulling out of the TiAl matrix, which is shown in Figure 8. During the experiment, they found that TiC particles existed in some coarse Ti2AlC particles, which improved the overall fracture toughness of composites by preventing the crack propagation in Ti2AlC particles.
Yue et al. [61] also studied the influence of the multi-step heat treatment on the fracture toughness of Ti2AlC/TiAl composites. The results showed that the flexural strength and fracture toughness of the composites reached 957.9 MPa and 20.73 MPa·m1/2 after heat treatment at 1390 °C. The fracture mode of the composites changed from the mixed mode of intergranular fracture and transgranular fracture to transgranular cleavage fracture after the multi-step heat treatment. Because the transgranular cleavage fracture consumes a large amount of energy inside the composites and inhibits the crack propagation, the flexural strength and fracture toughness of the composites after heat treatment can be greatly improved.

4.3. Hardness

Table 3 shows the hardness variation of TiAl matrix composites with different carbon contents. It can be seen from Table 3 that the hardness of composites was enhanced with the low addition of C. Zhang et al. [76] found that the hardness of the Ti-47Al-4Nb-2Cr-(0–0.8)C alloy increased by 38.9%, from 2.85 GPa to 3.96 GPa, with the increased carbon content. As shown in Figure 9, Gabrisch et al. [46] believed that this was because of the solid solution strengthening in the matrix by adding an appropriate amount of C, resulting in lattice distortion and increasing hardness. When the carbon content was excessive, the hardness decreased with the carbon content. Ruan et al. [72] found that the hardness of the Ti-(43-48)Al-(0–6.5)C alloy decreased by 18.5% as the carbon content increased from 0 at.% to 6.5 at.%. Yang et al. [75] considered that the main reason for the decrease in hardness was the increase in Ti2AlC formed in the TiAl matrix, which broke the inherent bonding force of the TiAl grain boundary.

4.4. Creep Properties

The improvement of the creep property is very important for improving the service life of TiAl alloys. Table 4 shows the creep test data of some TiAl matrix composites containing C. Zhang et al. [42] found that adding 0.05 at.%~0.2 at.% C reduced the transient creep strain, plastic strain, and creep rate at 200 h of TiAl alloys, and the improvement increased with the carbon content. Among them, adding 0.1 at.% C reduced the plastic creep strain by half. The creep rate under the same strain decreased by more than one order of magnitude. The results from Lapin et al. [78] showed that the creep resistance of Ti-46.4Al-5.1Nb-0.2B-1C was better than that of low-carbon alloy (Ti-47Al-5.2Nb-0.2B-0.2C) at the temperature of 800 °C. This was also confirmed by the results of Zhou et al. [79] and Zhang et al. [80]. For the Ti-(42.0-42.6)Al-8.7Nb-0.3Ta-(2.0-3.6)C alloy [81], the minimum creep rate of the alloy with 2.0 at.% C was lower, indicating that excessive carbon content will adversely affect the creep property. Li et al. [66] found that the precipitation of a large amount of Ti3AlC improved the creep resistance of the high-Nb-TiAl alloys with a trace element of C in the long-term creep process.
Kamyshnykova et al. [82] considered that the deformation structure of composites under low creep strain, which corresponds to the minimum creep rate, was mainly dislocation in the TiAl matrix, and that under high strain was mainly caused by deformation twinning. Worth et al. [83] believed that for the same alloys, the full-lamellar structure has a higher creep resistance than the duplex and equiaxial γ structure, and it is also because the fine lamellar spacing of the lamellar structure reduces the effective slip length of dislocation motion. The presence of carbides is an effective obstacle to dislocation motion, and dislocation bypasses these precipitates by climbing. At the same time, the carbides in the initial microstructure improve the strain hardening capacity of the alloys, and the carbides precipitated at the interrupted γ lamellar step play a dynamic hardening effect during the creep process, which is also the reason why the addition of C improves the homogeneity of creep deformation at high stress in the TiAl alloy with a lamellar microstructure [42].

5. Conclusions

The preparation of TiAl matrix composites can effectively solve the problem of poor microstructure stability of TiAl alloys at elevated temperature, which makes it possible to use TiAl matrix composites at high temperature. In this paper, the microstructure and mechanical properties of TiAl matrix composites are summarized and analyzed from the literature. There are two forms of C in TiAl matrix composites: solid solution in matrix and carbide precipitations. The C solid solubilized in the matrix mainly plays the role of solution strengthening. Carbides can be distributed in the interlamellar or grain boundary to improve the mechanical properties of the alloys due to its good thermal stability. Ti2AlC, one of the MAX phases, is regarded as an ideal reinforcing phase in the composite materials. For high-Nb-TiAl alloys, the minimum carbon content for carbide precipitation is generally between 0.75 and 1.0 at.%. The effects of C on the mechanical properties of TiAl matrix composites are mainly concluded as follows:
  • Proper amount of C can improve the tensile properties of TiAl matrix composites. An optimized tensile properties can be obtained with carbon content of about 0.2 at.%.
  • The flexural strength of the alloys can be greatly improved with addition of C, due to the formation of Ti2AlC, which can hinder the abnormal growth of grain.
  • For TiAl matrix composites, the hardness of the composites is higher due to solution strengthening when the carbon content is low; when superfluous C is added, the carbide precipitates destroy the inherent bonding force of TiAl grain boundary, leading to the decrease of hardness.
  • The minimum creep rate of TiAl can be reduced by one order of magnitude by adding C at about 0.5 at.%.
In general, when the carbon content is near the solubility limit in TiAl matrix composites, the composites show relatively comprehensive mechanical properties.

Author Contributions

Writing—original draft preparation, methodology, investigation, Y.Y.; design guide, writing assistance, funding acquisition, supervision, Y.L.; writing assistance, methodology, data analysis, C.L.; supervision, design guide, funding acquisition, writing assistance, J.L. All authors have read and agreed to the published version of the manuscript.

Funding

This work was funded by the National Key R&D Program of China [grant number 2021YFB3700501], the National Natural Science Foundation of China [grant number 51831001], the Funds for Creative Research Groups of China [grant number 51921001] and the Fundamental Research Funds for the Central Universities [FRF-MP-20-44].

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Conflicts of Interest

The authors declare no conflict of interest.

References

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Figure 1. The crystal structure of the Ti-Al intermetallic compounds: (a) Ti3Al; (b) TiAl; (c) TiAl3.
Figure 1. The crystal structure of the Ti-Al intermetallic compounds: (a) Ti3Al; (b) TiAl; (c) TiAl3.
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Figure 2. Atomic occupancy of interstitial C in (a) α2 (Ti3Al) and (b) γ (TiAl).
Figure 2. Atomic occupancy of interstitial C in (a) α2 (Ti3Al) and (b) γ (TiAl).
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Figure 3. Crystal structure diagram of (a) Ti2AlC; (b) Ti3AlC2; (c) TiC; (d) Ti3AlC.
Figure 3. Crystal structure diagram of (a) Ti2AlC; (b) Ti3AlC2; (c) TiC; (d) Ti3AlC.
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Figure 4. SEM micrographs showing the microstructure of the as-cast composite with various contents of carbon: (a) and (b) C1.4; (c) and (d) C3.6; 1 and 2—(Ti, Nb)2AlC, 3—(γ + α2), 4—γ, 5—(β/B2). Reproduced from [60], with permission from Elsevier, 2017.
Figure 4. SEM micrographs showing the microstructure of the as-cast composite with various contents of carbon: (a) and (b) C1.4; (c) and (d) C3.6; 1 and 2—(Ti, Nb)2AlC, 3—(γ + α2), 4—γ, 5—(β/B2). Reproduced from [60], with permission from Elsevier, 2017.
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Figure 5. BSEM micrographs showing the typical microstructure of heat-treated in situ composites: (a) C1.4; (b) C2.5; 1 and 2—(Ti, Nb)2AlC, 3—(γ + α2), 4—γ, 5—(β/B2). Reproduced from [60], with permission from Elsevier, 2017.
Figure 5. BSEM micrographs showing the typical microstructure of heat-treated in situ composites: (a) C1.4; (b) C2.5; 1 and 2—(Ti, Nb)2AlC, 3—(γ + α2), 4—γ, 5—(β/B2). Reproduced from [60], with permission from Elsevier, 2017.
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Figure 6. TEM micrographs of the composite (Ti(Ni)50Al45C5), showing smaller precipitates in place of lamellar disappearance: (a) as cast; (b) homogenized at 1273K. Reproduced from [31], with permission from Elsevier, 1999.
Figure 6. TEM micrographs of the composite (Ti(Ni)50Al45C5), showing smaller precipitates in place of lamellar disappearance: (a) as cast; (b) homogenized at 1273K. Reproduced from [31], with permission from Elsevier, 1999.
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Figure 7. TEM images of the lamellar morphologies, (a) TNC, (b) TNC-0.2C, (c) TNC-0.5C. Reproduced from [44], with permission from Elsevier, 2017.
Figure 7. TEM images of the lamellar morphologies, (a) TNC, (b) TNC-0.2C, (c) TNC-0.5C. Reproduced from [44], with permission from Elsevier, 2017.
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Figure 8. SEM images of typical fracture surface of the in-situ composite after three-point bending test, (a) The fracture surface with radial cracks in the vicinity of the tip notch region. (b) Laminated tearing of Ti2AlC particles and fine secondary carbide particles in the γ matrix. Reproduced from [74], with permission from Elsevier, 2019.
Figure 8. SEM images of typical fracture surface of the in-situ composite after three-point bending test, (a) The fracture surface with radial cracks in the vicinity of the tip notch region. (b) Laminated tearing of Ti2AlC particles and fine secondary carbide particles in the γ matrix. Reproduced from [74], with permission from Elsevier, 2019.
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Figure 9. Section of the diffraction pattern illustrating the peak shift due to the lattice expansion with increasing C contents. Reproduced from [46], with permission from Elsevier, 2013.
Figure 9. Section of the diffraction pattern illustrating the peak shift due to the lattice expansion with increasing C contents. Reproduced from [46], with permission from Elsevier, 2013.
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Table
Table 1. Tensile properties of the TiAl matrix composites.
Table 1. Tensile properties of the TiAl matrix composites.
CompositionTest Temperture/°CUltimate Tensile Strengh/MPaElongation/%Reference
Ti-47.5Al-3.7(Cr,V,Zr)255692.5[42]
Ti-47.5Al-3.7(Cr,V,Zr)-0.05C6032.3
Ti-47.5Al-3.7(Cr,V,Zr)-0.1C6121.9
Ti-47.5Al-3.7(Cr,V,Zr)-0.2C6611.2
Ti-47Al-2Nb-2Cr255000.8[44]
Ti-47Al-2Nb-2Cr-0.2C5601.05
Ti-47Al-2Nb-2Cr-0.5C5401.0
Ti-47Al-2Nb-2Cr8004003.7
Ti-47Al-2Nb-2Cr-0.2C5005.2
Ti-47Al-2Nb-2Cr-0.5C4704.65
Ti-46Al-7Nb-2Mo254850.058[65]
Ti-46Al-7Nb-2Mo-0.2C5380.032
Ti-46Al-7Nb-2Mo-0.5C6010.008
Ti-46Al-7Nb-2Mo7504850.38
Ti-46Al-7Nb-2Mo-0.2C4990.28
Ti-46Al-7Nb-2Mo-0.5C5240.17
Ti-46Al-8.5Nb257970.28[66]
Ti-46Al-8.5Nb-0.1C8200.07
Ti-46Al-8.5Nb7606812.78
Ti-46Al-8.5Nb-0.1C7230.3
Table
Table 2. Flexural strength and fracture toughness of the TiAl matrix composites.
Table 2. Flexural strength and fracture toughness of the TiAl matrix composites.
CompositionFlexural Strength/MPaFracture Toughness/MPa·m1/2Reference
Ti-50Al64312.1[70]
Ti-46.9Al-3.6C725.816.8
Ti-43.9Al-7.1C784.512.6
Ti-41Al-10.6C649.311.4
Ti-38.1Al-14.2C618.810.1
Ti-48Al3346.20[71]
Ti-46.3Al-4.7C469.058.39
Ti-41.9Al-9.4C743.849.17
Ti-50Al4607.19[75]
Ti-47Al-2Nb-2Cr-0.2C4867.78
Ti-47Al-2Nb-2Cr-0.5C4466.36
Ti-48Al334.686.83[72]
Ti-46.2Al-2.2C310.687.07
Ti-44.4Al-4.4C252.275.48
Ti-42.7Al-6.5C146.683.78
Table
Table 3. Hardness of the TiAl matrix composites.
Table 3. Hardness of the TiAl matrix composites.
CompositionHardness/GPaReference
Ti-47Al-4Al-2Cr2.85[76]
Ti-46.9Al-4Al-2Cr-0.2C3.59
Ti-46.9Al-4Al-2Cr-0.4C3.51
Ti-46.9Al-4Al-2Cr-0.6C3.63
Ti-46.9Al-4Al-2Cr-0.8C3.96
Ti-48Al3.58[77]
Ti-47.8Al-0.5C3.73
Ti-47.5Al-1C3.71
Ti-47Al-2C3.70
Ti-48Al2.98[72]
Ti-46.2Al-2.2C2.69
Ti-44.4Al-4.4C2.59
Ti-42.7Al-6.5C2.43
Ti-50Al2.93[75]
Ti-46.3Al-4.7C2.92
Ti-41.9Al-9.4C2.83
Table
Table 4. Creep properties of TiAl matrix composites.
Table 4. Creep properties of TiAl matrix composites.
CompositionT/°CStress/MPaMinimum Creep Rate/s−1Reference
Ti-47.5Al-3.7(Cr,V,Zr)7601382.2 × 10−9[42]
Ti-47.5Al-3.7(Cr,V,Zr)-0.05C1.2 × 10−9
Ti-47.5Al-3.7(Cr,V,Zr)-0.1C9.2 × 10−10
Ti-47.5Al-3.7(Cr,V,Zr)-0.2C7.4 × 10−10
Ti-46Al-8.5Nb-0.1C-0.2B7603007.9 × 10−9[66]
Ti-45Al-3Fe-2Mo7501502.2 × 10−8[79]
Ti-45Al-3Fe-2Mo-0.5C7.9 × 10−9
Ti-47.5Al-2.5V-1.0Cr-0.2Zr8003001.72 × 10−7[80]
Ti-47.5Al-2.5V-1.0Cr-0.2Zr-0.1C2.98 × 10−8
Ti-47Al-5.2Nb-0.2B-0.2C8002003.45 × 10−8[78]
Ti-46.4Al-5.1Nb-0.2B-1C8.31 × 10−9
Ti-42.6Al-8.7Nb-0.3Ta-2.0C8002009.9 × 10−9[81]
Ti-41.0Al-8.7Nb-0.3Ta-3.6C3.31 × 10−8
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Yang, Y.; Liang, Y.; Li, C.; Lin, J. Microstructure and Mechanical Properties of TiAl Matrix Composites Reinforced by Carbides. Metals 2022, 12, 790. https://doi.org/10.3390/met12050790

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Yang Y, Liang Y, Li C, Lin J. Microstructure and Mechanical Properties of TiAl Matrix Composites Reinforced by Carbides. Metals. 2022; 12(5):790. https://doi.org/10.3390/met12050790

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Yang, Ying, Yongfeng Liang, Chan Li, and Junpin Lin. 2022. "Microstructure and Mechanical Properties of TiAl Matrix Composites Reinforced by Carbides" Metals 12, no. 5: 790. https://doi.org/10.3390/met12050790

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