Room temperature mechanical behaviour of cast in-situ TiAl matrix composite reinforced with carbide particles

Intermetallics 105 (2019) 113–123

Contents lists available at ScienceDirect

Intermetallics journal homepage: www.elsevier.com/locate/intermet

Room temperature mechanical behaviour of cast in-situ TiAl matrix composite reinforced with carbide particles

T

J. Lapina,∗, M. Štamborskáa, K. Kamyshnykovaa,b, T. Pelachováa, A. Klimováa, O. Bajanaa a b

Institute of Materials & Machine Mechanics, Slovak Academy of Sciences, Dúbravská Cesta 9, 845 13 Bratislava, Slovak Republic Slovak University of Technology in Bratislava, Faculty of Materials Science and Technology in Trnava, Ulica Jána Bottu 25, 917 24 Trnava, Slovak Republic

ARTICLE INFO

ABSTRACT

Keywords: Intermetallics TiAl Composites Mechanical properties Fracture behaviour Finite element analysis

Room temperature mechanical behaviour of cast in-situ TiAl matrix composite reinforced with carbide particles was studied using hardness, tensile, compression, and three-point bending tests. The initial microstructure of the in-situ composite consists of TiAl matrix reinforced with coarse primary (HeTi2AlC and TiC) and fine secondary (HeTi2AlC and PeTi3AlC) particles. The in-situ composite shows plastic deformation of TiAl matrix and limited plastic deformation of primary Ti2AlC particles during the compression test. The compressive work hardening behaviour is affected by the plastic deformation of the matrix, plastic deformation of primary carbide particles and formation of cracks. The local equivalent strains in the compression specimens are numerically calculated using finite element analysis (FEA) and related to the size of primary carbide particles. The in-situ composite specimens show brittle fracture behaviour during tensile and three-point bending tests. The numerically calculated critical stress leading to a crack initiation in the notch tip region of three-point bending specimen is comparable with the ultimate tensile strength. The brittle fracture of the in-situ composite includes crack deviation, zone shielding by microcrack toughening, carbide fragmentation, delamination on the matrix-carbide interfaces and pull-out of the carbide particles from the matrix.

1. Introduction Cast TiAl-based alloys are attractive materials for high-temperature structural applications providing a unique set of physical and mechanical properties for stationary gas turbines, turbochargers, automotive engines and aircraft engines [1–10]. However, their inherent poor ductility at room temperature and insufficient strength at high temperatures limit their wide-scale applications [2,11,12]. Intermetallic matrix composites may improve the deficiency of these lightweight alloys at high temperatures due to a good combination of the properties of intermetallic matrix and reinforcement [13–22]. Recent studies have shown that in-situ TiAl matrix composites reinforced with homogenously distributed carbide particles can be prepared by vacuum induction melting followed by gravity or centrifugal casting of TiAlbased alloys with various content of carbon [23–26]. Additional strengthening of these in-situ composites can be achieved by fine secondary needle-like Ti3AlC (P-phase) and plate-like Ti2AlC (H-phase) precipitates similarly to that reported for several low carbon TiAl-based alloys [27–33]. Rössler et al. [34,35] have shown that either very small precipitates or coarse carbide particles are effective to retard creep deformation of in-situ TiAl matrix composites at low creep rates and

∗

high temperatures. The coarse carbide particles improve fracture toughness but alloying with carbon deteriorates room-temperature tensile ductility of the in-situ TiAl matrix composites [36,37]. In spite of the previous studies on in-situ TiAl matrix composites, only limited information is available about the effect of coarse primary carbide particles on their room temperature mechanical behaviour. Song et al. [14] and Shu et al. [38] have reported that an increase of volume fraction of Ti2AlC particles increases room temperature compressive strength of in-situ TiAl matrix composites prepared by powder metallurgy route. On the other hand, Lapin et al. [24] have reported a decrease of room temperature compressive strength with increasing volume fraction of primary Ti2AlC particles in the in-situ composites prepared by centrifugal casting. Since the centrifugal casting represents widely applied technology which might facilitate potential applications of the in-situ composites, it is of large interest to investigate how the coarse primary carbide particles, which have been suggested by Rössler et al. [34,35] to improve high temperature creep strength, affect room temperature mechanical behaviour of centrifugally cast in-situ composites. The aim of this paper is to study room temperature mechanical behaviour of cast in-situ composite reinforced with carbide particles

Corresponding author. E-mail address: [email protected] (J. Lapin).

https://doi.org/10.1016/j.intermet.2018.11.007 Received 1 August 2018; Received in revised form 8 November 2018; Accepted 9 November 2018 0966-9795/ © 2018 Elsevier Ltd. All rights reserved.

Intermetallics 105 (2019) 113–123

J. Lapin et al.

using hardness, tensile, compression and three-point bending tests. The samples of the in-situ composite are prepared by centrifugal casting of Ti-44.9Al-7.5Nb-4.9C-0.6Mo-0.1B (at.%) alloy. Finite element analysis (FEA) is used to calculate equivalent strains leading to fragmentation of coarse primary carbide particles and determine critical stress for crack initiation during compression and three-point bending tests, respectively. 2. Experimental material and procedure Samples of in-situ TiAl matrix composite with a diameter of 15 mm and length of 150 mm were prepared by induction melting of Ti-44.9Al7.5Nb-4.9C-0.6Mo-0.1B (at.%) alloy in graphite crucibles under argon atmosphere [39]. After holding at a temperature of 1680 °C for 60 s the melt was centrifugally cast into cold graphite mould at a rotation speed of 250 rpm [24,26]. The centrifugally cast (CC) samples were subjected to heat treatments consisting of hot isostatic pressing (HIP) at a temperature of 1250 °C and applied pressure of 200 MPa for 4 h in argon followed by annealing at a temperature of 900 °C for 25 h in air. Instrumented hardness measurements were performed on heat treated (HT) samples using an universal testing machine. Vickers hardness tests were carried out at an applied load of 50 N, holding time at the point of load application of 2 s and speed of load application of 15 N/s. Vickers microhardness measurements were performed at an applied load of 0.25 N and loading time of 10 s. Instrumented nanoindentation measurements of coexisting phases were carried out at an applied load of 0.01 N and holding time at the point of load application of 2 s on polished and slightly etched samples using a nanoindenter with Berkovich tip of the indenter. Threaded-head cylindrical tensile specimens with a gauge diameter of 6 mm and gauge length of 30 mm were prepared by machining from the HT samples. The surface of the gauge section was polished to a roughness better than 0.3 μm using a diamond paste. Room temperature tensile tests were carried out at an initial strain rate of 1 × 10−4 s−1 using universal testing machine. The elongation was measured by a contactless laser extensometer. Cylindrical compression specimens with a diameter of 8 mm and length of 12 mm were prepared by machining from the HT samples. The surface of the specimens was polished to a roughness better than 0.3 μm using a diamond paste. Room temperature compression tests were accomplished at an initial strain rate of 1 × 10−4 s−1 to a true strain of 16% using an universal testing machine. Compressive deformation was measured by a contactless laser extensometer. Compression tests to fracture were performed by Gleeble thermomechanical tester 3800 at room temperature using the procedure described elsewhere [40]. V-notch specimens with dimensions of 10 × 10 × 55 mm3, notch length of 2 mm and notch tip radius of 0.25 mm were prepared by wire electrical discharge machining from the HT samples. The surface of the specimens was machined to a roughness better than 0.8 μm. Room temperature quasi-static three-point bending tests were carried at a cross head speed of 8.33 × 10−6 m/s. The bending deformation was measured by a contactless laser extensometer. Two acoustic emission (AE) sensors were attached to the surface of the test specimen. The AE Vallen System sensors with a peak frequency of 600 kH were used. Acoustic emission events detected during testing were amplified by a pre-amplifier and analysed by the Vallen AE-Suite Software. The threshold was set to 50 dB and events of amplitude smaller than 50 dB were not recorded. Metallographic preparation of the samples consisted of standard grinding on abrasive papers and polishing on diamond pastes with various grain sizes up to 0.25 μm. Microstructural evaluation was performed by scanning electron microscopy (SEM), SEM in back scattered electron (BSE) mode, transmission electron microscopy (TEM), highresolution transmission electron microscopy (HRTEM), energy-dispersive X-ray spectroscopy (EDS) and X-ray diffraction (XRD) analysis. The XRD analysis was carried out by a diffractometer equipped with X-

Fig. 1. Initial mesh of compression specimen obtained by FEA.

ray tube with rotating Cu anode operating at 12 kW (wavelength of λ = 0.15418 nm). SEM samples were etched in a reagent of 150 ml H2O, 25 ml HNO3 and 10 ml HF. TEM samples were thinned mechanically to a thickness of about 40 μm and subsequently by ion milling to perforation. Size and volume fraction of coexisting phases were measured on digitalised micrographs using a computerised image analyser. The achieved microstructural data were treated statistically. The software Ansys Workbench was used for finite element analysis (FEA). A geometry model of the compression specimen was built up as a 3D cylinder with a diameter of 8 mm and length of 12 mm. Material model was built up using measured experimental data from the compression tests at room temperature. For 3D analysis, deformable hexahedron elements with a size of 0.43 mm were used to mesh the cylindrical compression specimen, as seen in Fig. 1. In total, the mesh was composed of 21 489 hexahedral elements, adding up to 90 552 nodes in total. A geometry model of the three-point bending specimen was built as a 3D cuboid with dimensions of 10 × 10 × 55 mm3 and notch length of 2 mm with notch root radius of 0.25 mm. One loading and two supporting pins were modelled as 3D cylinders. Material model was built up using measured experimental data from the three-point bending tests at room temperature. For calculations of strain distribution during three-point bending tests, deformable hexahedral elements with a size of 0.5 and 1 mm were used to mesh the cuboid specimen and cylinders, respectively, as seen in Fig. 2. In total, the mesh was composed of 45 621 hexahedral elements, adding up to 198 110 nodes in total. 3. Results and discussion 3.1. Microstructure characterisation The XRD analysis shows that the microstructure of the CC composite consists of γ(TiAl), α2(Ti3Al), H(Ti2AlC) and TiC phases, as seen in Fig. 3. The composite contains primary regular shaped (1), plate-like (2) and irregular shaped (3) carbide particles which are relatively uniformly distributed in the matrix composed of lamellar α2 + γ and single γ phase regions, as seen in Fig. 4. The TiC phase is identified in the core of many irregular shaped Ti2AlC particles, as shown in Fig. 4a and b. The origin of the TiC phase preserved in the microstructure of the

Fig. 2. Initial mesh of three-point bending specimen obtain by FEA. 114

Intermetallics 105 (2019) 113–123

J. Lapin et al.

→ α2 + γ + Ti2AlC. Fig. 4b shows clearly that the lamellar α2 + γ regions are formed in the vicinity of coarse carbide particles and single γ phase regions in the interdendritic region. Fig. 5 shows the initial microstructure of the test specimens prepared from the heat treated (HT) composite. The microstructure of the test specimens consists of coarse primary regular (1), plate-like (2) and irregular (3) shaped carbide particles which are relatively uniformly distributed in the γ matrix (4) containing fine secondary precipitates (5), as shown in Fig. 5a. The XRD pattern of the HT samples indicates existence of γ(TiAl) and Ti2AlC phases, as seen in Fig. 3. However, Table 1 and Fig. 6 show the existence of three chemically different phases in the HT in-situ composite such as γ(TiAl), Ti2AlC and TiC. The residual TiC phase can still be identified in the core of many irregular shaped Ti2AlC particles indicating that the applied HIP and annealing steps are insufficiently effective to achieve phase equilibria by a full solid phase transformation of TiC to Ti2AlC, as seen in Fig. 6a. The EDS map (Fig. 6b–f) and point analyses (Table 1) indicate that Nb substitutes Ti atoms in both Ti2AlC and TiC phases but Mo segregates preferentially to the TiAl matrix. Although the coarse primary carbide particles serve as heterogeneous nucleation sites for the β dendrites during solidification, the orientation relationship between the primary Ti2AlC and surrounding γ grains is found to be mainly random, as indicated in Fig. 5b. Besides the primary carbide particles, the TiAl matrix contains also secondary precipitates. Fig. 5c shows fine needle-like PeTi3AlC precipitates which were formed by dissolution of the α2 lamellae during the heat treatments. The inset in Fig. 5c indicates (001) (001)P and [120] [120]P orientation relationships between the γ matrix and the coherent P-type precipitates which confirm results reported recently by several authors [23,29,32,43]. Fig. 5d shows fine secondary precipitates distributed within the γ grains and along the γ/γ grain boundaries which are identified to belong to HeTi2AlC phase. The inset in Fig. 5d indicates (11¯0) (01¯11)H and [1¯1¯2] [2¯110]H orientation relationships between the γ matrix and the secondary H-type precipitates. Fig. 7 shows the typical examples of size distribution of the primary carbide particles before mechanical testing (BC) and after the compression tests (AC) in the vicinity of the contact area of the compression specimen with the pressure plate (region U) and centre of the compression specimen (region C) (see chapter 3.4). Both the length of major axis (Fig. 7a) and length of minor axis (Fig. 7b) can be fitted by a lognormal distribution function. Table 2 summarises values of mean length of major axis, mean length of minor axis, average shape factor and volume fraction of the regular, plate-like and irregular shaped primary carbide particles before mechanical testing (BC) and after the compression tests (AC) in the regions U and C. The shape factor F is calculated assuming relationship for circularity F = (4πA/P2), where A is the area of particle and P is the perimeter of particle. It is clear from this table that the irregular shaped carbide particles are represented by the highest volume fraction (10.1 vol.%) while the regular shaped (5.3 vol. %) and plate-like ones (4.7 vol.%) contribute nearly equally to the total

Fig. 3. The typical X-ray diffraction patterns of centrifugally cast (CC) and heat treated (HT) samples.

studied in-situ composite differs from that reported by Song et al. [13] for the composite prepared by powder metallurgy route through solid state phase transformations. The applied centrifugal casting leads to non-equlibrium solidification of the studied composite. According to the ternary TieAleC phase diagram reported by Witusiewicz et al. [41], the melt temperature of 1680 °C applied in the present study corresponds to L(liquid) + TiC1-x two phase region. During the cooling a new phase equilibrium is achieved according to a transformation pathway L + TiC1-x → L + Ti2AlC [41]. This reaction/transformation occurring at the L/TiC1-x interfaces leads to a full transformation of small TiC1-x particles to a regular shaped Ti2AlC. In the case of the coarse clustered TiC1-x particles, the reaction/transformation starts by the growth of solid Ti2AlC layer at the L/TiC1-x interfaces. When a continuous Ti2AlC layer is formed around the clusters, the further transformation is slowed-down significantly because of diffusion of alloying elements through the solid Ti2AlC layer separating the melt and remaining TiC1-x. Fast cooling rates associated with the centrifugal casting [6,42] lead to an incomplete transformation resulting in a preservation of small amount of untransformed TiC1-x phase in the core of many irregular shaped Ti2AlC particles. The plate like Ti2AlC particles nucleate and grow in the L + Ti2AlC phase region and their length is significantly influenced by cooling rate. The non-equlibrium solidification of the composite continues with the formation of β phase (Ti based solid solution with bcc crystal structure) which nucleates preferentially on the carbide particles according to phase transformation sequences: L + Ti2AlC → L + β + Ti2AlC → β + γ + Ti2AlC. The formation of the γ phase in the interdendritic regions results from Al segregation into remaining liquid during the β dendrite formation [26]. Further solid state phase transformation sequences can be characterised as follows: β + γ + Ti2AlC → α + β + γ + Ti2AlC → α + γ + Ti2AlC

Fig. 4. SEM micrographs showing microstructure of the CC composite: (a) Distribution of primary regular shaped (1), plate-like (2) and irregular shaped (3) carbide particles; (b) Distribution of lamellar α2 + γ and single γ regions in the matrix.

115

Intermetallics 105 (2019) 113–123

J. Lapin et al.

Fig. 5. The initial microstructure of the test specimens prepared from the HT composite: (a) SEM micrographs showing distribution of primary regular (1), plate-like (2) and irregular (3) shaped carbide particles in the γ matrix (4) containing fine secondary precipitates (5); (b) HRTEM micrograph of primary Ti2AlC particle/γ matrix interface and corresponding SAD pattern; (c) TEM micrograph and corresponding SAD pattern showing distribution of secondary PeTi3AlC particles formed in the γ matrix and along the γ/γ lamellar interfaces; (d) TEM micrograph and corresponding SAD pattern showing secondary HeTi2AlC precipitates formed within the γ matrix and along γ/γ grain boundaries.

Table 1 EDS analysis of chemical composition, nanohardness and indentation modulus of coexisting phases and HT in-situ composite. Phase

Element (at.%) Ti

Ti2AlC TiC γ(TiAl) Composite

42.3 41.8 41.9 42.1

± ± ± ±

1.8 0.6 0.3 0.2

Al

Nb

23.6 ± 0.6 1.2 ± 0.6 49.6 ± 0.5 44.9 ± 0.2

5.4 5.9 7.8 7.5

± ± ± ±

1.2 0.6 0.7 0.1

Mo

C

0.1 ± 0.1 – 0.7 ± 0.1 0.5 ± 0.1

28.6 ± 1.3 51.1 ± 0.5 – 4.9 ± 0.1

Nano-hardness (GPa)

Indentation modulus (GPa)

12.8 ± 2.1 28.4 ± 2.3 4.7 ± 1.1 –

231.1 ± 18.7 372.7 ± 23.5 165.8 ± 15.6 –

Fig. 6. EDS map analysis of the initial microstructure of test specimens. (a) SEM micrograph of primary irregular shaped Ti2AlC particles containing TiC phase; (b) carbon; (c) aluminium; (d) titanium; (e) niobium; (f) molybdenum.

116

Intermetallics 105 (2019) 113–123

J. Lapin et al.

Fig. 7. Log-normal size distribution of irregular shaped primary carbide particles: (a) Length of major axis; (b) Length of minor axis. BC – before mechanical testing, AC – after compression test. Regions U and C are designated in Fig. 9a.

composite and Vickers microhardness (HVmBC) of the TiAl matrix of the compression specimens before mechanical testing. The increase of the Vickers hardness HVAC and Vicker microhardness values HVmACU and HVmACC in the vicinity of the contact area of the specimen with the pressure plate (region U) and centre of the compression specimen (region C), respectively, clearly indicates work hardening of the in-situ composite and TiAl matrix during the compressive deformation to a true strain of 16%. The increase of the Vickers microhardness of the TiAl matrix strongly depends on the position in the compression specimen. Shu et al. [44] have observed similar increase of Vickers microhardness after compressive deformation of in-situ Ti2AlC/TiAl-xMn composites. However, the authors have not specified the positions of their microhardness measurements in the specimens. The measured indentation modulus of EI = 165 GPa for the studied in-situ composite (Table 3) is comparable with Young's modulus values ranging from 145 to 175 GPa reported for TiAl-based alloys [45].

Fig. 8. Dependence of compressive true stress on true strain. The numerically calculated data by FEA are shown in the figure.

3.3. Tensile tests The studied in-situ composite shows brittle tensile behaviour without plastic yielding. During room temperature tensile tests the stress increases nearly linearly with the strain up-to a tensile fracture. Table 4 summarises the measured room temperature mechanical properties of the studied in-situ composite. It should be noted that the average values listed in this table are calculated from five specimens tested at the same conditions for each type of mechanical test. The average ultimate tensile strength (UTS) is measured to be 392 MPa which is lower than that of 433 MPa measured by Li et al. [37] for Tie48Ale1C (at.%) in-situ composite with a lamellar α2 + γ matrix prepared by induction scull melting and casting into metallic mould but higher than that of 330 MPa measured by Lapin et el. [23] for Ti44.5Ale8Nb-5.2C-0.8Mo-0.1B (at.%) in-situ composite prepared by

volume fraction of the primary carbide particles (20.1 vol.%) which are relatively uniformly distributed in the γ matrix (79.9 vol.%). It should be noted that the volume fraction of the retained TiC particles is measured to be 0.8 vol.% in the HT composite. 3.2. Hardness and indentation modulus Table 1 summarises the average values of nanohardness and indentation modulus of the coexisting phases. It is clear that the nanohardness and indentation modulus of the retained TiC particles are significantly higher than those of the surrounding Ti2AlC phase or TiAl matrix. Table 3 shows the Vickers hardness (HVBC) of the in-situ

Fig. 9. (a) Longitudinal section of compression specimen tested to a true strain of 16%; (b) Corresponding 3D numerical modelling of equivalent plastic strains.

117

Intermetallics 105 (2019) 113–123

J. Lapin et al.

Table 2 Length of major axis, length of minor axis, shape factor and volume fraction of primary carbide particles before mechanical testing (BC) and after compression test (AC) in the regions U and C. Parameter

Region

Shape of particles Regular

Plate-like

Irregular

State

Length of major axis (μm) Length of minor axis (μm) Shape factor Volume fraction (vol.%)

U C U C U C

BC

AC

2.46 ± 0.04

1.96 1.83 1.25 1.02 0.70 0.68

1.55 ± 0.02 0.75 ± 0.01 5.4 ± 0.4

± ± ± ± ± ±

0.04 0.03 0.02 0.05 0.01 0.02

HVAC

HVmBC

HVmACU

HVmACC

EI (GPa)

310 ± 7

360 ± 9

301 ± 5

356 ± 9

433 ± 9

165 ± 5

AC

8.95 ± 0.06

3.55 2.73 0.85 0.78 0.43 0.47

1.05 ± 0.01 0.24 ± 0.01 4.6 ± 0.3

± ± ± ± ± ±

0.07 0.08 0.02 0.01 0.01 0.01

BC

AC

10.47 ± 0.55

2.54 1.84 1.13 0.86 0.62 0.59

7.19 ± 0.38 0.42 ± 0.02 10.1 ± 0.3

± ± ± ± ± ±

0.05 0.04 0.03 0.03 0.08 0.01

specimen tested to a true strain of 16% and corresponding 3D numerical calculations of equivalent plastic strains. The slightly barrelled shape of the specimen (Fig. 9a) corresponds quite well to the calculated one (Fig. 9b). The FEA indicates inhomogeneous distribution of the local calculated strains. The highest strain of about 20% is achieved in the central region of the compression specimen marked by a red rectangle (region C). The regions in the vicinity of the contact area of the specimen with the pressure plate of the testing machine marked by blue rectangles (region U) achieve the local strains ranging from 6 to 11%. The studied in-situ composite reinforced with bimodal sized coarse primary and fine secondary carbides is characterised by a strong binding at the TiAl/Ti2AlC and Ti2AlC/TiC interfaces. Fig. 10 shows the microstructure on a longitudinal section of the compression specimen deformed to a true strain of 16%. While the matrix shows no evidence of cracking, the compressive deformation in the regions U corresponding to a local strain of about 6% is connected with early stages of cracking of some coarse irregular shaped carbide particles and delamination at the carbide particle/matrix interfaces, as shown in Fig. 10a and b. The increase of the local strain to 20% in the region C leads to an intensive fragmentation of the irregular shaped and some plate-like carbide particles, as shown in Fig. 10c. In spite of relatively large local deformation, the initiation and propagation of cracks is constrained to take place within the carbide particles in the region C. However, Fig. 9a indicates that the compressive deformation in the vicinity of free surface of the compression sample is connected with the initiation and propagation of numerous large cracks which can lead even to the release of some grains from the specimen surface during testing. These cracks propagate within the TiAl matrix, carbide particles and along the γ/γ grain boundaries, as seen in Fig. 10d. Fig. 7 shows the typical examples of size distribution of the measured length of major axis (Fig. 7a) and length of minor axis (Fig. 7b) of the fragmented irregular shaped particles after the compression (AC) in the regions U and C. Table 2 summarises the results of statistical evaluation of mean length of major axis, mean length of minor axis and average shape factor of the fragmented regular, plate-like and irregular shaped carbide particles in the regions C and U. It is clear from this table that the fragmentation leads to a smaller size of the primary carbide particles and is more intensive in the region C (equivalent strain of 20%) compared to that in the regions U (equivalent strains ranging from 6 to 11%). The Vickers microhardness measurements of the TiAl matrix before

Table 3 Vickers hardness, Vickers microhardness and indentation modulus. HVBC

BC

induction melting in graphite crucibles and centrifugal casting into graphite mould. The measured UTS can be related to a critical stress leading to a crack initiation during three-point bending tests [23]. Leguillon et al. [46,47] have reported that the tensile strength of brittle materials is higher than flexural one because fewer randomly distributed defects in the material are involved in bending than in tension in case of specimens of the same size. However, such effect is minimised in the present work by selecting the size of tensile specimens which is close to the size of bending specimens stressed in tension. 3.4. Compression tests Two types of room-temperature compression tests are carried out in the present work: (i) compression tests to a true compressive strain of 16% and (ii) compression tests until the fracture of the specimens. Table 4 summarises average offset 0.2% compressive yield strength (CYS), ultimate compressive strength (UCS) and compressive strain to fracture (CS) calculated from five different compression specimens tested at the same conditions. The average CYS of 615 MPa is comparable with that of 617 MPa reported for Ti-44.5Al-7.9Nb-4.8C-0.6Mo0.1B (at.%) in-situ composite prepared by centrifugal casting [24] but lower than that of about 750 MPa reported by Ramaseshan et al. [48] for Tie45Ale5C (at.%) composite prepared by arc melting. The measured value of UCS of 1440 MPa is lower than those of 1679 or 2058 MPa reported recently for in-situ composites prepared by vacuum arc melting and spark plasma sintering, respectively [21,49]. The compressive true stress-true strain curves were calculated from the measured engineering compressive stress-strain data according to the procedure described elsewhere [40]. Fig. 8 shows the typical true stress-true strain compressive curve and numerical calculations using FEA. It is clear from this figure that the numerical calculations are in a very good agreement with the experimentally measured true stress-true strain curve. Fig. 9 shows a longitudinal section of the compression Table 4 Room-temperature mechanical properties. UTS (MPa)

CYS (MPa)

UCS (MPa)

CS (%)

K (MPa)

n

KS (MPam0.5)

392 ± 15

615 ± 8

1440 ± 25

28.3 ± 1.9

1585 ± 55

0.15 ± 0.02

13.5 ± 0.9

118

Intermetallics 105 (2019) 113–123

J. Lapin et al.

Fig. 10. BSE micrographs showing microstructure of longitudinal section of the compression specimen tested to a true strain of 16%: (a) Crack initiation within the coarse irregular shaped carbide particles in the region U, local strain of 6%; (b) Delamination at the carbide particle/matrix interface in the region U, local strain of 6%; (c) Fragmentation of irregular shaped carbide particles in the region C, local strain of 20%; (d) Crack propagation within the matrix and carbide particles in the vicinity of free surface of compression specimen, local strain of 19%.

attributed to an increase of dislocation density and mechanical twinning in the γ matrix [50–53]. However, Appel et al. [50] have observed nearly constant WHR up to a strain of about 8% in TiAl-based alloys. The studied in-situ composite shows a continuous decrease of the WHR with increasing strain which can be attributed to the initiation and propagation of cracks within the coarse carbide particles and along carbide particle/matrix interfaces, as seen in Fig. 10. Benitez et al. [54] have shown that the compressive strength of polycrystalline Ti2AlC depends strongly on grain size and increases from about 350 to about 1200 MPa by decreasing the grain size from 44 to 4.2 μm. Taking into account the procedure applied for in-situ composites described elsewhere [55], measured indentation modulus of the matrix of 165 GPa, indentation modulus of the carbide particles of 231.1 GPa and strain of 0.4% corresponding to the offset 0.2% yield strength of the composite, one can estimate a stress of about 930 MPa which is carried by the coarse carbide particles. This stress can be sufficiently high to cause plastic deformation of some favourably oriented primary Ti2AlC particles with the basal planes oriented at approximately 45° to load axis (soft orientations). The dislocations form and easily glide at such oriented basal planes at room temperature and arrange themselves either in pileups along slip planes or as dislocations walls normal to the basal planes [56–58]. The particles whose basal planes are parallel to the load axis can deform by kinking in the areas of maximum shear stress. The coarse Ti2AlC particles whose basal planes are perpendicular to the load axis can hardly deform plastically and undergo only elastic deformation. The fast decrease of the WHR in the region III at a true strain or true stress higher than 13.5% or 1170 MPa, respectively, indicates an intensive softening of the in-situ composite. Fig. 9a clearly shows that besides fragmentation of the coarse carbide particles and formation of cracks along the particle/matrix interfaces, this softening is connected with the initiation and propagation of cracks in the vicinity of the specimen free surface. The uniform plastic deformation stage of the compressive curve can be fitted by the Hollomon equation in the form.

Fig. 11. (a) Dependence of work hardening rate (WHR) on true strain. (b) The inset showing the dependence of WHR on true stress.

and after compression and compressive true stress-true strain curve show a clear work-hardening during compressive deformation of the studied in-situ composite. The work-hardening behaviour can be characterised by work hardening rate (WHR) Θ defined as

=

t t

(1)

where σt is the true stress and εt is the true strain. Fig. 11 shows dependence of calculated WHR on true strain and true stress. Three different regions can be well identified on the WHR curves, as illustrated in Fig. 11a and b. The region I corresponding to a true strain up to 1.9% or true stress up to 895 MPa is characterised by a fast decrease of the WHR. Appel et al. [50] have reported that a fast decrease of WHR with increasing compressive strain up to 2% in TiAl-based alloys is connected with non-uniform deformation starting in favourably oriented grains which leads to reduction of constraint stresses by localised flow but it is not conventional work hardening controlled by dislocation interactions. The region II observed at a true strain ranging from 1.9 to 13.5% or true stress from 895 to 1170 MPa is characterised by a gradual decrease of the WHR. The high WHR measured in the region II can be

t

=K

n t

(2)

where K is the material constant and n is the stress exponent. The linear regression analysis leads to K = 1585 MPa and n = 0.15, as summarised in Table 4. The measured value of the material constant of 1585 MPa is significantly lower than those of 2722 or 2945 MPa reported for TiAl or in-situ TiAl matrix composite reinforced with 6 vol.% 119

Intermetallics 105 (2019) 113–123

J. Lapin et al.

Fig. 12. SEM micrographs showing the typical fracture surface of the compression specimens: (a) Fracture surface with radial cracks along the grain boundaries; (b) Local plastic deformation of the fracture surface caused by sliding; (c) Bending and fracture of Ti2AlC particles; (d) Laminated tearing and bending of Ti2AlC particles; (e) Cracks formed in irregular shaped Ti2AlC arrested by TiC and fine secondary precipitates embedded in TiAl matrix; (f) Pull-out of carbide particles and delamination along carbide particle/matrix interfaces. P – precipitates.

of Ti2AlC particles fabricated by the combustion synthesis reaction, respectively [38]. The fracture morphology of the compression specimens is shown in Fig. 12. Fig. 12a indicates two different regions on the fracture surface and numerous radial cracks (perpendicular to the fracture surface) formed along the grain boundaries. Some local regions underwent an intensive plastic deformation which resulted in the formation of smooth local surfaces due to the sliding of two fracture surfaces during the compression, as seen in Fig. 12b. The coarse primary Ti2AlC particles show different level of plastic deformation or fracture morphology due to their different morphology and various crystallographic orientation to the load axis. Laminated tearing and crack arresting capability of the kink boundaries indicate an intrinsic plastic property of Ti2AlC that contributes to the reinforcement of the composite, as shown in Fig. 12c and d. The TiC particles preserved in the core of the irregular shaped carbides arrest the cracks and contribute to the enhancement of the damage tolerance of the in-situ composite, which is manifested as a

Fig. 13. The typical three-point bending load-deflection curve and corresponding AE events detected by AE sensors 1 and 2.

120

Intermetallics 105 (2019) 113–123

J. Lapin et al.

vicinity of the notch tip achieves a critical value corresponding to a flexural strength of the material [47]. Fig. 14 shows the tip region of the V-notch specimen after interrupted three-point bending test. The short cracks are initiated in the TiAl matrix, coarse carbide particles and along the carbide particle/matrix interfaces. The initiation of short cracks in the vicinity of the notch tip is observed at a load of 1690 N. This experimentally measured critical load is used to calculate stress distribution within the notched specimen and to determine a critical stress required for the crack initiation. Fig. 15 shows the distribution of numerically calculated equivalent (von-Mises) stresses σeq within the notched specimen corresponding to a load of 1690 N. A maximum local stress of 400 MPa calculated by FEA in the tip notch region is shown in detail in Fig. 15. This calculated maximum local stress is close to the measured average ultimate tensile strength of 392 MPa of the studied in-situ composite (Table 4). The fracture toughness of notched bending specimen can be calculated according to relationship [59]

Fig. 14. BSE micrograph showing crack initiation in the notch tip region during three-point bending test.

K S = 1.5

1.99 FLa0.5 × bw 2

a /w (1 a/ w )(2.15 (1 + 2a/ w )(1

3.93a/ w + 2.7(a/ w )2) 2a /w )1.5 (3)

where F is the fracture load, L is the span of the loading, b is the thickness of specimen, w is the width of specimen and a is the length of the notch. Since the ASTM specifications for fracture toughness testing concerning notch geometry and pre-cracking were not rigorously followed in the present study, the toughness values are denoted as KS. The measured value of fracture toughness for the studied in-situ composite of 13.4 MPam0.5 (Table 4) can be compared with those of 13.7 MPam0.5 for Ti-44.5Ale8Nb-4.8C-0.8Mo-0.1B in-situ composite prepared by centrifugal casting [24] or the values ranging from 7.3 to 17.8 MPam0.5 reported for TiAl/Ti2AlC in-situ composites prepared by reactive processing techniques [49,60,61]. The incorporation of Ti2AlC particles increases the fracture toughness of the in-situ composite compared to those of 8.0 MPam0.5 and 7.0 MPam0.5 reported by Rao et al. [62] and Wang et al. [61] for TiAl and Ti2AlC, respectively. Since the studied in–situ composite suffers from a lack of intrinsic toughening mechanisms due to absence of mobile dislocations during three-point bending testing, the improvement of the fracture toughness compared to that of TiAl can be related to crack deviation, zone shielding by microcrack toughening and bridging by coarse carbide particles. However, the effectiveness of these extrinsic toughening mechanisms depends strongly on the applied processing techniques, volume fraction, distribution and size of the carbide particles [24,48,49,60]. Fig. 16 shows brittle fracture mode of the in-situ composite after three-point bending test. The crack propagates along the crystallographic planes of some γ grains, grain boundaries and carbide particle/matrix interfaces. The coarse primary Ti2AlC particles with layered structure are fractured along their crystallographic planes. The laminated tearing of the primary carbides and the kink boundaries indicate that these particles contribute to the damage tolerance of the in-

Fig. 15. Equivalent (von-Mises) stress distribution at a critical load corresponding to a crack initiation in the notch tip region during three-point bending test.

crack deflection in the coarse Ti2AlC particles, as seen in Fig. 12e. The delamination at the particle/matrix interfaces leads to the pull out of some primary Ti2AlC from the matrix, which facilitate the final fracture of the compression specimens, as shown in Fig. 12f. 3.5. Three-point bending tests Fig. 13 shows the typical example of load-deflection curve of threepoint bending test with the standard V-notch specimen. The frequency of acoustic emission (AE) events increases with increasing load and deflection. Most of the AE events of amplitudes ranging from 70 to 90 dB occur at a deflection higher than 0.034 mm. The crack propagation is connected with the events of amplitudes up to 100 dB. The fracture of the specimen is observed at a fracture load of 1790 N and deflection of 0.041 mm. During the quasi-static three-point bending tests the notch creates a stress concentrator and increases local stresses in the specimen. The condition for crack initiation are fulfilled when local stress in the

Fig. 16. Typical fracture surface of the in-situ composite after three-point bending test, SEM. (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. P – precipitates.

121

Intermetallics 105 (2019) 113–123

J. Lapin et al.

situ composite when compared with a single TiAl matrix. The embedded carbides in the matrix retard the crack propagation and increase the fracture energy. As shown by Ramaseshan et al. [48] for Ti2AlC particles in lamellar α2+γ matrix, the increase of the fracture toughness of the in-situ composites results from crack trapping and its renucleation. The fracture toughness can be increased by increasing strength of the reinforcement [62]. Significantly harder TiC particles within the coarse irregular shaped Ti2AlC improve the fracture toughness of the studied in-situ composite by arresting and deviating the cracks formed within the coarse irregular shaped Ti2AlC particles.

[5] [6] [7]

[8]

4. Conclusions

[9]

The room temperature mechanical behaviour of the in-situ TiAl matrix composite reinforced with carbide particles prepared by centrifugal casting of Ti-44.9Al-7.5Nb-4.9C-0.6Mo-0.1B (at.%) alloy was studied. The achieved results can be summarised as follows:

[10] [11]

1. The initial microstructure of the test specimens prepared from the HIP-ed and annealed in-situ composite consists of the TiAl matrix reinforced with coarse primary regular, plate-like and irregular shaped carbide particles and fine secondary HeTi2AlC and PeTi3AlC precipitates. The cores of some irregular shaped primary Ti2AlC particles contain retained TiC phase. 2. The local equivalent strains in the compression specimens are numerically calculated using FEA and related to the size of fragmented primary carbide particles. The mean size of fragmented carbide particles decreases with increasing local equivalent strain in the compression specimen. 3. The in-situ composite shows plastic deformation of TiAl matrix and limited plastic deformation of coarse Ti2AlC particles during the early stages of the compressive deformation. The cracks preferentially initiate within the coarse irregular shaped carbide particles and at the carbide particle/matrix interfaces. 4. Three regions on work hardening rate curve characterised by different decrease of the WHR with increasing true strain or true stress can be related to the initial non-uniform deformation, plastic deformation of the γ matrix and favourably oriented Ti2AlC particles, cracking of coarse carbide particles and propagation of cracks in the vicinity of the specimen free surface. 5. The stress distribution in the notched specimen and critical stress leading to crack initiation in the notch tip region are numerically calculated using FEA. The calculated critical stress is close to the measured ultimate tensile strength of the studied in-situ composite. 6. The brittle fracture of the in-situ composite during three-point bending tests includes crack deviation, zone shielding by microcrack toughening, carbide fragmentation, delamination on the carbide particle/matrix interfaces and pull-out of the carbide particles from the TiAl matrix.

[12] [13] [14] [15] [16] [17]

[18]

[19] [20]

[21] [22] [23] [24]

Acknowledgements

[25]

This work was financially supported by the Slovak Research and Development Agency under the contract APVV-15-0660 and the Slovak Grant Agency for Science under the contract VEGA 2/0125/16.

[26]

References

[27]

[1] B.P. Bewlay, S. Nag, A. Suzuki, M.J. Weimer, TiAl alloys in commercial aircraft engines, Mater. A. T. High. Temp. 33 (2016) 549–559, https://doi.org/10.1080/ 09603409.2016.1183068. [2] Y.-W. Kim, S.-L. Kim, Advances in gammalloy materials–processes–application technology: successes, dilemmas, and future, JOM 70 (2018) 553–560, https://doi. org/10.1007/s11837-018-2747-x. [3] J. Lapin, M. Nazmy, Microstructure and creep properties of a cast intermetallic Ti46Al-2W-0.5Si alloy for gas turbine applications, Mater. Sci. Eng., A 380 (2004) 298–307, https://doi.org/10.1016/j.msea.2004.05.011. [4] T. Tetsui, T. Kobayashi, T. Ueno, H. Harada, Consideration of the influence of

[28] [29]

[30]

122

contamination from oxide crucibles on TiAl cast material, and the possibility of achieving low-purity TiAl precision cast turbine wheels, Intermetallics 31 (2012) 274–281, https://doi.org/10.1016/j.intermet.2012.07.019. M.T. Jovanović, B. Dimčić, I. Bobić, S. Zec, V. Maksimović, Microstructure and mechanical properties of precision cast TiAl turbocharger wheel, J. Mater. Process. Technol. 167 (2005) 14–21, https://doi.org/10.1016/j.jmatprotec.2005.03.019. P.X. Fu, X.H. Kang, Y.C. Ma, K. Liu, D.Z. Li, Y.Y. Li, Centrifugal casting of TiAl exhaust valves, Intermetallics 16 (2008) 130–138, https://doi.org/10.1016/j. intermet.2007.08.007. J. Aguilar, A. Schievenbusch, O. Kättlitz, Investment casting technology for production of TiAl low pressure turbine blades - process engineering and parameter analysis, Intermetallics 19 (2011) 757–761, https://doi.org/10.1016/j.intermet. 2010.11.014. H. Clemens, S. Mayer, Intermetallic titanium aluminides in aerospace applications – processing, microstructure and properties, Mater. A. T. High. Temp. 33 (2016) 560–570, https://doi.org/10.1080/09603409.2016.1163792. M. Bünck, T. Stoyanov, J. Schievenbusch, H. Michels, A. Gußfeld, Titanium aluminide casting technology development, JOM 69 (2017) 2565–2570, https://doi. org/10.1007/s11837-017-2534-0. K. Kamyshnykova, J. Lapin, Grain refinement of cast peritectic TiAl-based alloy by solid-state phase transformations, Kov. Mater. 56 (2018) 277–287, https://doi.org/ 10.4149/km_2018_5_277. Y.-W. Kim, Strength and ductility in TiAl alloys, Intermetallics 6 (1998) 623–628, https://doi.org/10.1016/S0966-9795(98)00037-5. Y.-W. Kim, S.-L. Kim, Effects of microstructure and C and Si additions on elevated temperature creep and fatigue of gamma TiAl alloys, Intermetallics 53 (2014) 92–101, https://doi.org/10.1016/j.intermet.2014.04.006. X.J. Song, H.Z. Cui, N. Hou, N. Wei, Y. Han, J. Tian, Q. Song, Lamellar structure and effect of Ti2AlC on properties of prepared in-situ TiAl matrix composites, Ceram. Int. 42 (2016) 13586–13592, https://doi.org/10.1016/j.ceramint.2016.05.152. X. Song, H. Cui, Y. Han, N. Hou, N. Wei, L. Ding, Q. Song, Effect of carbon reactant on microstructures and mechanical properties of TiAl/Ti2AlC composites, Mater. Sci. Eng., A 684 (2017) 406–412, https://doi.org/10.1016/j.msea.2016.12.069. Y. Tan, R. Chen, H. Fang, Y. Liu, H. Ding, Y. Su, J. Guo, H. Fu, Microstructure evolution and mechanical properties of TiAl binary alloys added with SiC fibers, Intermetallics 98 (2018) 69–78, https://doi.org/10.1016/j.intermet.2018.04.018. S. Cui, C. Cui, J. Xie, S. Liu, J. Shi, Carbon fibers coated with graphene reinforced TiAl alloy composite with high strength and toughness, Sci. Rep. 8 (2018) 1–8, https://doi.org/10.1038/s41598-018-20799-y. R. Chen, Y. Tan, H. Fang, L. Luo, H. Ding, Y. Su, J. Guo, H. Fu, Macro/microstructure evolution and mechanical properties of Ti33.3Al alloys by adding WC particles, Mater. Sci. Eng., A 725 (2018) 171–180, https://doi.org/10.1016/j.msea. 2018.04.025. W. Li, Y. Yang, M. Li, J. Liu, D. Cai, Q. Wei, C. Yan, Y. Shi, Enhanced mechanical property with refined microstructure of a novel γ-TiAl/TiB2 metal matrix composite (MMC) processed via hot isostatic press, Mater. Des. 141 (2018) 57–66, https://doi. org/10.1016/j.matdes.2017.12.026. P. Liu, X. Han, D. Sun, Q. Wang, Development and application of a ternary Ti-Al-N interatomic potential for Ti2AlN/TiAl composite, J. Alloy. Comp. 745 (2018) 63–74, https://doi.org/10.1016/j.jallcom.2018.02.168. H. Fang, R. Chen, Y. Yang, Y. Su, H. Ding, J. Guo, H. Fu, Role of graphite on microstructural evolution and mechanical properties of ternary TiAl alloy prepared by arc melting method, Mater. Des. 156 (2018) 300–310, https://doi.org/10.1016/j. matdes.2018.06.048. R. Chen, H. Fang, X. Chen, Y. Su, H. Ding, J. Guo, H. Fu, Formation of TiC/Ti2AlC and α2+γ in in-situ TiAl composites with different solidification paths, Intermetallics 81 (2017) 9–15, https://doi.org/10.1016/j.intermet.2017.02.025. H. Fang, R. Chen, X. Chen, Y. Yang, Y. Su, H. Ding, J. Guo, Effect of Ta element on microstructure formation and mechanical properties of high-Nb TiAl alloys, Intermetallics 104 (2019) 43–51, https://doi.org/10.1016/j.intermet.2018.10.017. J. Lapin, M. Štamborská, T. Pelachová, O. Bajana, Fracture behaviour of cast in-situ TiAl matrix composite reinforced with carbide particles, Mater. Sci. Eng., A 721 (2018) 1–7, https://doi.org/10.1016/j.msea.2018.02.077. J. Lapin, A. Klimová, Z. Gabalcová, T. Pelachová, O. Bajana, M. Štamborská, Microstructure and mechanical properties of cast in-situ TiAl matrix composites reinforced with (Ti,Nb)2AlC particles, Mater. Des. 133 (2017) 404–415, https://doi. org/10.1016/j.matdes.2017.08.012. A. Klimová, J. Lapin, T. Pelachová, Characterization of TiAl based alloys with various content of carbon, IOP Conf. Ser. Mater. Sci. Eng. (2017), https://doi.org/ 10.1088/1757-899X/179/1/012038. J. Lapin, K. Kamyshnykova, Processing, microstructure and mechanical properties of in-situ Ti3Al+TiAl matrix composite reinforced with Ti2AlC particles prepared by centrifugal casting, Intermetallics 98 (2018) 34–44, https://doi.org/10.1016/j. intermet.2018.04.012. T. Cegan, I. Szurman, Thermal stability and precipitation strengthening of fully lamellar Ti-45Al-5Nb-0.2B-0.75C alloy, Kov. Mater. 55 (2017) 421–430, https:// doi.org/10.4149/km_2017_6_421. T. Čegan, I. Szurman, M. Kursa, J. Holešinský, J. Vontorová, Preparation of TiAlbased alloys by induction melting in graphite crucibles, Kov. Mater. 53 (2015) 69–78, https://doi.org/10.4149/km_2015_2_69. L. Wang, H. Gabrisch, U. Lorenz, F.P. Schimansky, A. Schreyer, A. Stark, F. Pyczak, Nucleation and thermal stability of carbide precipitates in high Nb containing TiAl alloys, Intermetallics 66 (2015) 111–119, https://doi.org/10.1016/j.intermet. 2015.07.001. L. Wang, U. Lorenz, M. Münch, A. Stark, F. Pyczak, Influence of alloy composition and thermal history on carbide precipitation in γ-based TiAl alloys, Intermetallics

Intermetallics 105 (2019) 113–123

J. Lapin et al. 89 (2017) 32–39, https://doi.org/10.1016/j.intermet.2017.05.006. [31] L. Wang, C. Zenk, A. Stark, P. Felfer, H. Gabrisch, M. Göken, U. Lorenz, F. Pyczak, Morphology evolution of Ti3AlC carbide precipitates in high Nb containing TiAl alloys, Acta Mater. 137 (2017) 36–44, https://doi.org/10.1016/j.actamat.2017.07. 018. [32] H. Gabrisch, A. Stark, F.P. Schimansky, L. Wang, N. Schell, U. Lorenz, F. Pyczak, Investigation of carbides in Ti-45Al-5Nb-xC alloys (0 ≤ x ≤ 1) by transmission electron microscopy and high energy-XRD, Intermetallics 33 (2013) 44–53, https:// doi.org/10.1016/j.intermet.2012.09.023. [33] E. Schwaighofer, B. Rashkova, H. Clemens, A. Stark, S. Mayer, Effect of carbon addition on solidification behavior, phase evolution and creep properties of an intermetallic β-stabilized γ-TiAl based alloy, Intermetallics 46 (2014) 173–184, https://doi.org/10.1016/j.intermet.2013.11.011. [34] J. Rössler, A.G. Evans, The effect of reinforcement size on the creep strength of intermetallic matrix composites, Mater. Sci. Eng., A 153 (1992) 438–443, https:// doi.org/10.1016/0921-5093(92)90234-R. [35] J. Rössler, G. Bao, A.G. Evans, The effects of diffusional relaxation on the creep strength of composites, Acta Metall. Mater. 39 (1991) 2733–2738, https://doi.org/ 10.1016/0956-7151(91)90090-N. [36] Z. Wu, R. Hu, T. Zhang, F. Zhang, H. Kou, J. Li, Understanding the role of carbon atoms on microstructure and phase transformation of high Nb containing TiAl alloys, Mater. Char. 124 (2017) 1–7, https://doi.org/10.1016/j.matchar.2016.12. 008. [37] M. Li, S. Xiao, L. Xiao, L. Xu, J. Tian, Y. Chen, Effects of carbon and boron addition on microstructure and mechanical properties of TiAl alloys, J. Alloy. Comp. 728 (2017) 206–221, https://doi.org/10.1016/j.jallcom.2017.08.211. [38] S. Shu, F. Qiu, S. Lü, S. Jin, Q. Jiang, Phase transitions and compression properties of Ti2AlC/TiAl composites fabricated by combustion synthesis reaction, Mater. Sci. Eng., A 539 (2012) 344–348, https://doi.org/10.1016/j.msea.2012.01.108. [39] K. Kamyshnykova, J. Lapin, Vacuum induction melting and solidification of TiAlbased alloy in graphite crucibles, Vacuum 154 (2018) 218–226, https://doi.org/10. 1016/j.vacuum.2018.05.017. [40] M. Štamborská, J. Lapin, Effect of anisotropic microstructure on high-temperature compression deformation of CoCrFeNi based complex concentrated alloy, Kov. Mater. 55 (2017) 369–378, https://doi.org/10.4149/km_2017_6_369. [41] V.T. Witusiewicz, B. Hallstedt, A.A. Bondar, U. Hecht, S.V. Sleptsov, T.Y. Velikanova, Thermodynamic description of the Al-C-Ti system, J. Alloy. Comp. 623 (2015) 480–496, https://doi.org/10.1016/j.jallcom.2014.10.119. [42] L. Yang, L.H. Chai, Y.F. Liang, Y.W. Zhang, C.L. Bao, S.B. Liu, J.P. Lin, Numerical simulation and experimental verification of gravity and centrifugal investment casting low pressure turbine blades for high Nb-TiAl alloy, Intermetallics 66 (2015) 149–155, https://doi.org/10.1016/j.intermet.2015.07.006. [43] E. Schwaighofer, P. Staron, B. Rashkova, A. Stark, N. Schell, H. Clemens, S. Mayer, In situ small-angle X-ray scattering study of the perovskite-type carbide precipitation behavior in a carbon-containing intermetallic TiAl alloy using synchrotron radiation, Acta Mater. 77 (2014) 360–369, https://doi.org/10.1016/j.actamat. 2014.06.017. [44] S. Shu, F. Qiu, B. Xing, S. Jin, Y. Wang, Q. Jiang, Study of effect of Mn addition on the mechanical properties of Ti2AlC/TiAl composites through first principles study and experimental investigation, Intermetallics 28 (2012) 65–70, https://doi.org/ 10.1016/j.intermet.2012.03.053. [45] F. Appel, J.D.H. Paul, M. Oehring, Gamma Titanium Aluminide Alloys: Science and Technology, Wiley-VCH Verlag & Co. KGaA, Weinheim, Germany, 2011, https:// doi.org/10.1002/9783527636204.

[46] D. Leguillon, É. Martin, M.C. Lafarie-Frenot, Flexural vs. tensile strength in brittle materials, Compt. Rendus Mec. 343 (2015) 275–281, https://doi.org/10.1016/j. crme.2015.02.003. [47] D. Leguillon, Strength or toughness? A criterion for crack onset at a notch, Eur. J. Mech. Solid. 21 (2002) 61–72, https://doi.org/10.1016/S0997-7538(01)01184-6. [48] R. Ramaseshan, A. Kakitsuji, S.K. Seshadri, N.G. Nair, H. Mabuchi, H. Tsuda, T. Matsui, K. Morii, Microstructure and some properties of TiAl-Ti2AlC composites produced by reactive processing, Intermetallics 7 (1999) 571–577, https://doi.org/ 10.1016/S0966-9795(98)00069-7. [49] C. Yang, F. Wang, T. Ai, J. Zhu, Microstructure and mechanical properties of in situ TiAl/Ti2AlC composites prepared by reactive hot pressing, J. Eur. Ceram. Soc. 40 (2014) 8165–8171, https://doi.org/10.1016/j.ceramint.2014.01.012. [50] F. Appel, U. Sparka, R. Wagner, Work hardening and recovery of gamma base titanium aluminides, Intermetallics 7 (1999) 325–334, https://doi.org/10.1016/ S0966-9795(98)00109-5. [51] B. Viguier, Dislocation densities and strain hardening rate in some intermetallic compounds, Mater. Sci. Eng., A 349 (2003) 132–135, https://doi.org/10.1016/ S0921-5093(02)00785-2. [52] J.D.H. Paul, F. Appel, R. Wagner, The compression behaviour of niobium alloyed γtitanium aluminides, Acta Mater. 46 (1998) 1075–1085, https://doi.org/10.1016/ S1359-6454(97)00332-7. [53] F. Kauffmann, T. Bidlingmaier, G. Dehm, A. Wanner, H. Clemens, On the origin of acoustic emission during room temperature compressive deformation of a γ-TiAl based alloy, Intermetallics 8 (2000) 823–830, https://doi.org/10.1016/S09669795(00)00025-X. [54] R. Benitez, H. Gao, M. O'Neal, P. Lovelace, G. Proust, M. Radovic, Effects of microstructure on the mechanical properties of Ti2AlC in compression, Acta Mater. 143 (2018) 130–140, https://doi.org/10.1016/j.actamat.2017.10.019. [55] J. Lapin, Microstructure and mechanical properties of directionally solidified γ/γ′-α eutectic composites, Int. J. Mater. Prod. Technol. 18 (2003) 255–281, https://doi. org/10.1504/IJMPT.2003.003594. [56] R. Benitez, W.H. Kan, H. Gao, M. O'Neal, G. Proust, M. Radovic, Room temperature stress-strain hysteresis in Ti2AlC revisited, Acta Mater. 105 (2016) 294–305, https://doi.org/10.1016/j.actamat.2015.12.004. [57] M. Shamma, E.N. Caspi, B. Anasori, B. Clausen, D.W. Brown, S.C. Vogel, V. Presser, S. Amini, O. Yeheskel, M.W. Barsoum, In situ neutron diffraction evidence for fully reversible dislocation motion in highly textured polycrystalline Ti2AlC samples, Acta Mater. 98 (2015) 51–63, https://doi.org/10.1016/j.actamat.2015.07.023. [58] M.W. Barsoum, M. Radovic, Elastic and mechanical properties of the MAX phases, Annu. Rev. Mater. Res. 41 (2011) 195–227, https://doi.org/10.1146/annurevmatsci-062910-100448. [59] R. Cao, M.X. Lei, J.H. Chen, J. Zhang, Effects of loading rate on damage and fracture behavior of TiAl alloys, Mater. Sci. Eng., A 465 (2007) 183–193, https://doi.org/ 10.1016/j.msea.2007.02.026. [60] T.T. Ai, F. Wang, X.M. Feng, M.M. Ruan, Microstructural and mechanical properties of dual Ti3AlC2-Ti2AlC reinforced TiAl composites fabricated by reaction hot pressing, Ceram. Int. 40 (2014) 9947–9953, https://doi.org/10.1016/j.ceramint. 2014.02.092. [61] P. Wang, B. Mei, X. Hong, W. Zhou, Synthesis of Ti2AlC by hot pressing and its mechanical and electrical properties, Trans. Nonferrous Metals Soc. China 17 (2007) 1001–1004, https://doi.org/10.1016/S1003-6326(07)60215-5. [62] K.T.V. Rao, R.O. Ritchie, High-temperature fracture and fatigue resistance of a ductile β-TiNb reinforced γ-TiAl intermetallic composite, Acta Metall. 46 (1998) 4167–4180, https://doi.org/10.1002/9783527636204.

123