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The nucleation of microcracks under tensile stress in multi-phase high Nb-containing TiAl alloys
Intermetallics 106 (2019) 13–19
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The nucleation of microcracks under tensile stress in multi-phase high Nbcontaining TiAl alloys
T
Bin Zhu, Xiangyi Xue, Hongchao Kou∗, Ruifeng Dong, Jinshan Li State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, 710072, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Titanium aluminides Microcrack nucleation ECCI Heterogeneous deformation Plastic anisotropy B2 phase
A series of bending tests was conducted on specimens of different TiAl alloys to investigate the nucleation of microcracks under tensile stress. The distribution of microcracks and the deformation traces around them were characterized to study the role of plastic deformation intersecting with microstructural features in the microcrack nucleation. Results show that the microcrack nucleation plays an important role in the fracture of multiphase high Nb-containing TiAl alloys. The microcracks tend to nucleate at colony boundaries in near fully lamellar and fully lamellar alloys, while they prefer B2/γ phase boundaries in (B2+γ) alloy. The heterogeneous deformation in different lamellar colonies was caused by plastic anisotropy of the lamellar structure. It leads to the nucleation of microcracks at colony boundaries. Meanwhile, due to the stress concentration at lamellar interfaces in the transversal deformation mode, a few microcracks nucleate inside lamellar colonies which have an orientation angle Φ near to 90°. The existence of brittle B2 phase leads to strain incompatibility and facilitates microcrack nucleation at B2/γ boundaries.
1. Introduction After decades of R&D activities, gamma TiAl-based alloys have been successfully utilized as aero-engine low-pressure turbine blades [1], turbocharger wheels [2] and exhaust valves [3] in aerospace and automotive industries. In recent years, high Nb-containing TiAl alloys have attracted significant attention due to the improvement in service temperature and hot workability [4,5]. However, the extreme brittleness at ambient temperature could lead to a high propensity for catastrophic fracture, impeding the production and the application of these materials. The fracture processes of materials are generally divided into two stages, including the microcrack nucleation and the crack propagation. Fracture mechanics normally focus on crack propagation process and consider fracture toughness (the resistance to crack growth) as the main criterion for assessing the fracture tendency of materials [6]. The nucleation of microcracks in undamaged material is not very well understood. It was evidenced that the ductility and fracture toughness of the two-phase TiAl alloys are correlated to the microstructure in two opposite manners [7,8], i.e. the lamellar structure has a great fracture toughness and a poor ductility, whereas the duplex structure exhibits a low toughness with a high tensile elongation at ambient temperature. This phenomenon reveals that the tensile fracture of the two-phase TiAl
∗
alloys is controlled by crack nucleation rather than crack propagation. Thus, it is important to investigate the mechanism of the microcrack nucleation in TiAl alloys. Bieler et al. have systematically studied the damage nucleation at grain boundaries in the single γ phase TiAl alloys [9–11]. They have correlated the microcrack nucleation with the heterogeneous deformation at grain boundaries, and successfully established fracture initiation parameters on the basis of slip transfer across grain boundaries to predict the crack nucleation. However, there is still a lack of systematic investigation on the mechanisms of microcrack nucleation in the multi-phase high Nb-containing TiAl alloys. The studies for elucidating the role of plastic deformation intersecting with microstructural features in the microcrack nucleation are still necessary. In this work, the bending tests were conducted on TiAl alloys with different microstructures. During the bending tests, the upper part of the bending specimen was under compressive stress, while the lower part was under tensile stress. The microcracks were generated on the bottom of bending specimens. Since the top of the bending specimens was subjected to compressive stress, the damage tolerance of those specimens was better compared with uniaxial tensile tested specimens. The microcracks nucleated during bending tests could be maintained. The distribution of microcracks and the deformation traces around them were characterized to study the microcrack nucleation in multi-
Corresponding author. E-mail address: [email protected] (H. Kou).
https://doi.org/10.1016/j.intermet.2018.12.006 Received 4 September 2018; Received in revised form 5 December 2018; Accepted 9 December 2018 0966-9795/ © 2018 Elsevier Ltd. All rights reserved.
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investigation on these three alloys is reasonable. Grey acicular borides and bright yttrium oxide particles are observed in all three alloys. Detailed description of the microstructural parameters can be found in our previous work [13]. Prismatic bars with dimensions (thickness, b × height, h × length, l) of 2 × 4 × 35 mm3 were prepared for the bending test, as illustrated in Fig. 2. The specimens were wired-cut from TiAl alloys and grounded to the final dimensions. The loading mode of four point bending tests was also displayed in this figure. The tip radius of the four cylindrical indenters is 1.5 mm. The span between the two upper indenters, S1, is 20 mm. The support span, S2, is 30 mm. The bending tests were performed with a compressive speed of the two upper indenters of 0.01 mm/min at room temperature. Two strain gauges were glued on the central region of the top and bottom surfaces, respectively. The compressive load, P, and the longitudinal strains on the top, εc (compression), and the bottom, εt (tension), were recorded during bending tests. The bottom of each specimen was electropolished before mechanical tests to investigate the microcrack nucleation under tensile stress. The microstructures of raw materials and the distribution of nucleated microcracks on the fractured bending specimens were characterized by a VEGA II LMH scanning electron microscope (SEM) using the back scattered electron (BSE) mode. A detailed examination of the deformation traces in the vicinity of the microcracks was performed on a Zeiss GeminiSEM 500 using electron channeling contrast imaging (ECCI). Simultaneously, the crystal orientation in the same location was obtained by electron backscatter diffraction (EBSD). A step size of 1 μm was used for the EBSD measurement. The γ phase was indexed as normal face-centered cubic structure.
Table 1 Microstructure parameters of the multi-phase high Nb-containing TiAl alloys. Alloy
Nominal composition (at.%)
Processing
Microstructure
#1 #2 #3
Tie45Al-8.5Nb-0.2W-0.2B-0.02Y Tie45Al-8.5Nb-0.2W-0.2B-0.02Y Ti-42.5Ale8Nb-0.2W-0.2B-0.1Y
Cast Cast-HT Cast-forge-HT
NFL FL γ grains + B2 phase
phase high Nb-containing TiAl alloys.
2. Material and methods Three alloys with different microstructures were designed to investigate the effect of lamellar structure and phase constituents on the nucleation of microcracks in multi-phase high Nb-containing TiAl alloys. The compositions and processing of these alloys are listed in Table 1. Alloy #1 was sampled from an industrial PAM ingot in the ascast state. It has a near fully lamellar (NFL) structure with γ grains (in dark contrast) and residual B2 phase (in light contrast) located at α2/γ lamellar colony boundaries, as shown in Fig. 1(a). Alloy #2 was sampled from the same PAM ingot and heat treated (HT) at 1300 °C (α phase region) [12] for 10 h to remove B2 phase at colony boundaries. It displays a fully lamellar (FL) microstructure (see Fig. 1 (b)). To obtain alloy #3, a cylindrical rod was cut from a VAR ingot and canned-forged twice at 1260 °C. Then a heat treatment of 1168 °C/100 h (B2+γ field) was performed. As shown in Fig. 1(c), the microstructure of alloy #3 consists of equiaxial γ grains and B2 phase. This is similar to the microstructure at colony boundaries in the NFL alloy. It should be noted that all the three alloys have a high Nb content. Although there is a difference of Al and Y rates among these three alloys, the phase constituents and intrinsic mechanical properties are still similar. Thus the
Fig. 1. BSE images of different TiAl alloys: (a) alloy #1, NFL; (b) alloy #2, FL; (c) alloy #3, (B2+γ). 14
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Fig. 2. The specimen geometry and loading mode of the bending tests.
Fig. 3. Schematic image showing (a) the normal stress on the cross section during bending test and (b) the cross sectional area.
The strains εc and εt are correlated with the compressive load P. Both sides of the formulas can be divided by an infinitesimally small ΔP [15]. Equation (3) become
3. Results 3.1. Tensile stress-strain relationship obtained from flexural data During bending tests, pure bending with a radius of curvature ρ occurs on the specimen (Fig. 3(a)). The part of the bending specimen between two upper indenters is subjected to a constant bending moment M, which equals to P∙l1/2. In this part, only normal stress parallel to the length direction is generated on arbitrary cross sections as pure bending condition. The uniaxial stress, σ, is a function of the longitudinal strain, ε, and the function may differ between the tension and compression conditions. The compressive stress on the top surface is denoted as σc, and the tensile stress on the bottom is represented by σt. It is assumed that the neutral axis, where ε is 0, is located at y = 0. Thus the equilibrium of the normal stress and the bending moment on the cross section (Fig. 3(b)) are yt
∫ σdA = 0
σt =
−yc
∫ σydA = M (1)
−yc
Since ε = y/ρ and dA = bdy, the equilibrium equations become εt
εt
∫ σdε = 0 −εc
and
∫ σεdε = M /bρ2 . −εc
(2)
Then differentiate the equilibrium equation (2) by curvature ρ according to Leibnitz’ Rule [14], it can be deduced that
σt =
dM (εc + εt ) + 2M (dεc + dεt ) . bh2dεt
(
dεc dP
+
dεt dP
) + (ε
bh2
dεt dP
c
⎞ + εt )/2⎟ ⎠, (4)
which related the tensile stress-strain curve on the bottom with the load-strain data obtained from bending tests. The typical load-strain curves of different TiAl alloys obtained from the bending tests are displayed in Fig. 4 (a). Here the compressive strain εc and the tensile strain, εt, are given with the absolute values. The tensile stress-strain curves calculated from equation (4) are presented in Fig. 4 (b), implying the tensile deformation on the bottom of the bending specimens. The apparent elastic modulus derived from the elastic range is about 180 MPa, which agrees well with the measured values reported in other studies [16,17]. The curves deviate slightly from linearity before yielding, which reveals the significant pre-yield plastic deformation, especially for the NFL and FL alloys. The pre-yield plasticity is a representative phenomenon observed during deformation of TiAl alloys. This could be attributed to the plastic anisotropy and heterogeneous strain distribution [18,19]. The results show that the (B2+γ) alloy exhibits the lowest ductility among the three alloys, whereas the NFL and FL alloys have slightly higher fracture elongation.
yt
and
l1 ⎛⎜P ⎝
3.2. Characterization of microcrack nucleation (3)
Fig. 5(a, b and c) shows the montages of BSE images obtained from 15
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Fig. 4. (a) Typical load-strain curves obtained from the bending tests and (b) the calculated tensile stress-strain curves corresponding to the bottom of these specimens.
narrow mechanical twins, and the nucleation of microcracks is accompanied by mechanical twinning.
the fractured bending specimens of alloy #1, #2 and #3. The microcracks were highlighted in red lines. The tensile stress is along horizontal direction in these figures hereafter. It is seen that the distribution of microcracks is macroscopically heterogeneous on the three alloys. This is resulted from the slightly accidented surface after the electropolishing. The quantities of microcracks in NFL and FL alloys are less than these in the (B2+γ) alloy. It reveals the significant role of microcrack nucleation in the fracture process of TiAl alloys, since the (B2+γ) alloy has the worst ductility and the most microcracks at the same time. It should be noted that most of the microcracks are perpendicular to or inclined with the tensile direction in the NFL and FL alloys, whereas some of the cracks in the (B2+γ) alloy are parallel to the tensile direction. Those perpendicular or inclined cracks nucleate and propagate as the effect of the normal component of tensile stress. The parallel cracks in the (B2+γ) alloy may be resulted from the alteration of local stress state. This will be discussed in the latter. In addition, the cracks in the NFL and FL alloys are normally located at colony boundaries. Only a few cracks nucleate inside lamellar colonies and propagate along lamellar direction. Most of the cracks in the (B2+γ) alloy are correlated to B2 phase. The cracks are generally located in B2 phase or at B2/γ interfaces. Only a few cracks propagate along γ grain boundaries or through γ grains. Fig. 5(d)-(f) shows the ECCI image of the microcracks in alloy #1, #2 and #3, respectively. The microcrack in the NFL alloy is located at colony boundary. It is related to the B2 phase (indicated by white arrows). A part of this crack propagates into the lamellar colony. This may be correlated to the irregular γ grains and yttrium oxide particles (cavities indicated by the circles). Fig. 5(e) shows two microcracks nucleate on the FL alloy. One of the cracks is located at the colony boundary, and the other is situated in the lamellar colony and propagates along the lamellae direction. The crack in the (B2+γ) alloy propagates along γ grain boundary. It may be generated due to the white yttrium oxide particles (indicated by circles). There is no correlation between the borides and the microcrack nucleation in the three TiAl alloys. Straight deformation traces are observed in γ lamellae and γ grains in the vicinity of microcracks (Fig. 5(d)-(f)). Similar deformation traces have been discovered in γ grains by Simkin et al. [20] and in γ lamellae of TNM alloy by Lertner et al. [21]. These were identified as mechanical twins with a thickness of a few hundred nanometers. It is known that mechanical twinning is the primary deformation mode in high Nb-containing TiAl alloys at room temperature due to the relatively low stacking-fault energy [22]. In addition, we have observed numerous mechanical twins with a thickness of dozens of nanometers in the pre-yielding tensile tested TiAl specimens (undergoing work). Hence, the deformation traces observed by ECCI here are recognized as
4. Discussion 4.1. Plastic anisotropy of lamellar structure As has been investigated previously [23,24], the lamellar structure of TiAl alloys exhibits a significant plastic anisotropy due to the strong barrier strength of lamellar interface to dislocation glide and twin growth. The deformation behavior of lamellar colony is strongly dependent on the orientation angle Φ of lamellae interface with respect to the tensile axis. It is insensitive to the crystallographic orientation for a given orientation angle Φ. The deformation of lamellar structure can be classified into three modes by Lebensohn et al. [25]: (a) the mixed mode for Φ close to zero. The slip planes of the activated slip systems are inclined with the lamellar interface, whereas the slip directions are normally parallel to it. It results in a net strain parallel to the lamellar interface and a high yield stress; (b) the longitudinal mode for Φ close to 45°. In this deformation mode, the dislocations and twins do not cross the lamellar interfaces, leading to a relatively low yield stress; (c) the transversal deformation mode for Φ close to 90°. The slip planes and slip directions are both inclined to the lamellar interface, and it leads to the maximum yield stress. The plastic anisotropy of lamellar colonies gives rise to strain incompatibility and stress concentration at the colony boundaries. This may be the major reason for the microcrack nucleation in TiAl alloys with lamellar structure. Fig. 6 shows the ECCI of microcracks on the FL alloy. The spatial orientation of the lamellar interface and the orientation angle Φ between lamellar interface and the tensile direction are calculated from EBSD measurements as illustrated in our previous work [26] and presented in Fig. 6. As shown in Fig. 6(a), the orientation angles Φ of the two colonies are 23° and 75°, respectively. The deformation modes of the two colonies are most likely to be mixed mode and transversal mode, respectively. Both colonies are hard to deform, and the strain of the two colonies is in different directions. Thus, the microcrack nucleates at the colony boundary due to the strain incompatibility and stress concentration. As shown in Fig. 6(b), a microcrack nucleates inside the lamellar colony. Numerous inclined twins can be observed in γ lamellae in this colony. The origin of this microcrack is attributed to two factors. Firstly, the lamellar colony is in transversal deformation mode. The lamellar interfaces act as strong barriers to strain transfer and results in stress concentration at lamellar interfaces. Secondly, the lamellar interface is almost perpendicular to the specimen surface. This is favorable for the crack nucleation. It should be noted that there are still some 16
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Fig. 5. Montages of BSE images and ECCI in SEM showing the distribution and morphology of microcracks on the specimens of (a) and (d) alloy #1; (b) and (e) alloy #2; (c) and (f) alloy #3. The tensile stress is along horizontal direction.
Fig. 7(a). It shows that different twinning systems are activated at different regions of γ grain-A. The crystallographic elements of these twins were determined according to EBSD measurements [27]. The Euler angles of grain-A are (133.5, 29.1, 86.5°). The twinning planes of these twins activated in region-b are (11–1)γ and (−111)γ, and they are (111)γ and (−111)γ in region-c. In the γ phase with ordered L10 structure, only one distinct shear direction 1/6 < 11–2] exists for each {111} twinning plane. Since the γ phase was indexed as simple fcc structure in the EBSD experiment, as illustrated in Section 2, the c-axis of γ phase and the shear direction of these twins can't be confirmed. Thus, there are three possible orientations of c-axis in grain-A. The Schmid factors of the twinning systems activated in grain-A are calculated to explain the deformation mechanism, as listed in Table 2. The results include the Schmid factors for all three possible orientations. The indices of twining system listed in the column correspond to
residual B2 particles at colony boundaries after heat treatment, as indicated by white arrows in Fig. 6. These B2 particles may be also another important factor for microcrack nucleation.
4.2. Effect of B2 phase ECCI of the microcrack in the (B2+γ) alloy is shown in Fig. 7. The traces of all {111}γ planes for each adjacent γ grains are calculated according to the EBSD results and indicated in Fig. 7(a). The microcrack nucleates at B2/γ phase boundary and propagates along it. Numerous intersected twins are observed in γ grains around this microcrack, meanwhile there is no deformation generated in B2 phase. The heterogeneous strain distribution between B2 phase and γ phase are considered to be the main reason for the nucleation of this microcrack. Fig. 7(b) and (c) are magnified view of region-b and -c indicated in 17
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Fig. 6. ECCI of microcracks on the bending specimen of alloy #2. The ellipses show projected appearance of a unit circle on the lamellar interface: the major axis is the lamellar trace, the dotted line indicates the interface normal, and the shaded half is below the surface. The orientation angle Φ of different colonies are labelled in the images.
orientation I (the EBSD result). The results of Schmid factors in the same row correspond to the same deformation trace in Fig. 7. Since the twinning shear in γ phase is not reversible, the twinning under tensile stress is expected to be activated as the Schmid factor is positive and forbidden when it is negative. However, none of the three possible orientations fully satisfies the criterion, for which all the Schmid factors of the twinning on (11–1)γ, (−111)γ and (111)γ planes should be positive. It means the shear of some twins is in the opposite direction of the resolved shear stress, and the twinning in the γ grain cannot be predicted based on Schmid factors. This kind of phenomenon has also been observed in the deformation of TiAl alloys with lamellar structure [28], which was attributed to the alternation of local stress state during deformation. The grey-scale contrast in bright B2 phase reveals the existence of ω phase [29], which restricts dislocation mobility and increases the hardness of B2 phase. Due to the difference of mechanical properties between B2 and γ phases, the deformation concentrates in γ
Table 2 Twinning system and corresponding Schmid factor for three possible orientations of grain A indicated in Fig. 7. Twinning system (for orientation I)
1/6[11–2](111) 1/6[-11-2](-111) 1/6[1-1-2](1–11) 1/6[112](11–1)
Schmid factor I
II
III
−0.056 0.244 0.464 −0.061
−0.125 −0.106 −0.223 0.208
0.181 −0.138 −0.241 −0.147
grains and leads to strain discontinuity and stress concentration at B2/γ phase boundaries. The microcrack may nucleate as the strain accommodation is insufficient. Thus, it can be concluded that the existence of hard and brittle B2 phase promotes crack nucleation owing to the strain incompatibility at B2/γ phase boundaries.
Fig. 7. ECCI of microcracks on the bending specimen of alloy #3. The traces of all {111}γ planes are indicated for each γ grain. 18
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5. Conclusions [9]
Bending tests were carried out at room temperature to investigate the mechanisms of microcrack nucleation of multi-phase TiAl alloys. The results are summarized as follows:
[10] [11]
(1) Microcrack nucleation plays an important role in the fracture of TiAl alloys. The microcracks tend to nucleate at colony boundaries the TiAl alloys with NFL and FL microstructures, while they prefer to nucleate at the B2/γ phase boundaries in the (B2+γ) TiAl alloy. (2) The strain incompatibility and stress concentration caused by plastic anisotropy of the lamellar structure results in microcrack nucleation at colony boundaries. Simultaneously, the crack can nucleate inside lamellar colonies with an orientation angle Φ near to 90° due to the stress concentration at lamellae interfaces in the transversal deformation mode. (3) The existence of B2 phase changes local microscopic stress state. The heterogeneous deformation between B2 and γ phases results in strain discontinuity at B2/γ boundaries and facilitates microcrack nucleation.
[12] [13] [14] [15]
[16] [17] [18]
[19]
Acknowledgement [20]
The authors would like to acknowledge the National Natural Science Foundation of China [grant number 51571162] and the State Key Laboratory of Solidification Processing, NWPU, China [grant number 121-TZ-2015] for their financial supports.
[21]
[22]
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The nucleation of microcracks under tensile stress in multi-phase high Nbcontaining TiAl alloys
T
Bin Zhu, Xiangyi Xue, Hongchao Kou∗, Ruifeng Dong, Jinshan Li State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, 710072, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Titanium aluminides Microcrack nucleation ECCI Heterogeneous deformation Plastic anisotropy B2 phase
A series of bending tests was conducted on specimens of different TiAl alloys to investigate the nucleation of microcracks under tensile stress. The distribution of microcracks and the deformation traces around them were characterized to study the role of plastic deformation intersecting with microstructural features in the microcrack nucleation. Results show that the microcrack nucleation plays an important role in the fracture of multiphase high Nb-containing TiAl alloys. The microcracks tend to nucleate at colony boundaries in near fully lamellar and fully lamellar alloys, while they prefer B2/γ phase boundaries in (B2+γ) alloy. The heterogeneous deformation in different lamellar colonies was caused by plastic anisotropy of the lamellar structure. It leads to the nucleation of microcracks at colony boundaries. Meanwhile, due to the stress concentration at lamellar interfaces in the transversal deformation mode, a few microcracks nucleate inside lamellar colonies which have an orientation angle Φ near to 90°. The existence of brittle B2 phase leads to strain incompatibility and facilitates microcrack nucleation at B2/γ boundaries.
1. Introduction After decades of R&D activities, gamma TiAl-based alloys have been successfully utilized as aero-engine low-pressure turbine blades [1], turbocharger wheels [2] and exhaust valves [3] in aerospace and automotive industries. In recent years, high Nb-containing TiAl alloys have attracted significant attention due to the improvement in service temperature and hot workability [4,5]. However, the extreme brittleness at ambient temperature could lead to a high propensity for catastrophic fracture, impeding the production and the application of these materials. The fracture processes of materials are generally divided into two stages, including the microcrack nucleation and the crack propagation. Fracture mechanics normally focus on crack propagation process and consider fracture toughness (the resistance to crack growth) as the main criterion for assessing the fracture tendency of materials [6]. The nucleation of microcracks in undamaged material is not very well understood. It was evidenced that the ductility and fracture toughness of the two-phase TiAl alloys are correlated to the microstructure in two opposite manners [7,8], i.e. the lamellar structure has a great fracture toughness and a poor ductility, whereas the duplex structure exhibits a low toughness with a high tensile elongation at ambient temperature. This phenomenon reveals that the tensile fracture of the two-phase TiAl
∗
alloys is controlled by crack nucleation rather than crack propagation. Thus, it is important to investigate the mechanism of the microcrack nucleation in TiAl alloys. Bieler et al. have systematically studied the damage nucleation at grain boundaries in the single γ phase TiAl alloys [9–11]. They have correlated the microcrack nucleation with the heterogeneous deformation at grain boundaries, and successfully established fracture initiation parameters on the basis of slip transfer across grain boundaries to predict the crack nucleation. However, there is still a lack of systematic investigation on the mechanisms of microcrack nucleation in the multi-phase high Nb-containing TiAl alloys. The studies for elucidating the role of plastic deformation intersecting with microstructural features in the microcrack nucleation are still necessary. In this work, the bending tests were conducted on TiAl alloys with different microstructures. During the bending tests, the upper part of the bending specimen was under compressive stress, while the lower part was under tensile stress. The microcracks were generated on the bottom of bending specimens. Since the top of the bending specimens was subjected to compressive stress, the damage tolerance of those specimens was better compared with uniaxial tensile tested specimens. The microcracks nucleated during bending tests could be maintained. The distribution of microcracks and the deformation traces around them were characterized to study the microcrack nucleation in multi-
Corresponding author. E-mail address: [email protected] (H. Kou).
https://doi.org/10.1016/j.intermet.2018.12.006 Received 4 September 2018; Received in revised form 5 December 2018; Accepted 9 December 2018 0966-9795/ © 2018 Elsevier Ltd. All rights reserved.
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investigation on these three alloys is reasonable. Grey acicular borides and bright yttrium oxide particles are observed in all three alloys. Detailed description of the microstructural parameters can be found in our previous work [13]. Prismatic bars with dimensions (thickness, b × height, h × length, l) of 2 × 4 × 35 mm3 were prepared for the bending test, as illustrated in Fig. 2. The specimens were wired-cut from TiAl alloys and grounded to the final dimensions. The loading mode of four point bending tests was also displayed in this figure. The tip radius of the four cylindrical indenters is 1.5 mm. The span between the two upper indenters, S1, is 20 mm. The support span, S2, is 30 mm. The bending tests were performed with a compressive speed of the two upper indenters of 0.01 mm/min at room temperature. Two strain gauges were glued on the central region of the top and bottom surfaces, respectively. The compressive load, P, and the longitudinal strains on the top, εc (compression), and the bottom, εt (tension), were recorded during bending tests. The bottom of each specimen was electropolished before mechanical tests to investigate the microcrack nucleation under tensile stress. The microstructures of raw materials and the distribution of nucleated microcracks on the fractured bending specimens were characterized by a VEGA II LMH scanning electron microscope (SEM) using the back scattered electron (BSE) mode. A detailed examination of the deformation traces in the vicinity of the microcracks was performed on a Zeiss GeminiSEM 500 using electron channeling contrast imaging (ECCI). Simultaneously, the crystal orientation in the same location was obtained by electron backscatter diffraction (EBSD). A step size of 1 μm was used for the EBSD measurement. The γ phase was indexed as normal face-centered cubic structure.
Table 1 Microstructure parameters of the multi-phase high Nb-containing TiAl alloys. Alloy
Nominal composition (at.%)
Processing
Microstructure
#1 #2 #3
Tie45Al-8.5Nb-0.2W-0.2B-0.02Y Tie45Al-8.5Nb-0.2W-0.2B-0.02Y Ti-42.5Ale8Nb-0.2W-0.2B-0.1Y
Cast Cast-HT Cast-forge-HT
NFL FL γ grains + B2 phase
phase high Nb-containing TiAl alloys.
2. Material and methods Three alloys with different microstructures were designed to investigate the effect of lamellar structure and phase constituents on the nucleation of microcracks in multi-phase high Nb-containing TiAl alloys. The compositions and processing of these alloys are listed in Table 1. Alloy #1 was sampled from an industrial PAM ingot in the ascast state. It has a near fully lamellar (NFL) structure with γ grains (in dark contrast) and residual B2 phase (in light contrast) located at α2/γ lamellar colony boundaries, as shown in Fig. 1(a). Alloy #2 was sampled from the same PAM ingot and heat treated (HT) at 1300 °C (α phase region) [12] for 10 h to remove B2 phase at colony boundaries. It displays a fully lamellar (FL) microstructure (see Fig. 1 (b)). To obtain alloy #3, a cylindrical rod was cut from a VAR ingot and canned-forged twice at 1260 °C. Then a heat treatment of 1168 °C/100 h (B2+γ field) was performed. As shown in Fig. 1(c), the microstructure of alloy #3 consists of equiaxial γ grains and B2 phase. This is similar to the microstructure at colony boundaries in the NFL alloy. It should be noted that all the three alloys have a high Nb content. Although there is a difference of Al and Y rates among these three alloys, the phase constituents and intrinsic mechanical properties are still similar. Thus the
Fig. 1. BSE images of different TiAl alloys: (a) alloy #1, NFL; (b) alloy #2, FL; (c) alloy #3, (B2+γ). 14
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Fig. 2. The specimen geometry and loading mode of the bending tests.
Fig. 3. Schematic image showing (a) the normal stress on the cross section during bending test and (b) the cross sectional area.
The strains εc and εt are correlated with the compressive load P. Both sides of the formulas can be divided by an infinitesimally small ΔP [15]. Equation (3) become
3. Results 3.1. Tensile stress-strain relationship obtained from flexural data During bending tests, pure bending with a radius of curvature ρ occurs on the specimen (Fig. 3(a)). The part of the bending specimen between two upper indenters is subjected to a constant bending moment M, which equals to P∙l1/2. In this part, only normal stress parallel to the length direction is generated on arbitrary cross sections as pure bending condition. The uniaxial stress, σ, is a function of the longitudinal strain, ε, and the function may differ between the tension and compression conditions. The compressive stress on the top surface is denoted as σc, and the tensile stress on the bottom is represented by σt. It is assumed that the neutral axis, where ε is 0, is located at y = 0. Thus the equilibrium of the normal stress and the bending moment on the cross section (Fig. 3(b)) are yt
∫ σdA = 0
σt =
−yc
∫ σydA = M (1)
−yc
Since ε = y/ρ and dA = bdy, the equilibrium equations become εt
εt
∫ σdε = 0 −εc
and
∫ σεdε = M /bρ2 . −εc
(2)
Then differentiate the equilibrium equation (2) by curvature ρ according to Leibnitz’ Rule [14], it can be deduced that
σt =
dM (εc + εt ) + 2M (dεc + dεt ) . bh2dεt
(
dεc dP
+
dεt dP
) + (ε
bh2
dεt dP
c
⎞ + εt )/2⎟ ⎠, (4)
which related the tensile stress-strain curve on the bottom with the load-strain data obtained from bending tests. The typical load-strain curves of different TiAl alloys obtained from the bending tests are displayed in Fig. 4 (a). Here the compressive strain εc and the tensile strain, εt, are given with the absolute values. The tensile stress-strain curves calculated from equation (4) are presented in Fig. 4 (b), implying the tensile deformation on the bottom of the bending specimens. The apparent elastic modulus derived from the elastic range is about 180 MPa, which agrees well with the measured values reported in other studies [16,17]. The curves deviate slightly from linearity before yielding, which reveals the significant pre-yield plastic deformation, especially for the NFL and FL alloys. The pre-yield plasticity is a representative phenomenon observed during deformation of TiAl alloys. This could be attributed to the plastic anisotropy and heterogeneous strain distribution [18,19]. The results show that the (B2+γ) alloy exhibits the lowest ductility among the three alloys, whereas the NFL and FL alloys have slightly higher fracture elongation.
yt
and
l1 ⎛⎜P ⎝
3.2. Characterization of microcrack nucleation (3)
Fig. 5(a, b and c) shows the montages of BSE images obtained from 15
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Fig. 4. (a) Typical load-strain curves obtained from the bending tests and (b) the calculated tensile stress-strain curves corresponding to the bottom of these specimens.
narrow mechanical twins, and the nucleation of microcracks is accompanied by mechanical twinning.
the fractured bending specimens of alloy #1, #2 and #3. The microcracks were highlighted in red lines. The tensile stress is along horizontal direction in these figures hereafter. It is seen that the distribution of microcracks is macroscopically heterogeneous on the three alloys. This is resulted from the slightly accidented surface after the electropolishing. The quantities of microcracks in NFL and FL alloys are less than these in the (B2+γ) alloy. It reveals the significant role of microcrack nucleation in the fracture process of TiAl alloys, since the (B2+γ) alloy has the worst ductility and the most microcracks at the same time. It should be noted that most of the microcracks are perpendicular to or inclined with the tensile direction in the NFL and FL alloys, whereas some of the cracks in the (B2+γ) alloy are parallel to the tensile direction. Those perpendicular or inclined cracks nucleate and propagate as the effect of the normal component of tensile stress. The parallel cracks in the (B2+γ) alloy may be resulted from the alteration of local stress state. This will be discussed in the latter. In addition, the cracks in the NFL and FL alloys are normally located at colony boundaries. Only a few cracks nucleate inside lamellar colonies and propagate along lamellar direction. Most of the cracks in the (B2+γ) alloy are correlated to B2 phase. The cracks are generally located in B2 phase or at B2/γ interfaces. Only a few cracks propagate along γ grain boundaries or through γ grains. Fig. 5(d)-(f) shows the ECCI image of the microcracks in alloy #1, #2 and #3, respectively. The microcrack in the NFL alloy is located at colony boundary. It is related to the B2 phase (indicated by white arrows). A part of this crack propagates into the lamellar colony. This may be correlated to the irregular γ grains and yttrium oxide particles (cavities indicated by the circles). Fig. 5(e) shows two microcracks nucleate on the FL alloy. One of the cracks is located at the colony boundary, and the other is situated in the lamellar colony and propagates along the lamellae direction. The crack in the (B2+γ) alloy propagates along γ grain boundary. It may be generated due to the white yttrium oxide particles (indicated by circles). There is no correlation between the borides and the microcrack nucleation in the three TiAl alloys. Straight deformation traces are observed in γ lamellae and γ grains in the vicinity of microcracks (Fig. 5(d)-(f)). Similar deformation traces have been discovered in γ grains by Simkin et al. [20] and in γ lamellae of TNM alloy by Lertner et al. [21]. These were identified as mechanical twins with a thickness of a few hundred nanometers. It is known that mechanical twinning is the primary deformation mode in high Nb-containing TiAl alloys at room temperature due to the relatively low stacking-fault energy [22]. In addition, we have observed numerous mechanical twins with a thickness of dozens of nanometers in the pre-yielding tensile tested TiAl specimens (undergoing work). Hence, the deformation traces observed by ECCI here are recognized as
4. Discussion 4.1. Plastic anisotropy of lamellar structure As has been investigated previously [23,24], the lamellar structure of TiAl alloys exhibits a significant plastic anisotropy due to the strong barrier strength of lamellar interface to dislocation glide and twin growth. The deformation behavior of lamellar colony is strongly dependent on the orientation angle Φ of lamellae interface with respect to the tensile axis. It is insensitive to the crystallographic orientation for a given orientation angle Φ. The deformation of lamellar structure can be classified into three modes by Lebensohn et al. [25]: (a) the mixed mode for Φ close to zero. The slip planes of the activated slip systems are inclined with the lamellar interface, whereas the slip directions are normally parallel to it. It results in a net strain parallel to the lamellar interface and a high yield stress; (b) the longitudinal mode for Φ close to 45°. In this deformation mode, the dislocations and twins do not cross the lamellar interfaces, leading to a relatively low yield stress; (c) the transversal deformation mode for Φ close to 90°. The slip planes and slip directions are both inclined to the lamellar interface, and it leads to the maximum yield stress. The plastic anisotropy of lamellar colonies gives rise to strain incompatibility and stress concentration at the colony boundaries. This may be the major reason for the microcrack nucleation in TiAl alloys with lamellar structure. Fig. 6 shows the ECCI of microcracks on the FL alloy. The spatial orientation of the lamellar interface and the orientation angle Φ between lamellar interface and the tensile direction are calculated from EBSD measurements as illustrated in our previous work [26] and presented in Fig. 6. As shown in Fig. 6(a), the orientation angles Φ of the two colonies are 23° and 75°, respectively. The deformation modes of the two colonies are most likely to be mixed mode and transversal mode, respectively. Both colonies are hard to deform, and the strain of the two colonies is in different directions. Thus, the microcrack nucleates at the colony boundary due to the strain incompatibility and stress concentration. As shown in Fig. 6(b), a microcrack nucleates inside the lamellar colony. Numerous inclined twins can be observed in γ lamellae in this colony. The origin of this microcrack is attributed to two factors. Firstly, the lamellar colony is in transversal deformation mode. The lamellar interfaces act as strong barriers to strain transfer and results in stress concentration at lamellar interfaces. Secondly, the lamellar interface is almost perpendicular to the specimen surface. This is favorable for the crack nucleation. It should be noted that there are still some 16
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Fig. 5. Montages of BSE images and ECCI in SEM showing the distribution and morphology of microcracks on the specimens of (a) and (d) alloy #1; (b) and (e) alloy #2; (c) and (f) alloy #3. The tensile stress is along horizontal direction.
Fig. 7(a). It shows that different twinning systems are activated at different regions of γ grain-A. The crystallographic elements of these twins were determined according to EBSD measurements [27]. The Euler angles of grain-A are (133.5, 29.1, 86.5°). The twinning planes of these twins activated in region-b are (11–1)γ and (−111)γ, and they are (111)γ and (−111)γ in region-c. In the γ phase with ordered L10 structure, only one distinct shear direction 1/6 < 11–2] exists for each {111} twinning plane. Since the γ phase was indexed as simple fcc structure in the EBSD experiment, as illustrated in Section 2, the c-axis of γ phase and the shear direction of these twins can't be confirmed. Thus, there are three possible orientations of c-axis in grain-A. The Schmid factors of the twinning systems activated in grain-A are calculated to explain the deformation mechanism, as listed in Table 2. The results include the Schmid factors for all three possible orientations. The indices of twining system listed in the column correspond to
residual B2 particles at colony boundaries after heat treatment, as indicated by white arrows in Fig. 6. These B2 particles may be also another important factor for microcrack nucleation.
4.2. Effect of B2 phase ECCI of the microcrack in the (B2+γ) alloy is shown in Fig. 7. The traces of all {111}γ planes for each adjacent γ grains are calculated according to the EBSD results and indicated in Fig. 7(a). The microcrack nucleates at B2/γ phase boundary and propagates along it. Numerous intersected twins are observed in γ grains around this microcrack, meanwhile there is no deformation generated in B2 phase. The heterogeneous strain distribution between B2 phase and γ phase are considered to be the main reason for the nucleation of this microcrack. Fig. 7(b) and (c) are magnified view of region-b and -c indicated in 17
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Fig. 6. ECCI of microcracks on the bending specimen of alloy #2. The ellipses show projected appearance of a unit circle on the lamellar interface: the major axis is the lamellar trace, the dotted line indicates the interface normal, and the shaded half is below the surface. The orientation angle Φ of different colonies are labelled in the images.
orientation I (the EBSD result). The results of Schmid factors in the same row correspond to the same deformation trace in Fig. 7. Since the twinning shear in γ phase is not reversible, the twinning under tensile stress is expected to be activated as the Schmid factor is positive and forbidden when it is negative. However, none of the three possible orientations fully satisfies the criterion, for which all the Schmid factors of the twinning on (11–1)γ, (−111)γ and (111)γ planes should be positive. It means the shear of some twins is in the opposite direction of the resolved shear stress, and the twinning in the γ grain cannot be predicted based on Schmid factors. This kind of phenomenon has also been observed in the deformation of TiAl alloys with lamellar structure [28], which was attributed to the alternation of local stress state during deformation. The grey-scale contrast in bright B2 phase reveals the existence of ω phase [29], which restricts dislocation mobility and increases the hardness of B2 phase. Due to the difference of mechanical properties between B2 and γ phases, the deformation concentrates in γ
Table 2 Twinning system and corresponding Schmid factor for three possible orientations of grain A indicated in Fig. 7. Twinning system (for orientation I)
1/6[11–2](111) 1/6[-11-2](-111) 1/6[1-1-2](1–11) 1/6[112](11–1)
Schmid factor I
II
III
−0.056 0.244 0.464 −0.061
−0.125 −0.106 −0.223 0.208
0.181 −0.138 −0.241 −0.147
grains and leads to strain discontinuity and stress concentration at B2/γ phase boundaries. The microcrack may nucleate as the strain accommodation is insufficient. Thus, it can be concluded that the existence of hard and brittle B2 phase promotes crack nucleation owing to the strain incompatibility at B2/γ phase boundaries.
Fig. 7. ECCI of microcracks on the bending specimen of alloy #3. The traces of all {111}γ planes are indicated for each γ grain. 18
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5. Conclusions [9]
Bending tests were carried out at room temperature to investigate the mechanisms of microcrack nucleation of multi-phase TiAl alloys. The results are summarized as follows:
[10] [11]
(1) Microcrack nucleation plays an important role in the fracture of TiAl alloys. The microcracks tend to nucleate at colony boundaries the TiAl alloys with NFL and FL microstructures, while they prefer to nucleate at the B2/γ phase boundaries in the (B2+γ) TiAl alloy. (2) The strain incompatibility and stress concentration caused by plastic anisotropy of the lamellar structure results in microcrack nucleation at colony boundaries. Simultaneously, the crack can nucleate inside lamellar colonies with an orientation angle Φ near to 90° due to the stress concentration at lamellae interfaces in the transversal deformation mode. (3) The existence of B2 phase changes local microscopic stress state. The heterogeneous deformation between B2 and γ phases results in strain discontinuity at B2/γ boundaries and facilitates microcrack nucleation.
[12] [13] [14] [15]
[16] [17] [18]
[19]
Acknowledgement [20]
The authors would like to acknowledge the National Natural Science Foundation of China [grant number 51571162] and the State Key Laboratory of Solidification Processing, NWPU, China [grant number 121-TZ-2015] for their financial supports.
[21]
[22]
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