Orientation Dependence of Stress Rupture Properties of a Ni-based Single Crystal Superalloy at 760 C

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1 J. Mater. Sci. Technol., 2012, 28(3), Orientation Dependence of Stress Rupture Properties of a Ni-based Single Crystal Superalloy at 760 C Shaohua Zhang, Dong Wang, Jian Zhang and Langhong Lou Superalloys Division, Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang , China [Manuscript received February 21, 2011, in revised form April 18, 2011] The orientation dependence of creep rupture lives of a single crystal superalloy at 760 C/760 MPa was investigated. The orientations of the specimens tested were about 30 away from [001]. The results showed that specimens with orientations on the [001]-[011] boundary had the longest rupture life. The deformation of these specimen were controlled by a/2<110> slip and a few stacking faults with two orientations were observed. On the other hand, specimens with orientations near the [001]-[011] boundary or on the [001]- [ 111] boundary showed short rupture lives, and stacking faults with single orientation were observed in these specimens. The rupture properties and the deformation mechanisms were discussed based on the dislocation pattern and the calculated Schmid factors for different specimens. KEY WORDS: Single crystal superalloy; Stress rupture; Orientation 1. Introduction Turbine blades in single crystal (SC) form are the most important development in the materials technology of advanced turbines, because SC superalloys exhibit attractive elevated temperature capability [1,2]. However, SC superalloys are known as anisotropy [3], especially the creep properties tested at the condition of low temperature and high stress [4,5]. The previous work had studied the creep behavior of Mar-M200 at 760 C/689 MPa [4,5] and Mar- M247 at 774 C/724 MPa [6]. The results showed that the primary creep strain of specimens with orientation within 25 from [001] increased in the following order: [001], [001]-[011] boundary, between [001]-[011] and [001]-[ 111] boundaries, and [001]-[ 111] boundary. Large primary creep strains led to poor rupture properties. Transmission electron microscopy (TEM) observation confirmed that the extent of primary creep strain was related to the a/3<112> dislocations cutting into γ. The specimen with high Corresponding author. Prof., Ph.D.; Tel./Fax: ; address: jianzhang@imr.ac.cn (J. Zhang). Schimd factor on <112> {111} systems exhibited high primary creep strain. Similar work was carried out using CMSX-4 with misorientations within 20 at 750 C/750 MPa [7]. It was found that the stacking fault shear was absent in specimen near [013] and this specimen exhibited excellent creep properties. The reason was explained as that there was no appropriate a/2<110> dislocations in the matrix channel to nucleate a/3<112> dislocations. When the misorientation from [001] exceeded 25, the rupture properties of specimens near the [001]- [ 111] boundary were better than those near the [001]- [011] boundary [5,6]. It was reported that due to the operation of <112> dislocation, specimens with orientation near the [001]-[011] boundary required large rotation to reach the [001]-[ 111] boundary, which resulted in the cutting of γ by <112> type dislocations and poor creep properties [6]. Other works showed that specimen near [ 111] orientation had long creep life due to the low Schmid factor [8,9]. However, specimen near [011] orientation had low creep life due to the a/3<112> dislocations cutting into γ during the primary creep stage and the deformation on the single slip system [10,11]. Meanwhile, another work illu-

2 230 S.H. Zhang et al.: J. Mater. Sci. Technol., 2012, 28(3), Fig. 1 (a) The relationship between the creep lives tested at 760 C/760 MPa and orientations; (b) lattice rotation for specimens A, B and C strated that specimens near the [001]-[ 111] boundary were absent of stacking fault since the two operating <110>{111} slip systems on different {111} planes could not lead to interaction forming the <112> Burgers vectors [12]. However, the creep deformation mechanism of specimens with large deviation on the boundaries is never reported. In this paper, creep property data on the [001]- [ 111] and [001]-[011] boundary at 760 C/760 MPa were reported. The deformation mechanism was investigated by observation of lattice rotation, dislocation configuration and calculation of Schimd Factors. Life / h Specimen A Specimen B Lives (h) Elongations (%) Specimen C Elongations / % 2. Experimental The composition of the single crystal superalloy studied is 5Cr, 10Co, 11W, 1Mo, 6Al, 1Ti, and balance Ni (wt%). Single crystal bars were directionally solidified using liquid metal cooling (LMC) process [13]. Single crystal seeds with pre-arranged orientation were cut and used in order to obtain the SC specimens with desired orientation. The as-cast samples were cut from the SC bars by using an electric discharge machine and machined into stress rupture specimens of 25 mm in gauge length and 5 mm in diameter. The stress rupture tests were performed at 760 C/760 MPa by using a constant load creep testing machine. Two specimens were tested for each orientation. The samples for electron back-scattered diffraction (EBSD) examination were cut from different locations of the ruptured specimens (near the screw and near the fracture), ground and electro-polished in a solution of 10% perchloric acid and 90% ethanol. Specimen cut at a distance of 5 mm away from fracture surface were mechanically ground to 40 µm and finally electro-polished at 20 C in a solution of 10% perchloric acid and 90% ethanol to prepare TEM foils. They were then examined under a TECNAI G2 20 microscope. 3. Results The relationship between the creep lives tested Fig. 2 Creep properties of specimens A C tested at 760 C/760 MPa at 760 C/760 MP and the orientations are shown in Fig. 1(a). The ranked regions are reported by Mackay [6]. It can be concluded that the tested creep lives agree well with the ranked regions except for the specimen B, which is on the [001]-[011] boundary and has the extraordinary creep live. Therefore, the deformation mechanism of specimen B is further researched. For comparison, specimens A and C are also analyzed. Their orientations have nearly the same deviation degrees from [001] orientation. The lattice rotations for specimens A, B and C are shown in Fig. 1(b). Their initial deviation from [001] is 35, 31 and 36, respectively. Specimen A is on the [001]-[ 111] boundary, while specimen B is on the [001]-[011] boundary. Specimen C is 13 away from the [001]-[011] boundaries. The lattice rotation is observed after creep rupture tests. The results show that specimen A rotates towards [ 111] along the [001]-[ 111] boundary, while specimen B rotates towards [001] along [001]-[011] boundary. Specimen C rotates towards [ 111] orientation and is ruptured before reaching the [001]-[ 111] boundary. The creep properties of specimens with these orientations are illustrated in Fig. 2, and there are two specimens for each orientation. It can be seen that specimen B has the longest rupture life and the short-

3 S.H. Zhang et al.: J. Mater. Sci. Technol., 2012, 28(3), Fig. 3 Fracture surfaces: (a) specimen A; (b) specimen B; (c) specimen C est elongation. However, when the orientation deviates from the [001]-[011] boundary, i.e. specimen C, the rupture life drops significantly. The fracture surfaces of specimens A and C shown in Fig. 3(a) and (c) are elliptic and inclined, while the fracture surface of specimen B (Fig. 3(b)) is relatively circular and perpendicular to the applied stress. The ratios of major and minor diameters of the specimens after test are measured as 1.26, 1.10 and 1.35, respectively for the specimens A, B, and C. It is interesting to note that the cross section of specimen B contracts along [010]. However, the contraction directions of specimens A and C have an angle with [010]. The dislocation configurations in specimens A C are compared in Fig. 4. In specimens A (Fig. 4(a)) and C (Fig. 4(d)), many stacking faults with single orientation in γ are observed. These stacking faults are commonly observed in SC alloys tested at low temperature/high stress conditions [14 17]. For specimen B as seen in Fig. 4(b) and (c), most of the dislocations are constrained in γ matrix, and a few stacking faults with two orientations are observed in γ. The angle between the two orientations is approximately 70, as indicated by the arrow in Fig. 4(c). 4. Discussion Generally, a single a/2<110> dislocation may dissociate at γ/γ interface via the following schemes during low temperature/high stress creep [18 21] : a/2[011] a/3[112] + a/6[ 21 1] (1) a/2[ 101] a/3[ 211] + a/6[1 21] (2) a/2[101] a/3[211] + a/6[ 1 21] (3) When the applied stress is sufficiently high, the a/3<112> dislocation is able to enter the γ, leaving a superlattice stacking fault (SSF) behind it. In order to analyze the detailed deformation mechanism, the Schmid Factors (SF) for all specimens are calculated and listed in Table 1. Table 1 Calculated Schimd Factors for specimens A C on the most potentially activated systems Specimen Slip system Schmid Factor (111) [ 101] A ( 1 11) [011] (111) [ 211] (111) [ 101] B ( 111) [101] (111) [ 211] ( 111) [211] C (111) [ 101] (111) [ 211] The Schmid Factors are calculated by: SF = cosψ cos λ (4) where ψ is the angle between the stress direction and slip plane, λ is the angle between the stress direction and slip direction. The maximum value of SF is 0.5. And the value of SF represents that whether a slip system is easily operated or not. The most highly stressed dislocations in specimen A are a/2[ 101] and a/2[011]. These two dislocations can dissociate to a/3[ 211] and a/3[112], respectively, according to reactions (1) and (2). However, the ( 1 11) [112] system has a relatively low SF (0.31), while SF of (111) [ 211] system is nearly the same as that of (111) [ 101] and ( 1 11) [011] systems (Table 1). Moreover, SF of (111) [ 211] system will be larger than those of (111) [ 101] and ( 1 11) [011] systems during rotating towards [ 111] orientation. Therefore, (111) [ 211] system may be activated in the matrix. When a/3[ 211] cuts into γ, SSF in γ with single orientation is observed in specimen A (Fig. 4(a)). The macroscopic observations, i.e. lattice rotation and contraction of cross section shown in Figs. 1 and 3(a) also confirm that the <112> type dislocations are activated, which is in good agreement with previous work [4,7,22]. In specimen B, (111) [ 101] and ( 111) [101] systems exhibit very high SF. Therefore, these two systems can easily be activated compared to the two (111)

4 232 S.H. Zhang et al.: J. Mater. Sci. Technol., 2012, 28(3), Fig. 4 Dislocation configurations in transverse cross-section of the specimens tested to failure: (a) specimen A; (b) and (c) specimen B; (d) specimen C <112> systems. As a result, the deformation of specimen B is controlled by a/2<110> dislocations slipping in matrix channels. However, (111) [ 211] and ( 111) [211] systems with relatively high SFs may also be generated by reactions (2) and (3) during creep. Therefore, a few SSFs showing 70 with each other are observed in this specimen (Fig. 4(c)). Moreover, the lattice rotation and the contraction direction of specimen B also indicate that the deformation is controlled by a/2<110> slip systems. In specimen C, dislocation a/2 [ 101] on the highly stressed (111) [ 101] system can dissociate to a/3[ 211] according to reaction (2). The (111) [ 211] slip system has a high SF of Thus, SSF with single orientation is found in specimen C. The lattice rotation and the contraction direction are similar to specimen A. The a/3<112> dislocation cutting into γ can dramatically increase primary creep strain and lower rupture life [18]. Consequently, the rupture lives of specimens A and C are relatively short. Since specimen C is ruptured before reaching the [001]-[ 111] boundary, the matrix deformation in specimen C is also controlled by single a/2<110> slip system, which leads to the most severe asymmetric deformation (Fig. 3(c)). Due to the single slip in both γ channels and γ phase, specimen C shows the shortest rupture life. The duplex a/2<110> in γ channels and stacking fault locks in γ induced by a/3<112> dislocations decreased the creep rate and therefore, specimen B shows the longest rupture life and lowest creep elongation. 5. Summary The creep rupture tests of specimens with orientation about 30 C away from [001] are carried out at 760 C/760 MPa. Specimens with orientation on [001]-[011] boundary show the longest rupture lives. However, the rupture lives greatly decrease when the orientation moves into the triangle of the stereographic projection or on the [001]-[ 111] boundary. The deformation mechanisms are investigated based on the TEM observation, lattice rotation and calculated Schmid Factors. The two a/2<110> and a/3<112> dislocations can be activated in the specimens with orientation on [001]-[011] boundary. Meanwhile, two a/2<110> and single a/3<112> dislocations operated in the specimens with orientation on [001]-[ 111] boundary. When the orientation moves into the stereographic triangle, deformation is controlled by single slip in both γ and γ. Acknowledgements The project is sponsored by the National Basic Research Program of China (Grant No. 2010CB631201) and the National Natural Science Foundation of China (Grant Nos , , ). REFERENCES [1 ] W. Betteridge and S.W.K. Shaw: Mater. Sci. Tech-

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