Materials Science and Engineering A

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Materials Science and Engineering A 528 (2011) 3062 3068 Contents lists available at ScienceDirect Materials Science and Engineering A journal homepage: www.elsevier.

Lou a a Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b Shenyang National Laboratory for Materials Science, Institute of Metal

1 Materials Science and Engineering A 528 (2011) Contents lists available at ScienceDirect Materials Science and Engineering A journal homepage: Intermediate temperature creep of directionally solidified Ni-based superalloy containing local recrystallization G. Xie a, L. Wang a, J. Zhang a,b,, L.H. Lou a a Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang , China b Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang , China article info abstract Article history: Received 12 August 2010 Received in revised form 15 November 2010 Accepted 29 December 2010 Available online 5 January 2011 Keywords: Recrystallization Superalloy Creep Crack Creep rupture life Intermediate temperature creep of directionally solidified Ni-based superalloys containing local recrystallization (RX) was studied. A sharp reduction of the creep rupture life was observed with increasing the transverse RX area fraction on the cross section of the specimen. Creep curves of specimens containing local RX showed that the steady state creep stage was much shorter than that of the good specimens. The creep rates were also slightly higher when there was local RX. The creep and failure mechanism was discussed based on the dislocation morphology in the vicinity of the RX grain boundaries and the formation and propagation of cracks in the RX region according to the microstructure observation of the interrupted creep specimens Elsevier B.V. All rights reserved. 1. Introduction The use of directionally solidified (DS) columnar-grained or single crystal superalloys shows significant advantages over conventional cast alloys due to the elimination of highly stressed transverse grain boundaries. Due to the extremely good elevated temperature capability, these materials are widely used as turbine airfoils in advanced aero-engines and industrial gas turbines [1,2]. However, residual stress or plastic deformation generated in DS blades during manufacturing may induce recrystallization (RX) that degrades the performance or even threatens the safe application of the blades. Factors that influence the nucleation and growth of RX, including applied stress [3,4], heat treatment parameters [4 7], heat treatment atmosphere [7], as well as the microstructural features such as residual eutectics and carbides [8 10] in several DS superalloys have been reported. It is believed that RX may result in the reduction of the mechanical properties of DS superalloys [11 18]. For example, only 5 50% of the creep rupture life can be achieved in DZ22 alloy if a surface RX layer was introduced by shot peening. The detrimental effect was Corresponding author at: Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang , China. Fax: address: jianzhang@imr.ac.cn (J. Zhang). more pronounced at intermediate temperature/high stress condition [12]. Similar data tested in directionally solidified MAR-M247 was also reported by Khan et al. [13]. Besides the RX layer induced by shot-peening, it was reported that local RX showed similar detrimental effect on the creep rupture life of single crystal superalloy SRR99 [16]. In our previous work, the detailed studies on microstructural evolution of DS superalloy containing local RX during creep test at high temperature/low stress have been conducted [19]. It showed that the whole steady state (secondary creep regime) was actually the stage of crack propagation and connection within the RX region. Crack did not propagate into the region without RX until the starting of the tertiary creep stage [19]. In present paper, the effect of local RX on the intermediate temperature creep properties of two DS superalloys, and the role of RX are explored by monitoring the microstructural evolution during creep. The failure mechanism of DS superalloy containing local RX is also discussed. 2. Experimental The nominal compositions of the DS superalloys studied (DZ125L and DZ17G) are listed in Table 1. Alloys were directionally solidified into plate with the size of 220 mm 70 mm 16 mm using a Bridgman furnace. Details of the DS process were reported elsewhere [20]. The resulted DS slab has an average grain size of 2 3 mm and a primary dendrite arm spacing around 300 m /$ see front matter 2011 Elsevier B.V. All rights reserved. doi: /j.msea

G. Xie et al. / Materials Science and Engineering A 528 (2011) 3062 3068 3063 Table 1 Nominal chemical composition in wt.% of alloys investigated. Cr Co W Mo Al Ti Ta V C B Ni DZ125L 9 10 7 2 5 3.

After heat treatment (1220 C/2 h/ac + 1080 C/4 h/ac + 900 C/16 h/ac for DZ125L and 1220 C/4 h/ac + 980 C/16 h/ac for DZ17G), the local RX occurred.

2 G. Xie et al. / Materials Science and Engineering A 528 (2011) Table 1 Nominal chemical composition in wt.% of alloys investigated. Cr Co W Mo Al Ti Ta V C B Ni DZ125L Bal. DZ17G Bal. DS slabs were cut into small plates by electron discharge machining (EDM) and indented using Brinell hardness tester on one side. After heat treatment (1220 C/2 h/ac C/4 h/ac C/16 h/ac for DZ125L and 1220 C/4 h/ac C/16 h/ac for DZ17G), the local RX occurred. The indentation was then carefully removed by grinding. (The specimens with different depths of local RX can be obtained by applying different indentation loads and controlling the subsequent grinding process. Detailed procedure of sample preparation was reported in our previous work [17].) The specimens were available in the form of plate with 50 mm 5mm 2 mm gauge section for creep tests, and 15 mm 5mm 2 mm gauge section for creep rupture tests. The effect of depth of local RX on intermediate temperature creep rupture properties was studied firstly. Creep rupture tests were performed at 760 C/725 MPa for DZ125L and at 760 C/675 MPa for DZ17G. Some of the creep tests performed at 760 C/725 MPa using DZ125L were interrupted at different stages to investigate the microstructural evolution during creep. The stress axis was parallel to the DS direction, and specimens without local RX were used as reference materials. Specimens for optical microscope and scanning electron microscope (SEM) observation were prepared by the standard metallographic procedure. Transmission electron microscope (TEM, Philips TECNAI 20) was employed to characterize the dislocation pattern during creep. After interrupted creep tests, samples were cut from the RX site as well as from the un-recrystallized region of the specimen. These samples were first mechanically polished into 50 m discs and then ion polished into TEM foils. 3. Results The morphology of RX under the indentation after heat treatment is shown in Fig. 1. The local RX mainly formed in the dendritic core region, leaving the interdendritic region un-recrystallized. It was evident from the SEM observation that very fine particles precipitated in the RX grains and coarse particles formed along the RX grain boundary after heat treatment (Fig. 1(b)). This type of RX morphology resulted from the pinning effect of un-dissolved large, / eutectics, and/or carbides at interdendritic region during heat treatment has been observed in many superalloys [6,8,11]. The effect of the maximum depth of the local RX on the creep rupture properties at 760 C/725 MPa for DZ125L and at 760 C/675 MPa for DZ17G is summarized in Fig. 2(a) and (b). For DZ125L, the creep rupture life decreased rapidly with the increase of RX depth and the specimens almost lost bearing capacity completely when the local RX depth was beyond 800 m. Elongation to failure gradually decreased with the increase of RX depth. A sharp reduction of creep rupture life was observed by introducing about 400 m RX for DZ17G. Like DZ125L, elongation to failure also decreased with increasing the RX depth. Fig. 1. Morphology of local RX in DZ125L after heat treatment (1500 kg indentation load). (a) Low magnification showing the isolated RX grains distributed primarily at the dendrite core and (b) SEM image showing the coarse formed at the RX grain boundary. Fig. 2. Change of creep rupture properties with increase of RX depth. (a) DZ125L alloy at 760 C/725 MPa. (b) DZ17G alloy at 760 C/675 MPa.

3 3064 G. Xie et al. / Materials Science and Engineering A 528 (2011) Fig. 3. Normalized creep rupture life as a function of TRF in several DS superalloys. Normalized creep rupture life equals to the creep life of a specimen divided by the average creep life of the specimens without any local RX. TRF is defined as A RX/A, where A RX is the area of the RX on the cross section, and A is the area of the cross section of the specimen (see the insert) [14]. In order to quantitatively evaluate the effect of RX on creep rupture life, data obtained from present experiments as well as from literature was examined. The normalized creep rupture life of different DS superalloys was plotted as a function of transverse RX area fraction (TRF) as shown in Fig. 3. It was apparent that the creep rupture life of the DS superalloys was reduced by the local RX. A linear reduction of the creep rupture life was found when TRF was less than 10%, and then the specimen containing RX almost lost the bearing capacity. Creep curves of two specimens of DZ125L (specimen A, good specimen without RX and specimen B containing local RX within the gauge section) are compared in Fig. 4(a). The maximum depth of local RX in specimen B is about 600 m. Three regimes were observed, primary creep regime, followed by secondary (steady) creep and a very short tertiary (accelerated) creep regime. The secondary creep regime was reduced greatly when there was local RX in the specimen. The creep rate of different specimens as a function of time is shown in Fig. 4(b). The creep rate of specimen B at the beginning of the primary creep and at the secondary creep stage was higher than that of good specimen A. The microstructure interrupted at primary creep stage (point 1 in Fig. 4(b)) is shown in Fig. 5. It can be seen that a large number of the transverse RX grain boundaries cracked. The cracks connected with each other when there is little or no unrecrystallized DS material between them (Fig. 5(b)). Alternatively, the cracks deviated along the grain boundary (Fig. 5(c)). From this figure, one can also see that cracks initiated dispersedly at some sites of the same RX grain boundary (showed as white arrows) and then linked with each other. However, the region without RX remains intact at this stage. TEM observation demonstrated that deformation was inhomogenous at this stage. Within the RX grain where the fine distributed uniformly, very low dislocation density was observed. Dense dislocations in the matrix were only found occasionally (Fig. 6(a)). Slip bands were observed in the region without RX, as shown in Fig. 6(b). At RX grain boundary, dislocations mainly generated in the channel of the region without RX (Fig. 6(c)). Fig. 7 shows the optical micrograph of RX specimen interrupted at the end of the primary creep stage (point 2 in Fig. 4(b)). Comparing to Fig. 5, the width of cracks increased obviously. The primary crack had penetrated through the whole RX region and began to propagate into the region without RX along the interdendritic region (Fig. 7(b)). During the steady creep stage (point 3 in Fig. 4(b)), the primary crack continued to propagate in the region without RX (Fig. 8(a)). The length of crack cut into the region without RX increased slowly. In the periphery of RX region, a large amount of slip lines can be observed (Fig. 8(b)). At this stage, cracks were also observed at the transverse segment of the grain boundaries in the region without RX (Fig. 9). 4. Discussion Microstructural evolution during 760 C creep of the DS alloy is schematically shown in Fig. 10. At primary creep stage, formation and connection of micro-voids at RX grain boundaries due to the piled-up of dislocations resulted in the initiation of cracks (Figs. 10(a) and (b)). The cracks tended to connect with each other when there were recrystallized grains nearby, or otherwise deviated along the grain boundary (Figs. 5 and 10(c)). Consequently, a larger creep rate was found comparing to the good specimen. The whole RX region had been penetrated before the steady state creep stage (Fig. 7(b)). Stress concentration generated from the crack led to severe deformation and a number of slip lines produced at the tip of cracks near the RX region during the steady state creep stage (Fig. 8(b)). Cracks were also observed in the region without RX (Figs. 9 and 10(d)). The enhanced true stress and continuous propagation of cracks resulted in higher creep rate and shorter steady state creep regime in recrystallized specimen. Fig. 4. Comparison of (a) creep and (b) creep rate curves of DZ125L alloy without (A) and with (B) local RX as a function of time. Arrows in the figure (b) indicate the interrupted creep tests of curve B. Creep tests performed at 760 C/725 MPa.

G. Xie et al. / Materials Science and Engineering A 528 (2011) 3062 3068 3065 Fig. 5. (a) RX grain boundaries and cracks after interrupted creep test (point 1 in Fig. 4(b)).

4 G. Xie et al. / Materials Science and Engineering A 528 (2011) Fig. 5. (a) RX grain boundaries and cracks after interrupted creep test (point 1 in Fig. 4(b)). (b) and (c) are the different modes of crack propagation with different distributions of RX. Table 2 Comparison of failure mode of RX specimen at different creep conditions. Primary creep Secondary creep Tertiary creep High temperature/low stress creep Initiation of cracks Crack propagates within the RX region Intermediate temperature/high stress creep Crack initiates in the RX region Crack propagates in the DS and propagates into DS material material Crack propagates into region without RX and failure Linkage of cracks and failure Comparing to the creep behavior at high temperature/low stress, the deformation and failure mechanisms are obviously different at intermediate temperature/high stress (Table 2): (1) In the RX region, cracks had barely initiated (Fig. 10(a)) during the primary creep stage at high temperature [19]. However, cracks have already penetrated the RX region and propagated into regions without RX at intermediate temperature. (2) At high temperature, cracks propagated within the RX region during the secondary creep stage. As soon as the primary crack propagated into the region without RX, tertiary creep began [19]. On the other hand, cracks propagated in the region without RX during the whole secondary creep stage at intermediate temperature. Due to the large applied stress, pile-up or tangle of dislocations which is the main mode of crack initiation at intermediate temperature occurred at RX grain boundary quickly. On the other hand, initiation of cracks is controlled by two mechanisms at high temperature, i.e., the pile-up of dislocations as well as the accumulation of vacancies (Coble creep) [19,21,24]. The applied stress was relatively small at high temperature and there was strong recovery which may release part of stress concentration originated from dislocation pile-up [22,23]. Coble creep was an extremely slow process according to our previous work [24]. Therefore, the initiation of cracks was significantly delayed at high temperature. The results of creep rupture tests shown in Fig. 3 can therefore be interpreted by the above creep mechanism analysis. The bearing capacity of the recrystallized region at intermediate temperature/high stress was extremely low. Crack initiated and propagated rapidly in the RX region during the primary stage of creep, which resulted in the rapid reduction of the creep rupture life with the increase of the TRF.

3066 G. Xie et al. / Materials Science and Engineering A 528 (2011) 3062 3068 Fig. 6.

specimen interrupted at the beginning of primary creep stage (point 1 in Fig. 4(b)). Fig. 7.

5 3066 G. Xie et al. / Materials Science and Engineering A 528 (2011) Fig. 6. Characteristic of dislocation in (a) RX grain, (b) region without RX and (c) near the RX grain boundary of the recrystallized specimen interrupted at the beginning of primary creep stage (point 1 in Fig. 4(b)). Fig. 7. Microstructure of the RX specimen interrupted at the end of primary creep stage (point 2 in Fig. 4(b)). (a) Crack penetrated the whole recrystallized region. (b) Crack propagated into the region without RX.

G. Xie et al. / Materials Science and Engineering A 528 (2011) 3062 3068 3067 Fig. 8. Optical micrograph of the specimen interrupted at the steady state (point 3 in Fig. 4(b)).

Microstructure in the region without RX of the specimen interrupted at the steady state (point 3 in Fig. 4(b)). Cracks initiated at the transverse segment of the original DS grain boundary.

A rapid reduction of the creep rupture life was observed when TRF increased from 0 to 10%, and the bearing capacity almost lost when TRF was more than 10%.

6 G. Xie et al. / Materials Science and Engineering A 528 (2011) Fig. 8. Optical micrograph of the specimen interrupted at the steady state (point 3 in Fig. 4(b)). (a) Cracks in the RX region. (b) Cracks continued to propagate in the region without RX. 5. Conclusions Fig. 9. Microstructure in the region without RX of the specimen interrupted at the steady state (point 3 in Fig. 4(b)). Cracks initiated at the transverse segment of the original DS grain boundary. The present work indicates that 760 C creep property is extremely sensitive to the local RX induced by surface deformation. A rapid reduction of the creep rupture life was observed when TRF increased from 0 to 10%, and the bearing capacity almost lost when TRF was more than 10%. The primary and secondary creep regimes of RX specimens were shorter than those of good specimens. Due to the initiation and propagation of cracks at the RX grain boundaries, higher creep rate of the RX specimen at the beginning of the primary creep stage was observed. Cracks between RX grains almost linked with each other completely at the beginning of the primary creep stage and began to propagate into the region without RX at the end of this stage. The enhanced true stress and continued propagation of cracks resulted in the shorter steady creep regime and a slight increase of creep rate. Connection between primary crack and those cracks formed in the region without RX led to the final failure of the specimen at the tertiary creep regime. Fig. 10. Schematic of creep failure mechanism of DS superalloy containing local RX. (a) Initiation of cracks. (b) Crack propagation along RX grain boundaries. (c) Crack propagation in the RX region. (d) Crack propagates into region without RX outside the RX region and crack initiates in the region without RX.

7 3068 G. Xie et al. / Materials Science and Engineering A 528 (2011) Acknowledgements This work was financially supported by the National Basic Research Program (973 Program) of China under grant No. 2010CB and National Natural Science Foundation of China under grant No The authors are grateful for these supports. References [1] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, John Wiley & Sons, Inc., New York, 1987, pp [2] R.C. Reed, The Superalloys Fundamentals and Applications, Cambridge University Press, Beijing, 2006, pp [3] C. Zambaldi, F. Roters, D. Raabe, U. Glatzel, Mater. Sci. Eng. A (2007) [4] D.C. Cox, B. Roebuck, C.M.F. Rae, R.C. Reed, Mater. Sci. Technol. 19 (2003) [5] R. Bürgel, P.D. Portella, J. Preuhs, in: T.M. Pollock, R.D. Kissinger, R.R. Bowman, K.A. Green, M. McLean, S. Olson, J.J. Schirra (Eds.), Superalloys 2000, TMS, New York, 2000, pp [6] S.D. Bond, J.W. Martin, J. Mater. Sci. 19 (1984) [7] G. Xie, J. Zhang, L.H. Lou, Scripta Mater. 59 (2008) [8] L. Wang, G. Xie, J. Zhang, L.H. Lou, Scripta Mater. 55 (2006) [9] J.C. Xiong, J.R. Li, S.Z. Liu, J.Q. Zhao, M. Han, Mater. Charact. 61 (2010) [10] L. Wang, F. Pyczak, J. Zhang, R.F. Singer, Int. J. Mater. Res. 100 (2009) [11] C.Y. Jo, H.M. Kim, Mater. Sci. Technol. 19 (2003) [12] Y.R. Zheng, Z.C. Ruan, S.C. Wang, Acta Metall. Sin. 31 (1995) [13] T. Khan, P. Caron, Y.G. Nakagawa, J. Met. 38 (1986) [14] M.Z. Alam, D. Chatterjeea, B. Venkataramana, V.K. Varmab, D.K. Das, Mater. Sci. Eng. A 527 (2010) [15] C.Q. Sun, C.H. Tao, N.S. Xi, W.F. Zhang, C.X. Wu, Mater. Mech. Eng. 25 (2001) 4 7. [16] D.L. Wang, Ph.D. thesis, Institute of Metal Research, Shenyang, [17] G. Xie, L. Wang, J. Zhang, L.H. Lou, Metall. Mater. Trans. A 39A (2008) [18] J. Meng, T. Jin, X.F. Sun, Z.Q. Hu, Mater. Sci. Eng. A 527 (2010) [19] G. Xie, L. Wang, J. Zhang, L.H. Lou, in: R.C. Reed, K.A. Green, P. Caron, T.P. Gabb, M.G. Fahrmann, E.S. Huron, S.A. Woodard (Eds.), Superalloys 2008, TMS, New York, 2008, pp [20] B.C. Yan, J. Zhang, L.H. Lou, Mater. Sci. Eng. A 474 (2008) [21] M.E. Kassnera, T.A. Hayes, Int. J. Plast. 19 (2003) [22] M.H. Yoo, H. Trinkaus, Acta Metall. 34 (1986) [23] A. Epishin, T. Link, in: K.A. Green, T.M. Pollock, H. Harada, T.E. Howson, R.C. Reed, J.J. Schirra, S. Walston (Eds.), Superalloys 2004, TMS, New York, 2004, pp [24] L. Wang, G. Xie, J. Zhang, L.H. Lou, Mater. Sci. Forum (2007)

1 Introduction. Zhen-Xue Shi 1 Shi-Zhong Liu

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