RARE METALS Vol. 30, Spec. Issue, Mar 2011, p. 462 DOI: 10.1007/s12598-011-0325-2

Corresponding author: TIAN Sugui E-mail: tiansugui2003@163.com

Microstructure and creep behavior of A 9%W single crystal nickel-based superalloy LI Anan, TIAN Sugui, LIANG Fushun, LI Jingjing, WANG Xiaoliang, and ZHANG Te School of Materials Science and Engineering, Shenyang University of Technology, Shenyang 110870, China

Received 29 September 2010; received in revised form 29 December 2010; accepted 10 January 2011 © The Nonferrous Metals Society of China and Springer-Verlag Berlin Heidelberg 2011

Abstract

By means of the measurement of creep curves and microstructure observation, an investigation has been made into the microstructure evolu-tion and creep behaviors of 9%W single crystal nickel-base superalloy. Results show that the alloy displays an obvious sensibility on the ap-plied stress when applied stress is more than 160 MPa at 1040 ºC. In the ranges of the applied temperatures and stresses, the apparent creep activation energy is measured to be about 465 kJ/mol. In the initial stage of creep, the cubical γ′ phase in the alloy is transformed into the N-type rafted structure along the direction vertical to the applied stress axis, the deformed mechanism of the alloy during steady state creep is dislocations climbing over the rafted γ′ phase, the dislocation shearing into the rafted γ′ phase is thought to be the creep mechanism of the al-loy during later stage of creep. After crept up to fracture, the various morphology of the rafted γ′ phase is displayed in the different regions of the sample, the rafted γ′ phase vertical to the stress axis displays in the region far from the fracture, but the coarser twisted γ′ phase is detected in the regions near the fracture, which is attributed to the bigger plastic deformation occurred in the region near the fracture.

Keywords: single crystal nickel-base superalloy; element W; microstructure evolution; creep behavior; deformation features

1. Introduction

Single crystal nickel-base superalloys have been widely used, because they possess high volume fraction of γ′ strengthening phase and good high temperature properties, for preparing the blade parts of the advanced engine and combustion turbine [1-3]. The microstructure of the alloy consists of the cubical γ′ phase embedded coherent in the γ matrix phase, and the size, morphology and distribution of the cubical γ′ phase are related to the heat treatment process [4-5]. Adding the refractory elements Re and W can im-prove the high temperature properties and the creep resis-tance of superalloys [6-8]. Specially the refractory element W has the bigger atomic radius and the smaller partitioning ratio in γ′/γ phases, and has the higher solubility in the nickel-base alloy [9-10], so the effect of the solution strengthening of the alloy increase with the element of W content. Some literatures reported the creep behavior of sin-gle crystal nickel based superalloys [11-12], however, a few literature reports on the creep behaviors of the W-richer sin-gle crystal nickel based superalloy at high temperature.

In the paper, the creep behaviors and the deformation fea-tures of the W-richer single crystal nickel based superalloy are investigated by microstructure observation and creep

curves measuring, and deformation mechanism of the alloy during creep is briefly discussed.

2. Experimental

The single crystal nickel-base superalloy with [001] ori-entation has been produced by means of selecting crystal method in a vacuum directional solidification furnace under the condition of a high temperature gradient. The nominal chemical composition of the superalloy is Ni-Cr-Co-Al- Ta-Mo-9%W (wt,%). The heat treatment regimes of the sin-gle crystal nickel based superalloy bars are given as follows: 1280 ºC × 2 h + 1325 ºC × 4 h, A.C + 1080 ºC × 4 h, A.C + 870 ºC × 24 h, A.C.

After the crystal orientation was determined by Laue-back reflection, the heat treated single crystal bars were machined into the tensile creep samples along the [001] orientation, with a cross-section of 4.5 mm × 2.5 mm and the gauge length of 17 mm. The wider surface of the specimen was parallel to (100) plane. The uniaxial constant load tensile tests were conducted in the creep testing machine with GWT504 model. The creep curves of the alloy at the differ-ent conditions were measured. The microstructure of the al-

Li A.A. et al., Microstructure and creep behavior of A 9%W single crystal nickel-based superalloy 463

loy at the different states is observed by using SEM and TEM. The activation energy and stress exponent of the alloy during steady state creep are calculated according to the creep data.

3. Results and analysis

3.1. Creep features of the alloy

The creep curves of the single crystal nickel based super-alloy contain 9%W in different conditions are measured as shown in Fig. 1, indicating that the alloy display the obvious features of creep three stages, including the initial creep, steady state creep and accelerated creep stages.

The creep curves of the alloy under the applied stress of 137 MPa at different temperatures are shown in Fig. 1 (a), this indicates that the bigger initial strain of the alloy occurs when applying load at high temperature. As the creep goes on, the dislocation density in alloy increases to decrease the strain rate up to entering the steady state. The strain rate of the alloy during steady state creep is measured to be about 0.0144%/h at 1040 ºC, and the creep lifetime is measured to be 422 h. As the creep temperature enhances to 1060 ºC, the strain rate of the alloy during steady state creep increases to 0.0272%/h, the creep lifetime is measured to be 285 h. In the further, the creep lifetime of the alloy decreases to 138 h as the creep temperature enhances to 1072 ºC.

The creep curves of the alloy under the applied different stresses at 1040 ºC are shown in Fig. 1 (b), this indicates that the strain rate of the alloy during the steady state creep en-hances with the applied stress. Under the applied stress of 160 and 180 MPa, the strain rates of the alloy are measured, respectively, to be 0.0315%/h and 0.0529%/h, the creep life-times of the alloy are measured, respectively, to be 200 and 98 h. This indicates that the alloy displays an obvious sensi-bility on the applied stress when the applied stress is more than 160 MPa.

3.2. Constitutive equation and relative parameters

The initial strain of the single crystal nickel based super-alloy occurs when applied the tensile stress. As the creep goes on, the strain rate of the alloy decreases up to entering the steady state stage, in which the strain rate of the alloy maintains constant, and the strain rate of the alloy during steady state creep may be expressed by Dorn creep law given as follows [11]:

)exp(RTQ

A nAss −= σε

(1)

where, ssε is the strain rate during the steady state creep, A being the constant related to material structure, Aσ being applied stress, n being the apparent stress exponent, R being the gas constant, T being thermodynamics temperature, Q being the apparent active energy.

According to the data in Fig. 1, the dependences of the strain rates of the alloy during steady state creep on the ap-plied temperatures, stresses are shown in Fig. 2. Thereinto, the relationship between the strain rate and the applied tem-peratures is shown in Fig. 2 (a), the relationship between the strain rate and the applied stresses is shown in Fig. 2 (b). Therefore, the apparent creep activation energy and stress exponent of the alloy are measured to be Q = 465 kJ/mol and n = 4.78, respectively. It can be deduced according to the data that the strain rate of the alloy during steady state creep is controlled by the climbing of dislocations in the ranges of the applied temperatures and stresses.

3.3. Influence of heat treatment on microstructure

The microstructures of 9%W superalloy at different stages of heat treatment, which the samples are eroded by electrolysis method, are shown in Fig. 3. When the alloy is solution treated at 1325 ºC for 4 h, the γ′ phase and eutectic structure in the alloy is completely dissolved, after cooled in air, the fine γ′ phase about 200 nm in size is dispersedly pre-cipitated in the γ matrix of the alloy, as shown in Fig. 3 (a).

Fig. 1. Creep curves of 9%W single crystal nickel-base superalloy at different conditions: (a) various temperatures under applied stress of 137 MPa; (b) applied different stresses at 1040 ºC.

464 RARE METALS, Vol. 30, Spec. Issue, Mar 2011

Fig. 2. Relationship between the strain rate and the applied temperatures, stresses during the steady state creep: (a) strain rate & temperature; (b) strain rate & applied stress.

Fig. 3. Morphology after the alloy heat treated at different conditions: (a) solution at 1325 ºC; (b) first aged; (c) full heat treatment.

After aging treated for 4 h at 1080 ºC, the fine γ′ phase grows into the cubical configuration about 0.4 - 0.5 μm in size, but the smooth corners is displayed in the cubical γ′ phase, as shown in Fig. 3 (b). After aging treated for 24 h at 870 ºC, no obvious change in size is detected in the cubical γ′ phase, but the cubical extent of the γ′ phase increases, as shown in Fig. 3 (c). This indicates that no TCP phase is de-tected in the alloy, and the composition segregation in the dendrite / interdendrite of the alloy may be reduced [13] to adjust the size and distribution of the cubical γ′ phase during heat treated at high temperature.

3.4. Microstructure evolution of the alloy during creep

After the alloy is crept for 422 h up to fracture under the applied stress of 137 MPa at 1040 ºC, the morphologies in the different regions of the sample are shown in Fig. 4. The various microstructures are displayed in the different regions due to the ones supporting the various stress. Therefore, the deformation extent of the alloy in the different regions may be analyzed according to the configuration of γ′ phase.

The schematic diagram of observing regions in the sam-ple is shown in Fig. 4 (a), no stress is applied in the region A in which the coarsening of γ′ phase occurs only to form the meshlike structure along the vertical or horizontal directions, as shown in Fig. 4 (b). But the γ′ phase in the region B has

been transformed into the rafted structure along the direction vertical to the applied stress axis, the thickness of the rafted γ′ phase in the region B is about 0.6 μm, as shown in Fig. 4 (c). The morphology of the rafted γ′ phase in the region C is similar to the one in the region B, but the size of the rafted γ′ phase in thickness increases slightly as shown in Fig. 4 (d). In the further, the twisted configuration of the rafted γ′ phase appear in the region D due to the bigger plastic strain, and the size of the rafted γ′ phase in thickness increases to 0.8 - 0.9 μm, as shown in Fig. 4 (e). The much more coarsening and twisting of the rafted γ′ phase occurs in the region E near the fracture, so that the size of the γ′ phase in thickness in-creases to about 1 μm and the orientation of the rafted γ′ phase have about 45° angle relative to the direction of the applied stress axis as shown in Fig. 4 (f). It may be considered by analysis that the coarsening and twisted of the rafted γ′ phase in the region E are attributed to the severe plastic strain.

Under the applied stress of 137 MPa at 1040 ºC, the mi-crostructure of the alloy crept for different time are shown in Fig. 5. After crept for 2 h, no fully rafted structure of γ′

Fig. 4. Morphology of γ′ rafts in the different regions after the alloy crept up to fracture: (a) Schematic diagram of marking regions in specimen, (b), (c), (d), (e) and (f) being SEM mor-phologies corresponding to A, B, C, D and E regions of the specimen, respectively.

Li A.A. et al., Microstructure and creep behavior of A 9%W single crystal nickel-based superalloy 465

Fig. 5. Morphology of alloy crept for different time under ap-plied stress of 137 MPa at 1040 ºC: (a) 2 h; (b) crept for 30 h; (c) crept for 422 h up to fracture.

phase is formed in the alloy, the only several γ′ phase is linked to form the rafted structure along the direction verti-cal to the stress axis as marked by letter A in Fig. 5 (a). After crept for 30 h, the creep of the alloy enters the steady state stage, the γ′ phase has been transformed into the rafted structure along the direction vertical to the stress axis; the thickness of the rafted γ′ phase is about 0.5 μm; and the dis-location networks appear in the interfaces between γ′ and γ phases, as marked by arrow in Fig. 5 (b).

After crept for 422 h up to fracture, the dislocation net-works exists still in the interfaces of γ′ and γ phases, and significant amount of dislocations shear into the rafted γ′ phases as shown in Fig. 5 (c), And the trace of the disloca-tion displays the feature with 90° folded line as marked by arrow in Fig. 5 (c), indicating that the cross-slipping of dis-location occurs during creep of the alloy. Thereinto, the di-rection of the straight-like dislocations is at 45° agree rela-tive to the orientation of the rafted γ′ phase, this indicates that the trace of the dislocations shearing into the rafted γ′ phase is the direction along the biggest shearing stress, and the rafted γ′ phase in the alloy has lost the creep resistance in the later stage of creep.

4. Discussion

In the initial stage of creep under the applied stress of 137 MPa at 1040 ºC, the deformation feature of the alloy is that dislocations move in the matrix channel, no dislocations shear into the γ′ phase. When the creep enters the steady stage, the deformation mechanism of alloy is that the dislo-cations climb over the rafted γ′ phase and the dislocation networks appear in the interface of γ′ and γ phases as shown in Fig. 5 (b). The dislocations networks may relax the lattice mismatch stress in the interface of γ′/γ phases [14-15] for decreasing the interfacial energy to stabilize the microstruc-ture. Meanwhile, the deformed dislocations in the matrix during creep move to the interfaces to change the original direction for promoting the dislocation climbing over the rafted γ′ phase, therefore, the interface dislocation plays the

coordinating role on the deformed strengthening and recov-ery softening during creep of the alloy. It can be concluded that the existence of the interfacial dislocation networks may delay the dislocations shearing into the rafted γ′ phase for improving the creep resistance of alloy.

In the later period of the creep, some of the dislocation networks in the interfaces have been damaged because sig-nificant amount of dislocations are piled up in the interfaces for causing the stress concentration as marked by letter A in Fig. 5 (c), which may result in the dislocation shearing into the rafted γ′ phase for enhancing the strain rate up to the oc-currence of the creep fracture.

It is indicated according to the morphology of in different region of the alloy crept for 422 h up to fracture at 137 MPa/ 1040 ºC that the orientation of the straight-like rafted γ′ phase in the region far from the fracture is vertical to the stress axis, but the coarsening and twisting extent of the rafted γ′ phase in the region near fracture increase, as shown in Fig. 4 (f), which indicates that the coarsening and twisting of the rafted γ′ phase increase with the strain of the alloy due to the reducing the cross-section of the sample and increas-ing the effective stress. As the creep goes on in the later stage of creep, significant amount of dislocations are alter-nately activated along the direction with the maximum shearing stress for twisting the rafted γ′ phase, which results in the micro-cracks formed in the interfaces of the rafted γ′/γ phases, and propagated up to the occurrence of creep frac-ture. This is thought to be the main reason of occuring the creep fracture of the alloy.

5. Conclusion

(1) Under the applied stress of 137 MPa at 1040 ºC, the 9%W single crystal nickel based superalloy displays a lower strain rate and longer creep lifetime, and displays an obvious sensibility on the applied stress when the one is more than 160 MPa at 1040 ºC. In the range of the applied stresses and temperatures, the apparent creep activation energy of the al-loy is measured to be Q = 465 kJ/mol.

(2) In the initial period of creep at 137 MPa / 1040 ºC, the cubical γ′ phase in the alloy is transformed into the N-type rafted structure along the direction vertical to the applied stress axis. As the creep enters the steady stage, the defor-mation mechanism of the alloy is that the dislocations climb over the γ′ phase, and then the deformation feature of the al-loy in the later stage of creep is that dislocations shear into the rafted γ′ phase.

(3) After the alloy is crept up to fracture, the various mor-phologies of the rafted γ′ phase displays in the different re-gions of the sample. The rafted γ′ phase formed along the direction vertical to the stress axis appears in the region far

466 RARE METALS, Vol. 30, Spec. Issue, Mar 2011

from the fracture, and coarsening and twisting configuration of the rafted γ′ phase are displayed in the region near the fracture, which is attributed to the severe plastic deformation of the region.

Acknowledgement

The work was supported by the National Natural Science Foundation of China (No. 50571070).

References

[1] Hu Z.Q., Liu L.R., Jin T., and Sun X.F., Development of the Ni-base single crystal superalloys, Aeroengine, 2005, 31 (3): 1.

[2] Luo Z.P., Wu Z.T., and Miller D.J., The dislocation micro-structure of A nickel-base single-crystal superalloy after ten-sile fracture, Mater Sci Eng, A., 2003, A354: 358.

[3] Liang X.H., Zhou K.S., Liu M., and Deng C.G., Recrystalliza-tion on interface between NiCoCrAlYTa coating and nickel-based super-alloy, Rare Metal Materials and Engi-neering, 2009, 38 (3): 545.

[4] Nystrom J.D., Pollck T.M., Murphy W.H., and Garg A, Dis-continuous cellular precipitation in a high-refractory nickel- base supcralloy, Metall Mater Trans, A, 1997, 28: 2443.

[5] Yu X.H., and Yro Y.Y., Design of quaternary Ir-Nb-Ni-Al re-fractory supcralloys, Metall. Trans, A, 2000, (31): 173.

[6] Kobayash T., Harada H., and Zhan J.X., Influence of heat treatment on microstructure and mechanical properties of a 1st generation single-crystal superalloy, Journal of the Japan Institute of Metals, 2006, 7: 47.

[7] Wang K.G., Li J.R., and Cao C.X., Present situation of study on creep behavior of single crystal superalloys, Journal of Mater. Eng., 2004, 1: 3.

[8] Walston S., Cetel A., and MacKay R., Joint development of a fourth generation single crystal superalloy, Superalloys, 2004, 15.

[9] Nathal M.V., and Ebert L.J., The influence of cobalt, tantalum, and tungsten on the microstructure of single crystal nickel- base superalloys, Metall Trans., 1985, (16A): 1849.

[10] Yang Dayun, Zhang Xuan, Jin Tao, Sun Xiaofeng, Guan Hengrong, and Hu Zhuangqi, The influence of cobalt, tung-sten, and titanium on stress-rupture properties of nickel- base single crystal superalloy, Rare Metal Mater. Eng., 2005, 34 (8): 1295.

[11] Mukherjee A.K., Bird J.E, and Dorn J.E., Experimental cor-relation for high-temperature creep, Trans ASM, 1969, 62: 155.

[12] Kearsey R.M., Beddoes J.C., Jaansalu K.M., et al, The effects of Re, W and Ru on microsegregation behaviour in single crystal superalloy//SUPERALLOYS 2004-Proceedings of the Tenth International Symposium on Superalloys. United States: Metals and Materials Society, 2004. 801.

[13] Zhang J.H., Hu Z.Q., Xu Y.B., and Wang Z G. Dislocation structure in a single crystal nickel-base superalloy during low cycle fatigue, Metall. Trans. (A), 1992, 23: 1253.

[14] Gabrisch H., Mukherji D., and Wahi R.P., Deformation in-duced dislocation networks at the γ/γ′ interfaces in the single crystal superalloy, Philosophical Magazine, 1996, A74 (1): 229.

[15] Gabb T.P., Draper S.L., and Hull D.R., The role of interfacial dislocation networks in high temperature creep of superalloys, Material Science and Engineering , 1989, A118 (3): 59.