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Microstructural evolution and compressive deformation of a new Ni–Fe base superalloy after long term thermal exposure at 700 °C
Materials Science & Engineering A 619 (2014) 364–369
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructural evolution and compressive deformation of a new Ni–Fe base superalloy after long term thermal exposure at 700 1C Y. Yuan a,n, Z.H. Zhong b,nn, Z.S. Yu c,a, H.F. Yin a, Y.Y. Dang a, X.B. Zhao a, Z. Yang a, J.T. Lu a, J.B. Yan a, Y. Gu a a
Thermal Power Research Institute Co. Ltd., No. 136, Xingqing Road, Xi'an 710032, China School of Materials Science and Engineering, Hefei University of Technology, No.193, Tunxi Road, Hefei 230009, China c The Key State Laboratory for Mechanical Behavior of Materials, Xi'an Jiaotong University, Xi'an 710049,China b
art ic l e i nf o
a b s t r a c t
Article history: Received 29 July 2014 Received in revised form 24 September 2014 Accepted 26 September 2014 Available online 5 October 2014
A new Ni–Fe base superalloy for applications of advanced ultra-supercritical (A-USC) boiler tubes was heat treated to three different statuses, i.e. aging treatment, aging treatment plus 1000 h and 7000 h thermal exposure at 700 1C. The specimens with various heat treatments have been compressed at 700 1C. The yield strength first increased with thermal exposure time and decreased after further thermal exposure. After aging treatment and 1000 h thermal exposure, deformation was heterogeneous, and dislocation slip combing climb was dominant process. However, after 7000 h thermal exposure, deformation became homogeneous. Stacking faults shearing the γ0 precipitates were also observed. The coarsening of γ0 precipitates and α-Cr precipitation occurred during long term thermal exposure. The influence of microstructural evolution on the yield strength and deformation mechanisms is discussed. & 2014 Elsevier B.V. All rights reserved.
Keywords: Ni–Fe base superalloy Compressive deformation Thermal exposure
1. Introduction Advanced ultra-supercritical (A-USC) technology is an effective method to improve thermal efficiency and reduce CO2, NOx and SOx emissions for coal fired power plants [1]. Currently, the main steam parameters of A-USC power plants have been designed to be 700 1C and 35 MPa in Europe [2], or even higher in USA, i.e. 760 1C and 35 MPa [3]. For 700 1C-class A-USC power plants, the thermal efficiency is projected to be above 50%. The requirements of high temperature materials for use in A-USC boilers are 100,000 h creep strength greater than 100 MPa at 750 1C and 200,000 h coal-ash corrosion resistance of less than 2 mm metal loss [4]. The ferritic and austenitic heat resistant steels cannot meet these requirements. Ni-based superalloys are thought to be the promising candidates for high temperature components in A-USC power plants. For example, Inconel 740/740 H [5–7], CCA617 [8–10], Haynes 282 [11] and Nimonic 263 [12] have been widely evaluated. According to the reported results, 740 H has the highest rupture strength at 700 1C with rupture life of 105 h [13]. In China, the research and development of 700 1C A-USC technology is also ongoing. For high temperature materials, a
n
Corresponding author. Tel.: þ 86 29 8210 2731; fax: þ 86 29 8210 2090. Corresponding author. E-mail addresses: [email protected] (Y. Yuan), [email protected] (Z.H. Zhong). nn
http://dx.doi.org/10.1016/j.msea.2014.09.095 0921-5093/& 2014 Elsevier B.V. All rights reserved.
new Ni–Fe base superalloy, has been developed at Thermal Power Research Institute, Xi'an, China. This new alloy is a candidate for applications of 700 1C A-USC boiler tubes. It is a γ0 strengthened superalloy, and the volume faction of γ0 after aging treatment is around 17%. Its temperature capability with rupture life of 105 h at 100 MPa is roughly 754 1C, comparable with 740 H (765 1C) [14]. This new alloy contains more than 20 wt% Fe, and no Co addition. Compared to 740 H, it has much lower cost and better hot workability. For this new alloy, it is important to deeply understand the microstructures and related deformation mechanisms after long term thermal exposure. In the present study, the compressive testing of the new alloy after aging treatment and long term thermal exposure was conducted at 700 1C. The microstructural evolution and related deformation mechanisms are discussed.
2. Experiment The Ni–Fe base alloy (Ni-(20–30)wt%Fe-(18–25)wt%Cr-(3–7)wt %TiþAl) of 7 kg was prepared via vacuum induction melting. The ingot was homogenized, hot forged and rolled at 1200 1C. Finally, a 10 mm thick plate was attained. All the test specimens used in this study were cut from this plate in the longitudinal direction. The alloy for testing was first solutioned at 1100 1C. Then, aging treatment was carried out at 750 1C for 8 h and furnace cooling
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with a cooling rate of 50 1C/h to 650 1C, and holding at 650 1C for 16 h, with subsequent air cooling. Two specimens were heat treated plus thermal exposure at 700 1C for 1000 h and 7000 h, Table 1 Comparison of yield strength of HT-X with other Ni-base alloys. Alloy
Heat treatment
Yield strength (MPa)
HT-X HT-X HT-X 740 [15] GH2984 [15] Nimonic 263 [15]
Aged Agedþ 700 1C/1000 h Agedþ 700 1C/7000 h Aged Aged Aged
682 731 664 650 539 490
365
respectively. After heat treatment, cylindrical samples, 3 mm in diameter and 6 mm in length, were machined for compressive testing. Compressive tests with a strain rate of 2.5 10 4 s 1 were performed at 700 1C. For each condition, two samples were tested, and the average results were used for discussion. The microstructures were characterized using scanning electron microscopy (SEM) and transmission electron microscopy (TEM). After compressive testing, TEM discs cut perpendicular to stress axis were manually ground to 50 μm, and then perforated by a twin-jet electro-polisher at 40 V/18 mA and 10 1C. The electrolyte consisted of 225 ml acetic acid, 225 ml butylcellosolve and 50 ml perchloric acid. A Tecnai 20 microscope operating at 200 kV was used.
Fig. 1. The microstructures of HT-X after aging treatment (a) and (b), thermal exposure at 700 1C for 1000 h (c) and 7000 h (d) and (e). α-Cr precipitated within the grains after thermal exposure, and its size increased with increasing thermal exposure time. (f) SAD of α-Cr along [111]. (g) EDS spectrum of α-Cr.
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3. Results and discussion 3.1. Yield strength The yield strength of the new alloy (designated as HT-X hereafter) after various heat treatments is shown in Table 1. For comparison, the yield strength of 740, GH2984 and Nimonic 263 from Ref. [15] is also listed in Table 1. All HT-X specimens, even after 7000 h thermal exposure at 700 1C, had greater yield strength than other Ni-base alloys. With increasing thermal exposure time, yield strength of HT-X first increased from 682 MPa to 731 MPa, and then decreased to 664 MPa. 3.2. Microstructural evolution Figs. 1 and 2 show the microstructures after aging treatment and thermal exposure. After aging treatment, the grain size was about 135 μm (see Fig. 1(a)). The alloy mainly consisted of M23C6 type carbides at grain boundaries, and γ0 precipitates dispersed homogeneously in the γ matrix. M23C6 type carbides were indicated by arrows in Fig. 1(b). The γ0 precipitates have an average size of around 18 7 73 nm, as shown in Fig. 2(a). After thermal exposure at 700 1C, the grain size and morphology of carbides had no significant variation, but randomly distributed α-Cr phase precipitated inside the grains. The size of α-Cr increased with increasing thermal exposure time, as seen in Fig. 1 (c) and (d). A few microns to a few tens of microns plate-like α-Cr were formed in the specimen after thermal exposure for 7000 h (see Fig. 1(d) and (e)). Fig. 1(f) is a selected area diffraction (SAD) of bcc α-Cr phase along [111], which can be indexed using lattice parameter a ¼ ¼0.2975 nm [16]. The energy dispersive spectrum (EDS) from α-Cr is shown in Fig. 1(g). Note that Fe, Ni and Mo were also detected in α-Cr. The formation of α-Cr has been reported in other Ni-base superalloys containing higher Cr content [17,18]. The morphology of α-Cr is consistent with present investigations. The coarsening of γ0 precipitates occurred during thermal exposure, as shown in Fig. 2. The average size of γ0 precipitates for thermal exposure of 1000 h and 7000 h is 36 76 nm and 657 11 nm, respectively. 3.3. Deformation microstructures Fig. 3 shows the deformation microstructures of specimen compressed at 700 1C after aging treatment. Dislocation slip bands and dislocation free areas were frequently observed, as shown in Fig. 3(a) and (b). It indicates that the deformation is inhomogeneous. The dislocations may first initiate at the positions with
stress concentration. As plastic deformation continued, dislocations gradually filled the free dislocation areas between slip bands. Generally, the dislocation density in the vicinity of M23C6 carbides at grain boundaries was higher than that in the interior of grains, implying that these carbides impeded dislocation motion and strengthened grain boundaries. In TEM bright field image, Fig. 3 (c), and dark field image, Fig. 3(d), sinuous dislocations were observed, which is a typical character of dislocation slip combing with climb. After thermal exposure at 700 1C for 1000 h, the deformation microstructure was similar with that observed in the specimen after aging treatment (see Fig. 4). However, in the specimen experienced thermal exposure at 700 1C for 7000 h, the deformation was homogeneous, and no dislocation slip bands were observed (see Fig. 5(a)). In TEM dark field images, Fig. 5(b), many dislocation loops around the γ0 precipitates were formed, and stacking faults (SFs) shearing the γ0 precipitates were observed. 3.4. Discussion Both the yield strength and deformation process are closely associated with the microstructural evolution. According to the present investigations, the main difference in microstructures after thermal exposure at 700 1C is the coarsening of γ0 precipitates and the precipitation of α-Cr. For precipitation strengthened alloy, there exists an optimum size of precipitates to attain the maximum strength. Based on the commonly used equations [19], the optimum γ0 precipitate size is calculated: " # 1=2 γ 6γ APB f r Weakly coupled : Δτ ¼ APB f ð1Þ 2b πT Strongly coupled : Δτ ¼
!1=2 rffiffiffi 3 μb 1=2 w 2πγ APB r f 1 2 r π 3=2 wμb2
ð2Þ
Here, Δτ is critical resolved shear stress (CRSS), γAPB is anti-phase boundary (APB) energy, b is Burgers vector of a/2 o1104 , f is the volume fraction of γ0 precipitates, r is the radius of γ0 precipitates, μ is shear modulus, T is the line tension of dislocation (approxi2 mately 1=2μb ), w is a dimensionless constant. Taking γAPB ¼ 0.12 J m 2 [20], b¼ 0.25 nm, f ¼0.17, μ ¼60 GPa [21], w¼1 [22], the theoretical CRSS as a function of γ0 precipitates size can be obtained, as illustrated in Fig. 6. The transition of weakly coupled dislocation mode to strongly coupled occurs at a radius of about 20 nm. This value is very close to the precipitates size, 3676 nm in diameter, after 1000 h thermal exposure. It implies that the size
Fig. 2. Dark field TEM images showing the γ0 precipitates after aging treatment (a), thermal exposure at 700 1C for 1000 h (b) and 7000 h (c). The average size of γ0 precipitates was 18 73 nm, 367 6 nm and 65 711 nm, respectively.
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Fig. 3. The deformation microstructures compressed at 700 1C after aging treatment. Dislocation bands (a) and dislocation free areas (b) were frequently observed. The sinuous dislocations in TEM bright field image (c) and dark field image (d) implied dislocation slip combing with climb.
Fig. 4. Compressive deformation microstructures after 1000 h thermal exposure at 700 1C.
of γ0 precipitates gradually increases to the optimum value during thermal exposure. Thus, the coarsening of γ0 precipitates can well account for the increase of yield strength in specimen after 1000 h thermal exposure. Further coarsening of γ0 precipitates decreases the yield strength, as seen in Fig. 6. Another strengthening
contribution during early stage of thermal exposure may arise from α-Cr precipitation, but the effect of α-Cr precipitates evolution on the strength is still unclear. With increasing thermal exposure time to 7000 h at 700 1C, the deformation process became homogeneous, and SFs sheared γ0
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Fig. 5. TEM bright field image (a) and dark field image (b) showing the compressive deformation microstructures after 7000 h thermal exposure at 700 1C. Stacking faults shearing the γ0 precipitates were indicated by arrows in (b).
Fig. 6. Theoretical critical resolved shear stress vs. radius of γ0 precipitates.
precipitates. According to present investigations, the size of γ0 precipitates increased from 18 73 nm to 657 11 nm after 7000 h thermal exposure. Assuming the volume fraction of γ0 precipitates is a constant, the inter-particle spacing L increases due to γ0 coarsening, therefore Orowan stress Gb=L decreases. The decreasing of Orowan stress facilitates the homogeneous dislocation slip. On the other hand, the larger size of γ0 precipitate, the longer time of climbing over the particle. Thus, for larger γ0 precipitate, the dislocation may not have enough time to climb over it, and dislocation accumulation occurs at the γ/γ0 interface. Stress concentration assists the dislocation dissociation. Many models concern the mechanisms of SF formation in the ordered γ0 precipitates [23–26]. One possibility is that a/2 o110 4 dislocation may dissociate into two Shockley partial dislocations. The leading partial dislocation shears the γ0 precipitates, and stacking fault is formed in the precipitates. Similar dislocation reaction has been reported in a Ni–Co base superalloy [27].
4. Conclusion The compressive tests of a new Ni–Fe base superalloy, HT-X, for applications of A-USC boiler tubes have been conducted at 700 1C after aging treatment and thermal exposure.
The yield strength first increased with thermal exposure time and decreased after further thermal exposure. The optimum radius of γ0 precipitates to obtain the maximum strength was calculated to be about 20 nm, which is consistent with the γ0 precipitates size after 1000 h thermal exposure. The variation of yield strength is mainly attributed to the coarsening of γ0 precipitates. After aging treatment, HT-X mainly consisted of M23C6 carbides and γ0 precipitates. The average grain size was 135 μm, and γ0 precipitates size was 18 73 nm. After thermal exposure, both the grain size and the morphology of M23C6 carbides had no obvious change, but the coarsening of γ0 precipitates from 18 7 3 nm to 65711 nm occurred. Another microstructural difference is that αCr precipitated within the grains, and the size of plate-like α-Cr increased significantly with thermal exposure time. The influence of α-Cr on the mechanical properties of HT-X needs further investigations. After aging treatment and 1000 h thermal exposure at 700 1C, heterogeneous deformation was observed. Dislocation slip combing climb was dominant process. However, after 7000 h thermal exposure at 700 1C, deformation was homogeneous, and stacking faults shearing the γ0 precipitates were observed. The difference is due to the larger γ0 precipitates size, which may cause stress accumulation at γ/γ0 interface and facilitate the dislocation dissociation.
Acknowledgments This work has been financially supported by China Huaneng group, Huaneng Power International, Inc., and the National Natural Science Foundation of China under Grant no. 51301130, 51301131 and 51401071. References [1] K. Nicol, Status of Advanced Ultra-supercritical Pulverised Coal technology, IEA Clean Coal Center, London, 2013. [2] R. Narula, R. Purgert, J. Phillips, H. Wen, Proceedings of the International Conference on Supercritical Technologies and Best Practices, India, 2013. [3] R. Viswanathan, R. Purgert, S. Goodstine, J. Tanzosh, G. Stanko, J.P. Shingledecker, B. Vitalis, Advances in materials technology for fossil power plants, in: Proceedings of the 5th International Conference, ASM International, USA, 2008. [4] D.H. Bechetti, J.N. Dupont, J.J. deBarbadillo, B.A. Baker, Metall. Mater. Trans. A 45 (2014) 3051–3063. [5] R. Viswanathan, K. Coleman, U. Rao, Int. J. Press. Vessel Pip. 83 (2006) 778–783.
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[6] N.D. Evans, P.J. Maziasz, R.W. Swindeman, G.D. Smith, Scr. Mater. 51 (2004) 503–507. [7] Y. Chong, Z.D. Liu, G. Andy, W. Liu, Y.Q. Weng, Mater. Sci. Eng. A 589 (2014) 153–164. [8] Q.Y. Wu, H.J. SONG, R.W. Swindeman, J.P. Shingledecker, V.K. Vasudevan, Metall. Mater.Trans. A 39 (2008) 2569–2585. [9] R. Krishna, S.V. Hainsworth, H.V. Atkinson, A. Strang, Mater. Sci. Technol. 26 (2010) 797–802. [10] M. Akbari-Garakani, M. Mehdizadeh, Mater. Des. 32 (2011) 2695–2700. [11] P.S. Weitzel, J.M. Tanzosh, B. Boring, N. Okita, T. Takahashi, N. Ishikawa, Proceedings of Power-Gen Asia, Thailand, 2012. [12] L.G. Klingensmith, Proceedings of the 4th Symposium on Heat Resistant Steels and Alloys for High Efficiency USC Power Plants, China, 2011. [13] J.J. de Barbadillo, B.A. Baker, R. Gollihue, Proceeding of the 4th Symposium on Heat Resistant Steels and Alloys for High Efficiency USC Power Plants, China, 2011. [14] Y.F. Gu, Technical report on Ni-base superalloys for applications of A-USC power plants (unpublished), 2014. [15] J.T. Guo, X.K. Du, Acta Metall. Sin. 41 (2005) 1221–1227. [16] J.X. Dong, X.S. Xie, Acta Metall. Sin. 41 (2005) 1159–1166. [17] J.F. Radavich, A. Fort, Superalloys 718, 625, 706 and Various Derivatives, The Minerals, Metals & Materials Society, Warrendale, PA, USA, 1994.
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[18] J.F. Radavich, Superalloys 718, 625, 706 and Various Derivatives, The Minerals, Metals & Materials Society, Warrendale, PA, USA, 1997. [19] R.C. Reed, The Superalloys: Fundamentals and Applications, Cambridge University Press, New York (2006) 78–81. [20] M. Dollar, I.M. Bernstein,The minerals, metals & materials society superalloys, in: Proceedings of the 6th International Symposium on Superalloys, 1988. [21] Y. Chong, Z.D. Liu, G. Andy, W. Liu, Y.Q. Weng, Mater. Sci. Eng. A 589 (2014) 153–164. [22] B. Reppich, Acta Metall. 30 (1982) 87–94. [23] Y.H. Zhang, Q.Z. Chen, D.M. Knowles, Mater. Sci. Technol. 17 (2001) 1551–1555. [24] R.R. Unocic, G.B. Viswanathan, P.M. Sarosi, S. Karthikeyan, J. Li, M.J. Mills, Mater. Sci. Eng. A 483-484 (2008) 25–32. [25] L. Kovarik, R.R. Unocic, J. Li, P. Sarosi, C. Shen, Y. Wang, M.J. Mills, Prog. Mater. Sci. 54 (2009) 839–873. [26] R.R. Unocic, N. Zhou, L. Kovarik, C. Shen, Y. Wang, M.J. Mills, Acta Mater. 59 (2011) 7325–7339. [27] Y. Yuan, Y.F. Gu, Z.H. Zhong, T. Osada, C.Y. Cui, T. Tetsui, T. Yokokawa, H. Harada, J, Microscopy 248 (2012) 34–41.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructural evolution and compressive deformation of a new Ni–Fe base superalloy after long term thermal exposure at 700 1C Y. Yuan a,n, Z.H. Zhong b,nn, Z.S. Yu c,a, H.F. Yin a, Y.Y. Dang a, X.B. Zhao a, Z. Yang a, J.T. Lu a, J.B. Yan a, Y. Gu a a
Thermal Power Research Institute Co. Ltd., No. 136, Xingqing Road, Xi'an 710032, China School of Materials Science and Engineering, Hefei University of Technology, No.193, Tunxi Road, Hefei 230009, China c The Key State Laboratory for Mechanical Behavior of Materials, Xi'an Jiaotong University, Xi'an 710049,China b
art ic l e i nf o
a b s t r a c t
Article history: Received 29 July 2014 Received in revised form 24 September 2014 Accepted 26 September 2014 Available online 5 October 2014
A new Ni–Fe base superalloy for applications of advanced ultra-supercritical (A-USC) boiler tubes was heat treated to three different statuses, i.e. aging treatment, aging treatment plus 1000 h and 7000 h thermal exposure at 700 1C. The specimens with various heat treatments have been compressed at 700 1C. The yield strength first increased with thermal exposure time and decreased after further thermal exposure. After aging treatment and 1000 h thermal exposure, deformation was heterogeneous, and dislocation slip combing climb was dominant process. However, after 7000 h thermal exposure, deformation became homogeneous. Stacking faults shearing the γ0 precipitates were also observed. The coarsening of γ0 precipitates and α-Cr precipitation occurred during long term thermal exposure. The influence of microstructural evolution on the yield strength and deformation mechanisms is discussed. & 2014 Elsevier B.V. All rights reserved.
Keywords: Ni–Fe base superalloy Compressive deformation Thermal exposure
1. Introduction Advanced ultra-supercritical (A-USC) technology is an effective method to improve thermal efficiency and reduce CO2, NOx and SOx emissions for coal fired power plants [1]. Currently, the main steam parameters of A-USC power plants have been designed to be 700 1C and 35 MPa in Europe [2], or even higher in USA, i.e. 760 1C and 35 MPa [3]. For 700 1C-class A-USC power plants, the thermal efficiency is projected to be above 50%. The requirements of high temperature materials for use in A-USC boilers are 100,000 h creep strength greater than 100 MPa at 750 1C and 200,000 h coal-ash corrosion resistance of less than 2 mm metal loss [4]. The ferritic and austenitic heat resistant steels cannot meet these requirements. Ni-based superalloys are thought to be the promising candidates for high temperature components in A-USC power plants. For example, Inconel 740/740 H [5–7], CCA617 [8–10], Haynes 282 [11] and Nimonic 263 [12] have been widely evaluated. According to the reported results, 740 H has the highest rupture strength at 700 1C with rupture life of 105 h [13]. In China, the research and development of 700 1C A-USC technology is also ongoing. For high temperature materials, a
n
Corresponding author. Tel.: þ 86 29 8210 2731; fax: þ 86 29 8210 2090. Corresponding author. E-mail addresses: [email protected] (Y. Yuan), [email protected] (Z.H. Zhong). nn
http://dx.doi.org/10.1016/j.msea.2014.09.095 0921-5093/& 2014 Elsevier B.V. All rights reserved.
new Ni–Fe base superalloy, has been developed at Thermal Power Research Institute, Xi'an, China. This new alloy is a candidate for applications of 700 1C A-USC boiler tubes. It is a γ0 strengthened superalloy, and the volume faction of γ0 after aging treatment is around 17%. Its temperature capability with rupture life of 105 h at 100 MPa is roughly 754 1C, comparable with 740 H (765 1C) [14]. This new alloy contains more than 20 wt% Fe, and no Co addition. Compared to 740 H, it has much lower cost and better hot workability. For this new alloy, it is important to deeply understand the microstructures and related deformation mechanisms after long term thermal exposure. In the present study, the compressive testing of the new alloy after aging treatment and long term thermal exposure was conducted at 700 1C. The microstructural evolution and related deformation mechanisms are discussed.
2. Experiment The Ni–Fe base alloy (Ni-(20–30)wt%Fe-(18–25)wt%Cr-(3–7)wt %TiþAl) of 7 kg was prepared via vacuum induction melting. The ingot was homogenized, hot forged and rolled at 1200 1C. Finally, a 10 mm thick plate was attained. All the test specimens used in this study were cut from this plate in the longitudinal direction. The alloy for testing was first solutioned at 1100 1C. Then, aging treatment was carried out at 750 1C for 8 h and furnace cooling
Y. Yuan et al. / Materials Science & Engineering A 619 (2014) 364–369
with a cooling rate of 50 1C/h to 650 1C, and holding at 650 1C for 16 h, with subsequent air cooling. Two specimens were heat treated plus thermal exposure at 700 1C for 1000 h and 7000 h, Table 1 Comparison of yield strength of HT-X with other Ni-base alloys. Alloy
Heat treatment
Yield strength (MPa)
HT-X HT-X HT-X 740 [15] GH2984 [15] Nimonic 263 [15]
Aged Agedþ 700 1C/1000 h Agedþ 700 1C/7000 h Aged Aged Aged
682 731 664 650 539 490
365
respectively. After heat treatment, cylindrical samples, 3 mm in diameter and 6 mm in length, were machined for compressive testing. Compressive tests with a strain rate of 2.5 10 4 s 1 were performed at 700 1C. For each condition, two samples were tested, and the average results were used for discussion. The microstructures were characterized using scanning electron microscopy (SEM) and transmission electron microscopy (TEM). After compressive testing, TEM discs cut perpendicular to stress axis were manually ground to 50 μm, and then perforated by a twin-jet electro-polisher at 40 V/18 mA and 10 1C. The electrolyte consisted of 225 ml acetic acid, 225 ml butylcellosolve and 50 ml perchloric acid. A Tecnai 20 microscope operating at 200 kV was used.
Fig. 1. The microstructures of HT-X after aging treatment (a) and (b), thermal exposure at 700 1C for 1000 h (c) and 7000 h (d) and (e). α-Cr precipitated within the grains after thermal exposure, and its size increased with increasing thermal exposure time. (f) SAD of α-Cr along [111]. (g) EDS spectrum of α-Cr.
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3. Results and discussion 3.1. Yield strength The yield strength of the new alloy (designated as HT-X hereafter) after various heat treatments is shown in Table 1. For comparison, the yield strength of 740, GH2984 and Nimonic 263 from Ref. [15] is also listed in Table 1. All HT-X specimens, even after 7000 h thermal exposure at 700 1C, had greater yield strength than other Ni-base alloys. With increasing thermal exposure time, yield strength of HT-X first increased from 682 MPa to 731 MPa, and then decreased to 664 MPa. 3.2. Microstructural evolution Figs. 1 and 2 show the microstructures after aging treatment and thermal exposure. After aging treatment, the grain size was about 135 μm (see Fig. 1(a)). The alloy mainly consisted of M23C6 type carbides at grain boundaries, and γ0 precipitates dispersed homogeneously in the γ matrix. M23C6 type carbides were indicated by arrows in Fig. 1(b). The γ0 precipitates have an average size of around 18 7 73 nm, as shown in Fig. 2(a). After thermal exposure at 700 1C, the grain size and morphology of carbides had no significant variation, but randomly distributed α-Cr phase precipitated inside the grains. The size of α-Cr increased with increasing thermal exposure time, as seen in Fig. 1 (c) and (d). A few microns to a few tens of microns plate-like α-Cr were formed in the specimen after thermal exposure for 7000 h (see Fig. 1(d) and (e)). Fig. 1(f) is a selected area diffraction (SAD) of bcc α-Cr phase along [111], which can be indexed using lattice parameter a ¼ ¼0.2975 nm [16]. The energy dispersive spectrum (EDS) from α-Cr is shown in Fig. 1(g). Note that Fe, Ni and Mo were also detected in α-Cr. The formation of α-Cr has been reported in other Ni-base superalloys containing higher Cr content [17,18]. The morphology of α-Cr is consistent with present investigations. The coarsening of γ0 precipitates occurred during thermal exposure, as shown in Fig. 2. The average size of γ0 precipitates for thermal exposure of 1000 h and 7000 h is 36 76 nm and 657 11 nm, respectively. 3.3. Deformation microstructures Fig. 3 shows the deformation microstructures of specimen compressed at 700 1C after aging treatment. Dislocation slip bands and dislocation free areas were frequently observed, as shown in Fig. 3(a) and (b). It indicates that the deformation is inhomogeneous. The dislocations may first initiate at the positions with
stress concentration. As plastic deformation continued, dislocations gradually filled the free dislocation areas between slip bands. Generally, the dislocation density in the vicinity of M23C6 carbides at grain boundaries was higher than that in the interior of grains, implying that these carbides impeded dislocation motion and strengthened grain boundaries. In TEM bright field image, Fig. 3 (c), and dark field image, Fig. 3(d), sinuous dislocations were observed, which is a typical character of dislocation slip combing with climb. After thermal exposure at 700 1C for 1000 h, the deformation microstructure was similar with that observed in the specimen after aging treatment (see Fig. 4). However, in the specimen experienced thermal exposure at 700 1C for 7000 h, the deformation was homogeneous, and no dislocation slip bands were observed (see Fig. 5(a)). In TEM dark field images, Fig. 5(b), many dislocation loops around the γ0 precipitates were formed, and stacking faults (SFs) shearing the γ0 precipitates were observed. 3.4. Discussion Both the yield strength and deformation process are closely associated with the microstructural evolution. According to the present investigations, the main difference in microstructures after thermal exposure at 700 1C is the coarsening of γ0 precipitates and the precipitation of α-Cr. For precipitation strengthened alloy, there exists an optimum size of precipitates to attain the maximum strength. Based on the commonly used equations [19], the optimum γ0 precipitate size is calculated: " # 1=2 γ 6γ APB f r Weakly coupled : Δτ ¼ APB f ð1Þ 2b πT Strongly coupled : Δτ ¼
!1=2 rffiffiffi 3 μb 1=2 w 2πγ APB r f 1 2 r π 3=2 wμb2
ð2Þ
Here, Δτ is critical resolved shear stress (CRSS), γAPB is anti-phase boundary (APB) energy, b is Burgers vector of a/2 o1104 , f is the volume fraction of γ0 precipitates, r is the radius of γ0 precipitates, μ is shear modulus, T is the line tension of dislocation (approxi2 mately 1=2μb ), w is a dimensionless constant. Taking γAPB ¼ 0.12 J m 2 [20], b¼ 0.25 nm, f ¼0.17, μ ¼60 GPa [21], w¼1 [22], the theoretical CRSS as a function of γ0 precipitates size can be obtained, as illustrated in Fig. 6. The transition of weakly coupled dislocation mode to strongly coupled occurs at a radius of about 20 nm. This value is very close to the precipitates size, 3676 nm in diameter, after 1000 h thermal exposure. It implies that the size
Fig. 2. Dark field TEM images showing the γ0 precipitates after aging treatment (a), thermal exposure at 700 1C for 1000 h (b) and 7000 h (c). The average size of γ0 precipitates was 18 73 nm, 367 6 nm and 65 711 nm, respectively.
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Fig. 3. The deformation microstructures compressed at 700 1C after aging treatment. Dislocation bands (a) and dislocation free areas (b) were frequently observed. The sinuous dislocations in TEM bright field image (c) and dark field image (d) implied dislocation slip combing with climb.
Fig. 4. Compressive deformation microstructures after 1000 h thermal exposure at 700 1C.
of γ0 precipitates gradually increases to the optimum value during thermal exposure. Thus, the coarsening of γ0 precipitates can well account for the increase of yield strength in specimen after 1000 h thermal exposure. Further coarsening of γ0 precipitates decreases the yield strength, as seen in Fig. 6. Another strengthening
contribution during early stage of thermal exposure may arise from α-Cr precipitation, but the effect of α-Cr precipitates evolution on the strength is still unclear. With increasing thermal exposure time to 7000 h at 700 1C, the deformation process became homogeneous, and SFs sheared γ0
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Fig. 5. TEM bright field image (a) and dark field image (b) showing the compressive deformation microstructures after 7000 h thermal exposure at 700 1C. Stacking faults shearing the γ0 precipitates were indicated by arrows in (b).
Fig. 6. Theoretical critical resolved shear stress vs. radius of γ0 precipitates.
precipitates. According to present investigations, the size of γ0 precipitates increased from 18 73 nm to 657 11 nm after 7000 h thermal exposure. Assuming the volume fraction of γ0 precipitates is a constant, the inter-particle spacing L increases due to γ0 coarsening, therefore Orowan stress Gb=L decreases. The decreasing of Orowan stress facilitates the homogeneous dislocation slip. On the other hand, the larger size of γ0 precipitate, the longer time of climbing over the particle. Thus, for larger γ0 precipitate, the dislocation may not have enough time to climb over it, and dislocation accumulation occurs at the γ/γ0 interface. Stress concentration assists the dislocation dissociation. Many models concern the mechanisms of SF formation in the ordered γ0 precipitates [23–26]. One possibility is that a/2 o110 4 dislocation may dissociate into two Shockley partial dislocations. The leading partial dislocation shears the γ0 precipitates, and stacking fault is formed in the precipitates. Similar dislocation reaction has been reported in a Ni–Co base superalloy [27].
4. Conclusion The compressive tests of a new Ni–Fe base superalloy, HT-X, for applications of A-USC boiler tubes have been conducted at 700 1C after aging treatment and thermal exposure.
The yield strength first increased with thermal exposure time and decreased after further thermal exposure. The optimum radius of γ0 precipitates to obtain the maximum strength was calculated to be about 20 nm, which is consistent with the γ0 precipitates size after 1000 h thermal exposure. The variation of yield strength is mainly attributed to the coarsening of γ0 precipitates. After aging treatment, HT-X mainly consisted of M23C6 carbides and γ0 precipitates. The average grain size was 135 μm, and γ0 precipitates size was 18 73 nm. After thermal exposure, both the grain size and the morphology of M23C6 carbides had no obvious change, but the coarsening of γ0 precipitates from 18 7 3 nm to 65711 nm occurred. Another microstructural difference is that αCr precipitated within the grains, and the size of plate-like α-Cr increased significantly with thermal exposure time. The influence of α-Cr on the mechanical properties of HT-X needs further investigations. After aging treatment and 1000 h thermal exposure at 700 1C, heterogeneous deformation was observed. Dislocation slip combing climb was dominant process. However, after 7000 h thermal exposure at 700 1C, deformation was homogeneous, and stacking faults shearing the γ0 precipitates were observed. The difference is due to the larger γ0 precipitates size, which may cause stress accumulation at γ/γ0 interface and facilitate the dislocation dissociation.
Acknowledgments This work has been financially supported by China Huaneng group, Huaneng Power International, Inc., and the National Natural Science Foundation of China under Grant no. 51301130, 51301131 and 51401071. References [1] K. Nicol, Status of Advanced Ultra-supercritical Pulverised Coal technology, IEA Clean Coal Center, London, 2013. [2] R. Narula, R. Purgert, J. Phillips, H. Wen, Proceedings of the International Conference on Supercritical Technologies and Best Practices, India, 2013. [3] R. Viswanathan, R. Purgert, S. Goodstine, J. Tanzosh, G. Stanko, J.P. Shingledecker, B. Vitalis, Advances in materials technology for fossil power plants, in: Proceedings of the 5th International Conference, ASM International, USA, 2008. [4] D.H. Bechetti, J.N. Dupont, J.J. deBarbadillo, B.A. Baker, Metall. Mater. Trans. A 45 (2014) 3051–3063. [5] R. Viswanathan, K. Coleman, U. Rao, Int. J. Press. Vessel Pip. 83 (2006) 778–783.
Y. Yuan et al. / Materials Science & Engineering A 619 (2014) 364–369
[6] N.D. Evans, P.J. Maziasz, R.W. Swindeman, G.D. Smith, Scr. Mater. 51 (2004) 503–507. [7] Y. Chong, Z.D. Liu, G. Andy, W. Liu, Y.Q. Weng, Mater. Sci. Eng. A 589 (2014) 153–164. [8] Q.Y. Wu, H.J. SONG, R.W. Swindeman, J.P. Shingledecker, V.K. Vasudevan, Metall. Mater.Trans. A 39 (2008) 2569–2585. [9] R. Krishna, S.V. Hainsworth, H.V. Atkinson, A. Strang, Mater. Sci. Technol. 26 (2010) 797–802. [10] M. Akbari-Garakani, M. Mehdizadeh, Mater. Des. 32 (2011) 2695–2700. [11] P.S. Weitzel, J.M. Tanzosh, B. Boring, N. Okita, T. Takahashi, N. Ishikawa, Proceedings of Power-Gen Asia, Thailand, 2012. [12] L.G. Klingensmith, Proceedings of the 4th Symposium on Heat Resistant Steels and Alloys for High Efficiency USC Power Plants, China, 2011. [13] J.J. de Barbadillo, B.A. Baker, R. Gollihue, Proceeding of the 4th Symposium on Heat Resistant Steels and Alloys for High Efficiency USC Power Plants, China, 2011. [14] Y.F. Gu, Technical report on Ni-base superalloys for applications of A-USC power plants (unpublished), 2014. [15] J.T. Guo, X.K. Du, Acta Metall. Sin. 41 (2005) 1221–1227. [16] J.X. Dong, X.S. Xie, Acta Metall. Sin. 41 (2005) 1159–1166. [17] J.F. Radavich, A. Fort, Superalloys 718, 625, 706 and Various Derivatives, The Minerals, Metals & Materials Society, Warrendale, PA, USA, 1994.
369
[18] J.F. Radavich, Superalloys 718, 625, 706 and Various Derivatives, The Minerals, Metals & Materials Society, Warrendale, PA, USA, 1997. [19] R.C. Reed, The Superalloys: Fundamentals and Applications, Cambridge University Press, New York (2006) 78–81. [20] M. Dollar, I.M. Bernstein,The minerals, metals & materials society superalloys, in: Proceedings of the 6th International Symposium on Superalloys, 1988. [21] Y. Chong, Z.D. Liu, G. Andy, W. Liu, Y.Q. Weng, Mater. Sci. Eng. A 589 (2014) 153–164. [22] B. Reppich, Acta Metall. 30 (1982) 87–94. [23] Y.H. Zhang, Q.Z. Chen, D.M. Knowles, Mater. Sci. Technol. 17 (2001) 1551–1555. [24] R.R. Unocic, G.B. Viswanathan, P.M. Sarosi, S. Karthikeyan, J. Li, M.J. Mills, Mater. Sci. Eng. A 483-484 (2008) 25–32. [25] L. Kovarik, R.R. Unocic, J. Li, P. Sarosi, C. Shen, Y. Wang, M.J. Mills, Prog. Mater. Sci. 54 (2009) 839–873. [26] R.R. Unocic, N. Zhou, L. Kovarik, C. Shen, Y. Wang, M.J. Mills, Acta Mater. 59 (2011) 7325–7339. [27] Y. Yuan, Y.F. Gu, Z.H. Zhong, T. Osada, C.Y. Cui, T. Tetsui, T. Yokokawa, H. Harada, J, Microscopy 248 (2012) 34–41.