MICROSTRUCTURAL STABILITY OF ALLOY 617 MOD. DURING THERMAL AGING

Proceedings of the ASME Symposium on Elevated Temperature Application of Materials for Fossil, Nuclear, and Petrochemical Industries March 25-27, 214 ETAM214-11 MICROSTRUCTURAL STABILITY OF ALLOY 617 MOD. DURING THERMAL AGING GUO Yan ZHOU Rongcan Institute Co. Ltd.,Xi an 7132, Institute Co. Ltd.,Xi an 7132, TANG Liying HOU Shufang WANG Bohan Institute Co. Ltd.,Xi an 7132 Institute Co. Ltd., Xi an 7132 Institute Co. Ltd. Xi an 7132 investigated, and several have emerged as promising candidates for A-USC power plants due to their excellent creep-rupture strength, high oxidation and hot corrosion resistance, including Alloy 617mod., Nimonic 263 and Inconel 74H[5 8]. Alloy 617 mod. is a modified version of Alloy 617. Alloy 617 belongs to the group of nickel-base alloys usually referred to as solid-solution strengthened materials [9,1]. Microstructure evolution during creep of Alloy 617 has been reported [11, 12]. The effect of microstructure on the hardness of Alloy 617 was investigated by Wu et al. [13, 14]. References [7, 14] reported microstructure and creep property of Alloy 617 mod. The properties of Alloy 617 mod. are strongly influenced by microstructural changes that occur in the process of aging. Thus, this paper focuses on microstructural stability of Alloy 617 mod. and its effect on certain mechanical properties such as hardness and toughness. ABSTRACT The microstructural stability of Alloy 617 mod. during aging for up to 3 h at temperatures of 7oC and 75oC was investigated. The precipitates of the aged alloy included M23C6 carbides located both inside grains (intragranular) and at grain boundaries (intergranular) and γ phase dispersed within grains. During aging, the intergranular precipitates showed a good stability. Intragranular particles increased substantially after the aging for 3h at 75oC. Inter and intragranular carbide particles resulted in the precipitation hardening of the aged alloys. The precipitation of γ phase particles during aging at 7 C and 75 C is also an important factor for an enhanced hardness and an obvious decrease of the impact absorbed energy. Additionally, the intergranular cracks apparently lead to a decrease in the impact absorbed energy for the aged alloys due to carbide particles at grain boundaries. The impact absorbed energies of the aged alloys were fairly stable within the dwelling time from 3 h to 3 h and were in the range of 63~65J and 75~83J for the 7oC and 75oC aging, respectively. 2. EXPERIMENTAL The chemical composition of the Alloy 617 mod. investigated in this study is given in Table 1. Table 1 Chemical composition of Alloy 617 mod. KEY WORDS Alloy 617mod., microstructure, aging, hardness, impact absorbed energy 1. INTRODUCTION In recent years, efficient ultra supercritical (USC) units have operated at steam temperatures of approximately 6 with steam pressures of 25 3 MPa. Certain advanced ferric and austenitic stainless steels are widely used in such units, e.g. P92, TP347HFG, S3432 and TP31HNbN. Currently there is a strong incentive to increase the steam conditions in order to improve the efficiency from the viewpoint of energy conservation and reduction of CO2 emissions. Advanced Ultra Supercritical (A-USC) power plants with target steam conditions of 7 C or higher and 35 MPa are under consideration. The heat-resistant steels above mentioned are not suitable for service in such conditions. Thus, collaborative projects to investigate alternative materials have been conducted in Europe, Japan and the United States [1-4]. A number of austenitic steels and nickel-base alloys have been Co Fe Ti Al B Cu W 12..97.41.84.6.2.14 C Si Mn Cr Mo Ni Pb.6.15.1 21.9 8.9 54.8.4 Nb S P Sn Sb Bi As.6.8.7.9.2.2.8 The as-received alloy had been solution treated at 118 C for 3 min. It was then aged for 3, 1 and 3h at 7 C and 75 C, respectively. The hardness of the tested samples was measured with a HB-3C Brinell hardness tester. The 55 mm 1 mm 7.5 mm V-notched Charpy impact samples were used to measure the absorbed energy at room temperature on a PKP45 impact testing machine. For the as-received alloy, metallographic samples were ground to 1-grit and mechanically polished, and then etched using a solution of HCl, HNO3 and H2O. For the aged alloy, metallographic samples 1

were ground to 15-grit and electro-polished with a solution of H2SO4 and CH3OH, then electro-etched with a solution of H2SO4 and H2O. Metallographic inspections were carried out using a FEI Quanta-4HV scanning electron microscope (SEM). Energy dispersive spectroscopy (EDS) micro-analysis was performed in SEM. Fracture surfaces of the impact tested samples were also analyzed by SEM. The samples used for transmission electron microscopy (TEM) analysis were prepared as follows. A foil of about 5 μm thick was cut and mechanically ground to a thickness of 4μm, from which TEM disks of 3 mm in diameter were punched. Twin-jet electro-polishing was performed using a 1% perchloric acid+9% acetic acid solution at room temperature, with a polishing current of 3 ma. TEM observation was carried out in a JEM-2CX TEM operating at 2 kv. The phase identification was performed using the selected area electron diffraction (SAED) pattern. 24 7 C 75 C Impact aborbed energy/j 2 16 12 8 4 5 1 15 2 25 3 Aging time/h Fig.2 Effect of aging time and temperature on the impact absorbed energy of the alloys Fig.3 shows the SEM photographs of the fractured surfaces after the impact test of the alloy under different conditions. In the case of the as-received alloy, the fracture was predominantly ductile in nature, as indicated by the occurrence of large dimples. After aging for 3 h or more, the fracture mode changed from ductile to brittle because of the carbides formed at grain boundaries after aging. At the same aging temperature, the number and size of dimples remained unchanged with an increase in the aging time. On the other hand, as for the alloy aged for the same time, an increase in amount of dimpled surface with rising aging temperature is indicative of an enhanced toughness. 3. Results The hardness of the as-received and the aged samples at room temperature is presented in Fig. 1. The alloy had a low hardness value in the as-received condition. A substantial increase in hardness was observed during the first period of the aging (between and 3h). After that, the hardness changed slightly for the aged alloys in the period of 3 3h and were in the range of 225~236HB and 25~214HB in the process of aging at 7oC and 75oC, respectively. Fig.2 reveals the effect of aging temperatures and times on the absorbed energy of the alloy, which is the characterization for the toughness. The as-received alloy exhibited the highest toughness and the drop in toughness occurred predominately during the first aging period. Then the absorbed energies of the aged alloys were fairly stable within the dwelling time from 3 h to 3 h and were in the range of 63~65J and 75~83J for the 7oC and 75oC aging, respectively. The toughness of the Alloy 617 mod. is greater than that of Alloy 617 and Inconel 74H [15, 16]. 7 C 75 C 28 24 Hardness/HB 2 16 12 8 4 5 1 15 2 25 3 Aging time/h Fig.1 Effect of the aging temperature and time on the hardness of the alloys 2

Fig.3 The fractured surfaces after the impact test for Alloy 617 mod.: (a) as-received; (b) aged for 3 h at 7 oc; (c) aged for 3 h at 7 oc; (d) aged for 3 h at 75 oc; (e) aged for 3 h at 75 oc Fig.4 reveals the microstructural evolution with aging temperature and time, as viewed in the SEM. Some large particles distributed within grains with diameters of 2.4~3.3μm were rich in Ti and N, and were therefore probably TiN. The amount of small precipitates increased clearly with prolonging aging time and temperature. Especially after 3 h aging at 75, an obvious increase of the precipitate particles was observed (Fig.4d). Fig.5 TEM images and SAED patterns of the Alloy 617 mod. aged for 3h at 7 oc TEM images and SAED patterns of the alloy aged for 3 h at 7 is shown in Fig.6. Fig.6 (a) (b) depicted M23C6 and spherical γ particles within grains, with diameters of 3 14 nm and 6 9 nm, respectively. The precipitation of M23C6 particles with a length of 14~4 nm at discontinuous grain boundaries was visible in Fig.6(c) and (d). Compared with Fig.5, it is apparent that increasing the aging time induced an increase in γ volume fraction [14]. Fig.4 SEM images of the Alloy 617 mod. with different conditions: (a) aged for 3 h at 7 oc; (b) aged for 3 h at 7 oc; (c) aged for 3 h at 75 oc; (d) aged for 3 h at 75 oc Fig.6 TEM images and SAED patterns of the Alloy 617 mod. aged for 3h at 7 oc Fig.5 shows the TEM images and SAED patterns of the alloy aged for 3 h at 7. As presented in Figs.5 (a) (b), M23C6 with a size of 2 11 nm and spherical γ particles with a diameter of 2 5 nm precipitated within grains. The precipitation of M23C6 particles of 1~3 nm discontinuously at grain boundaries was visible in Fig.5(c). These carbide particles are a complex face centered cubic (fcc) structure dispersed within the matrix, containing elements of Cr, Mo and Fe [9]. Fig.7 presents TEM images and SAED patterns of the alloy aged for 3 h at 75. The size of M23C6 and γ particles were 25 13 nm and 2 7nm, respectively (see Fig.7(a)). Fig.7 (b) showed the dark-field image of M23C6 particles. In contrast with Fig.5, it is indicative of a decrease of the γ volume fraction with an increase in the aging temperature as for the aged alloys. 3

Fig.7 TEM images and SAED patterns of the Alloy 617 mod. aged for 3 h at 75 o C: (a) intragranular M23C6 and γ, (b) dark-field image of intragranular M23C6 Fig. 8 reveals the morphologies of the precipitates of the alloys aged for 3h at 75 C, as viewed in the TEM. There were M 23 C 6 particles with a length of 5~33 nm and γ particles with a diameter of 7 ~ 13 nm within grains (Fig.8(a)(b)). M 23 C 6 particles with a size of 27~5 nm at grain boundaries was visible in Fig.8(c) and (d), which is smaller than that of Alloy 617. The size of intragranular particles of the aged alloys under different aging temperatures and times is summarized in Table 2. Fig.8 TEM images and SAED patterns of the Alloy 617 mod. aged for 3 h at 75 C: (a) γ phase inside grain,(b) M 23 C 6 and γ phase inside grain, (c)(d) M 23 C 6 phase at grain boundaries Table 2 Effect of the aging temperatures and time on size of the precipitates inside grains Temperatu re ( C) 7 75 Time (h) γ (nm) M23C6 (nm) 3 2~5 2~11 1 4~65 25~13 3 6~9 3~14 3 2~7 25~13 1 4~8 5~15 3 7~13 5~33 4. Discussion From the above results, it was clear that the precipitates of the aged alloy were M 23 C 6 with a complex fcc structure and γ phase with an ordered fcc structure. M 23 C 6 particles were located both at grain boundaries and inside grains, while γ phase precipitated within grains. Increasing the aging temperature and prolonging the aging time resulted in an increase of the size of the γ phase particles but did not affect coherency between γ and matrix during aging. The precipitates discontinuously distributed at grain boundaries underwent no obvious change during the aging. Reference [14] reported the number of γ particles in the aged Alloy 617 mod. increased greatly after 3 h aging at 7 and 75 C; however, γ particles decreased for the alloy aged for 3h at 8 C. The alloy aged for 3h exhibited an enhanced hardness compared with as-received alloy. In the period of 3-3h, the hardness of the aged alloy remained almost unchanged at 7 and 75. The hardening of the aged alloy correlates well with the distribution and size of the precipitates and with the nature of the bonding between the precipitates and the matrix. For example, coherency between γ phase and the matrix was maintained during aging and resulted in a coherent stress for the aged alloys, thus giving rise to an enhanced hardness. On the other hand, the hardening of the aged alloy also results from the precipitation of carbides as discrete particles both inside grains and at grain boundaries. The intragranular carbides contribute to the strengthening effect by acting as barriers for dislocation motion and by stabilizing dislocations. The discrete nature of the grain boundary carbides enhances hardness because it causes pinning of the boundary and decreases grain boundary sliding. The as-received alloy exhibited the highest toughness and the drop in toughness occurred predominately during the first aging period. This is believed to be the result of the carbides formed at grain boundaries after aging. The grain boundaries are weakened by grain boundary carbides and the separation occurs by decohesion of the carbide/matrix interface. At the same time, the separations between the carbides and the matrix serve as initiation sites for fracture. Hence, cracks may mainly take place at grain boundaries during the impact test. The separated surfaces appear smooth due to the distribution of the grain boundary carbides. This inter-granular weakness appears to be the cause of the decrease in toughness for the aged alloys. On the other hand, the precipitate strengthening of γ phase of the aged alloys limits plastic deformation to the area near grain boundaries, resulting in the occurrence of crack along grain boundaries, which is an another important factor for a decrease of toughness. During aging, the distribution and size of carbides along grain boundaries exhibited no substantial change. The toughness of the aged alloy maintained a high level and remained almost unchanged in the period of 3-3h that is manifested by the unchanged number and size of dimples for the aged alloys at 7 and 75, respectively. It is likely that boron is beneficial with respect to the bonding at grain 4

boundaries and retards the initiation and propagation of cracks. Reference [17] reported that addition of boron changed the fracture mode and dramatically improved the strength and ductility. Compared with microstructure of the alloys aged at 7 and 75, it was evident that after aging for 3 h, increasing aging temperature induced a decrease of γ volume fraction, thus leading to a decrease in hardening effect and increased toughness. 5. Conclusions 1) The precipitates of the aged alloy were M 23 C 6 carbides located both inside grains and at grain boundaries and γ phase dispersed within grains. 2) During aging, the distribution and size of the precipitates along grain boundaries underwent no obvious change. Increasing the aging temperature and prolonging the aging time caused growth of γ particles but did not affect coherency between γ and matrix. 3) An enhanced hardness and a relatively high level of absorbed energy were achieved for the aged Alloy 617mod. compared to Alloy 617. ACKNOWLEDGMENTS This work was supported by National Energy Applied Technology Research &Demonstration Project (NY21112-1), National High Technology Research and Development Program (212AA551) and CSEE Youth Science & Technology Innovation Project (No.3) REFERENCES Put references here. [1] Bugge J, Kjær S, Blum R. High-efficiency coal-fired power plants development and perspectives [J].Energy, 26,31 (1):1437 1445. [2] Roster J,Gotting M,Del Genovese D,et al. Wrought Ni-base superalloys for steam turbine applications beyond 7 o C [J].Advanced Engineering Materials,25,5(7) :469-483. [3] C. DavidTung, C. John Lippold, Proc of the12th International Symposium on Superalloys, SevenSprings, PA, US, Sept. 9 13, 212. [4] J. Bratberg, H. Mao, L. Kjellqvist, A. Engstr om, P.Mason and Q. Chen, Proc of the 12th International Symposium on Superalloys, Seven Springs, PA, US, Sept. 9-13, 212. [5] J. Sauders,M.Monteriro,F.Rizzo,et al. The oxidation behavior of metals and alloys at high temperatures in atmospheres containing water vapor: A review [J]. Progress in Materials Science,28,53 (5):775-837. [6] D.Tung,J. Lippold.Weld Solidification behavior of Ni-base superalloys for use in advanced supercritical coal-fired power plants[j].proceeding to the 12th International Symposium on Superalloys, Seven Springs Mountain Resort, Seven Springs, PA, U.S., September 9-13, 212. [7] Maile K. Use of advanced Alloy 617 mod for critical components of the future 7 coal fired power plant [J].29 Symposium on Advanced Power Plant Heat Resistant Steels and Alloys, Shanghai, China, 29. [8] S.Q. Zhao, Y. Jiang, J.X. Dong and X.S. Xie, Experimental Investigation and Thermodynamic Calculation on Phase Precipitation of INCONEL 74[J]. Acta Metall. Sin. (Engl. Lett.) 19(26) 425. [9] Klöwer J, Husemann R U, Bader M, et al. Evelopment of Nickel based on alloy 617 for components in 7 power plants[j].6th International conference on creep, fatigue and creep-fatigue Interaction,India, 212. [1] Gariboldi E, Cabibbo M, Spigarelli S.Investigation on precipitation phenomena of Ni-22Cr-12Co-9Mo alloy aged and crept at high temperature[j].international Journal of Pressure Vessels and Piping,28,85(1/2):63-71. [11] S. Rahman, G. Priyadarshan, K.S. Raja, C. Nesbittand M. Misra, Investigation of the secondary phases of alloy 617 by scanning Kelvin probe force microscope[j]. Mater. Lett. 62 (28) 2263. [12] J. Roster, M. Gotting, D. Del Genovese, et al. Wrought Ni-base superalloys for steam turbine applications beyond 7 [J]. Adv. Eng. Mater. 5 (23) 469. [13] Q. Wu, H. Song, W. Robert, J. P. Shingledecker,V.K. Vasudevan, Microstructure of long-term aged IN617Ni-base superalloy[j]. Metall. Mater. Trans. A 39 (28)2569. [14] Q. Wu,V. K. Vasudevan, J. P. Shingledecker, R. W. Swindeman, Microstructural characterization of advanced boiler materials for ultra supercritical coal power plants[j].proceedings to the Fourth International Conference on Advances in Materials Technology for Fossil Power Plants,25. [15] Y. GUO, B.H. WANG, S.F.HOU. Aging Precipitation Behavior and Mechanical Properties of Inconel 617 Superalloy [J]. Acta Metall. Sin. (Engl. Lett.), 213, 26 (3): 37-312. [16] Xie Xishan, Zhao Shuangqun, Dong Jianxin, et al.structural stability and improvement of Inconel alloy 74 for ultra supercritical power plants[j].journal of Chinese Society of Power Engineering, 211, 31(8):638-643. [17] D. Mukherji, J. Rösler, M. Krüger, et al. The effects of boron addition on the microstructure and mechanical properties of Co-Re-based high-temperature alloys[j]. Scripta Materialia 212, 66(1): 6-63. Corresponding author: GUO Yan Senior Engineer, Ph.D.; Tel: +86 29 821241-2; Fax: +86 29 83255651; E-mail address:guoyan9732@gmail.com 5