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Temperature effect on the creep behavior of alloy 617 in air and helium environments
Nuclear Engineering and Design 271 (2014) 291–300
Contents lists available at ScienceDirect
Nuclear Engineering and Design journal homepage: www.elsevier.com/locate/nucengdes
Temperature effect on the creep behavior of alloy 617 in air and helium environments Woo-Gon Kim a,∗ , Jae-Young Park b , Gyeong-Geun Lee a , Sung-Deok Hong a , Yong-Wan Kim a a b
Korea Atomic Energy Research Institute, Daejeon 305-353, Republic of Korea Pukyong National University, Busan 608-739, Republic of Korea
a b s t r a c t The temperature effect on creep and oxidation behaviors in air and helium (He) environments was investigated at 950, 900 and 850 ◦ C for Alloy 617, which is considered as a prime candidate material for VHTR components. Creep data were obtained with different stress levels at the three temperatures. Oxidation microstructural features such as the surface oxide layer, internal oxidation and decarburization were analyzed by observing each crept specimen. At 950 and 900 ◦ C, the creep rupture time in He environment was shorter than that in air, and the thickness of the surface oxide layer was thicker than in air. The deterioration of creep resistance in the He environment was due to a thicker oxide-layer thickness, which reduced the effective area carrying the creep load. On the other hand, at 850 ◦ C, the creep rupture time was almost the same regardless of both environments. The surface oxide-layer thickness was reduced by about 50% compared to the temperatures of 950 and 900 ◦ C. The temperature of 850 ◦ C could be assumed as a boundary temperature at which the He effects disappeared. It was found that the relationship between the creep data and surface oxide-layer thickness was in accordance with both environments. © 2013 Elsevier B.V. All rights reserved.
1. Introduction A very high temperature gas reactor (VHTR) is one of the Gen-IV reactors aiming at safe, long-lived, proliferation-resistant and economical nuclear power plants. Its high operating temperature of over 800 ◦ C enables a high energy efficiency and the production of hydrogen gas using Sulfur–Iodine process. The heat of the primary helium (He) circuit transfers to the secondary helium loop through the intermediate heat exchanger (IHX). The IHX component needs high-temperature creep resistance in the He environment, and requires good oxidation resistance, corrosion resistance, and phase stability at high temperatures. Among them, creep and oxidation properties are importantly considered, as the integrity of the components should be preserved during a prolonged period in a VHTR coolant (Kim et al., 2007, 2008, 2009, 2010, 2012a,b). Of the existing alloys, nickel-base Alloy 617 is the leading candidate for use in the next generation nuclear plant (NGNP) heat exchangers because it has the highest creep strength of solid solution alloys under consideration for temperature above 850 ◦ C. It was reported that the creep rupture time varied widely with the environments interactions between He impurities and the alloy (Jang et al., 2008; Dewson and Li, 2005; Cook, 1984). The impurities in He
∗ Corresponding author at: Korea Atomic Energy Research Institute, 989-111 Daedeok-daero, Yuseong-gu, Daejeon 305-353, Republic of Korea. Tel.: +82 42 868 2493. E-mail address: [email protected] (W.-G. Kim). 0029-5493/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.nucengdes.2013.11.050
coolant during steady-state operation of the VHTRs were reported as H2 O, H2 , CO, CO2 , CH4 , O2 and N2 . However, it has not been well established on the temperature effect in He environment to be drastically affected to the rupture time. The relationships between creep and oxidation features are not well understood. Therefore, it is necessary to investigate the temperature effect on the creep and oxidation behaviors in air and He environments of Alloy 617. In this study, the creep behaviors of Alloy 617 were comparatively investigated by creep tests with different stresses in air and He environments at 950, 900, and 850 ◦ C. The temperature effect on the creep and oxidation behaviors was evaluated, and relationships between these behaviors were found and discussed. 2. Experimental procedures Commercial grade nickel-based superalloy, Alloy 617 (Inconel 617), was used in this study. The material was a hot-rolled plate with a thickness of 15.875 mm. Creep test specimens in the air and He environments were fabricated in cylindrical form with a 30 mm gauge length and 6 mm diameter. The gage section was parallel to the longitudinal rolling direction. Circular grooves were machined at both ends beyond the shoulder region of these specimens to attach extensometers for monitoring the elongation during creep testing. The loading frames used in the creep tests had a lever arm ratio of 20:1. The overall creep rates were determined from elongations measured by extensometers equipped at the circular grooves. Creep strain data with elapsed times were taken automatically by
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3. Results and discussion 3.1. Comparison of creep behaviors in air and He environments at 950, 900 and 850 ◦ C The creep property data of Alloy 617 such as the rupture time, minimum creep rate, rupture elongation, and reduction of area were obtained in the air and He environments at 950, 900, and 850 ◦ C. The creep results in both environments were compared at the three temperatures.
Fig. 1. Creep testing apparatus in He environment.
Fig. 2. Typical creep curves obtained in air and He environments at 900 ◦ C.
a PC. K-type thermocouples were used to monitor temperatures within the gage section of the specimens. A split furnace was used to heat the specimens of Alloy 617 in the air and He environments to 950, 900, and 850 ◦ C. The pull rod and jig used for the creep tests were manufactured with a nickel-based superalloy material to sufficiently endure oxidation and thermal degradation during creep. Before creep tests in the He environment, a vacuum chamber made for the quartz tube was purged three or four times by a vacuum pump to remove some impurities in the chamber. During the creep tests, pure He gas of 99.999% was supplied to the specimens equipped in the quartz tube. Impurity concentration in pure He gas was O2 < 1.0 ppm, N2 < 5.0 ppm and H2 O < 1.0 ppm. Flow rate of the He gas during the creep test was controlled under 20 cm3 /min. The control of the other impurities such as CO2 , CO, CH4 , and H2 in He gas simulating a VHTR condition will be investigated in the further tests. Creep testing apparatus in the He environment is shown in Fig. 1. Photos showing each component such as the quart tube, He flow meter, strain indicator, specimen and jigs, furnace controller, and data acquisition PC are presented.
Fig. 3. A photo of crept specimen under 18 MPa at 950 ◦ C in air of Alloy 617.
Fig. 4. Comparison of stress vs. rupture time in the air and He environments at (a) 950 ◦ C, (b) 900 ◦ C, and (c) 850 ◦ C.
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Fig. 5. Comparison of minimum creep rate vs. stress in air and He environments at 950, 900, and 850 ◦ C. Fig. 6. Comparison of rupture time vs. minimum creep rate in air and He environments at 950, 900, and 850 ◦ C.
Fig. 2 shows typical creep curves for stresses in the air and He environments at 900 ◦ C. Alloy 617 showed little primary creep strain, and a well-defined secondary creep stage was not exhibited in the full creep curves. The onset of a tertiary creep was unclear, and a tertiary creep stage was initiated from a low strain level. There were no differences in the shapes of the creep curves between the air and He environments. Alloy 617 revealed good ductility (>30%) in spite of the long duration above 104 h at 950, 900, and 850 ◦ C.
From a series of creep tests, long-term creep data reaching 14,000 h at 950 ◦ C were achieved in the air, as shown in Fig. 3. Creep rupture occurred at three positions within the gage length. The rupture elongation was about 30%, and the reduction of area was as small as 22%. It was identified that the creep rupture was developed by the incorporation of the many small cavities formed during the creep, rather than the necking of the specimen.
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Fig. 8. Typical OM micrographs of a crept specimen under 28 MPa in air at 950 ◦ C.
Fig. 7. Comparison of creep rupture elongation vs. stress in air and He environments at 950, 900, and 850 ◦ C.
Fig. 4 shows a comparison of the log stress vs. log rupture time in the air and He environments at 950, 900 and 850 ◦ C. For 950 and 900 ◦ C, in the high stress ranges above about 30 MPa, the creep stress between the air and He environments was almost similar. However, in the low stress ranges below about 30 MPa, there were differences in the creep stress (or creep rupture time)
between the two environments. Namely, the creep stress in air was higher than that in He environment, but for 850 ◦ C, the creep data of the log–log plot was almost superimposed regardless of both environments, as shown well in the figure. Therefore, it is obvious that the creep properties in air and He environments have a temperature effect. The reason for this is closely related to the oxide layer thickness formed during the creep. The oxide-layer thicknesses in both environments are measured at the three temperatures, and the results are discussed in a later section. Fig. 5 shows a comparison of the minimum creep rate vs. stress at 950, 900, and 850 ◦ C in air and He environments. The relationships between the minimum creep rate and stress revealed a good linearity regardless of both environments at the three temperatures. It is clear that Alloy 617 followed Norton’s power law well under this creep condition because the deformation map corresponded to the power-law creep region. The creep mechanism is governed by a climb of dislocation (Kim et al., 2010). The creep rates between both environments at 950 and 900 ◦ C plot a significant difference at low stress levels. The creep rate in He was faster than that in air. However, the creep rates at 850 ◦ C were almost the same regardless of both environments. Fig. 6 shows a plot of Monkman–Grant (M–G) relationships. In the equation of M–G relationship, log tr + mε˙ m = C, the m values of the slope were obtained in the air and He environments at the three temperatures of 950, 900, and 850 ◦ C. At 950 and 900 ◦ C, it means that the creep rupture time of the He specimens was shorter than that of the air one. But at 850 ◦ C, their values were almost equal to
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Fig. 9. Precipitates formed in the grip-section region and gage-section region of the specimen crept in He environment under 20 MPa at 950 ◦ C.
m = 1.0 regardless of the environment. This means that the creep properties were not changed at 850 ◦ C under the He environment. It is obvious that the He effects on the creep properties disappeared at 850 ◦ C. This is assumed to be a temperature effect on the creep properties under this He-agent condition. Thus, the temperature of 850 ◦ C can be regarded as a boundary temperature at which the He effects disappeared. The data on the creep rupture elongation were similar in both environments, although the data were scattered, as shown in Fig. 7.
3.2. Creep fracture micrographs Fig. 8 shows typical OM micrographs at two magnifications of the crept specimen under 28 MPa in air at 950 ◦ C. Photo (a) is a low magnification of the whole gage section. While photo (b) is a high magnification showing the creep cavities formed inside the region. The creep cavities and cracks developed along the grain boundaries perpendicular to the applied stress. The cavities were widely distributed over the gage section. Creep failure finally occurred by the incorporation of these cavities (or voids) introduced from the creep damage. These cavities were initially generated by precipitates formed along the grain boundary. The precipitates become site initiating cavities. Fig. 9 shows the precipitates observed in the (a) grip section (stress free region) and (b) gage section (applied stress region) of the specimens crept in the He environment under 20 MPa (rupture time of 2020 h) at 950 ◦ C. The grip-section sample of the stress free region, taken from the grip end, precipitated through the carbides in the matrix and on the grain boundaries, as clearly shown by the carbides.
On the contrary, in the gage-section sample (near the failure region), which was applied by creep stress, the carbides were very scattered throughout matrix, and the grain boundaries were not clear. Also, the carbides were coarse and blocky, and the cavities occurred on the grain boundaries where the carbides were formed. Cylindrically-shaped coarse carbides, Crrich M23 C6 , developed mainly on the grain boundaries. In the precipitate-denuded regions, the carbides developed near the grain boundaries. According to other studies on Alloy 617, M23 C6 carbide was found to be abundant at temperatures of 649 to 1093 ◦ C. Other carbides (Cr23 C6 ), carbon nitrides [Cr Mo(C, N)], and nitride (TiN) were exhibited but rare (Kim et al., 2012a; Jang et al., 2008). After the creep failure, the fracture surface of each tested specimen, conducted under the high and low stress levels in the air and He environments at the three temperatures of 950, 900, and 850 ◦ C, was examined by scanning electron microscopy (SEM). Fig. 10 shows typical SEM micrographs taken at two magnifications showing fracture modes under 70 MPa and 35 MPa in the air and He environments at 850 ◦ C. The fracture modes are similar in both environments. The plane of the surfaces is not flat at the high stress level of 70 MPa. The fracture surface shows a dominant dimple rupture and secondary cracking with the indication of intergranular cracking. However, the fracture surface at the low stress level of 35 MPa shows a combination of dimple rupture and dominant intergranular fracture. The plane of the surfaces is very even when compared with the high stress of 70 MPa. In particular, surface oxidation is evident owing to a prolonged exposure time under 35 MPa. The intergranular cracking is initiated and developed by incorporation of minor cavities, formed on the grain boundary for prolonged creep duration.
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Fig. 10. Typical SEM micrographs of the fracture surface under the high and low stress levels of 70 and 35 MPa in air and He environments at 850 ◦ C.
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Fig. 11. Typical oxide layers formed for Alloy 617 (28 MPa in He environment at 900 ◦ C).
3.3. Oxidation features in air and He environments with temperatures Three microstructure features are investigated in this study related to oxidation behavior during creep exposure time of Alloy 617. They are the surface oxide-layer thickness, the depth of the internal oxidation and the depth of the carbide depleted zone (or Cr depleted zone). These three regions are labeled in the micrographs, as shown in Fig. 11. By quantifying the time dependence of the change in the microstructure features, the mechanism of degradation can be ascertained in the air and He environments at the three temperatures of 950, 900 and 850 ◦ C. An energy dispersive X-ray (EDX) analysis was used to identify the metallic element constituents at different locations. On the outer surface, a Cr-rich (mainly Cr2 O3 ) oxide layer was formed owing to Cr diffusion from the matrix during the creep exposure time. Just below the Cr2 O3 layer, an Al-rich (mainly Al2 O3 ) discrete internal oxide layer was formed. The internal oxidation was formed by the diffusion of oxygen through the surface oxide and matrix. Then, below the internal oxidation layer, a
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thick carbide-depleted zone (or Cr depleted zone) was formed by depletion of the carbides at the grain boundary near the surface oxides as the oxidation progressed gradually. The depletion of grain boundary carbides is caused by a reaction between the grain boundary carbides and oxygen provided by diffusion from the environment or the oxidation reaction of the surface oxide layers. Carbide dissolution occurs from chromium migration to support oxide formation on the surface, as Alloy 617 is a chromium-forming alloy. It is known that the optimal distribution of grain boundary carbide improves the creep resistance as the grain boundary carbide can prevent grain boundary migration and sliding (Christ et al., 1987). Fig. 12 shows comparisons of typical SEM micrographs of the oxide layers formed for creep failure specimens in the (a) air and (b) He environments under an identical stress of 20 MPa at 950 ◦ C. As shown in (a) and (b), different oxidation features were found between the two environments, as follows. First, for the surface oxide layer, the thickness in He was thicker than in air, although the helium gas used in this study was at high purity at 99.999% containing small amount of impurities, such as H2 O < 1.0 ppm, O2 < 1.0 ppm, N2 < 5.0 ppm. The structure was more porous and smooth than that in the air samples. The connection of the surface oxide layer in the He samples was broken by cracks. In the air samples, the connection is continuous without cracks or breaks, and the structures were very dense over the gage length. The dense oxide layer is formed owing to a high oxygen concentration in the atmosphere, and it provides a proper protection layer from surface oxidation. Then, for the internal oxide layer, the He samples were deeper in Al-oxide intergrowth, and the interconnection between the surface oxide layer and matrix was weak. The structure was also porous and bared. As the temperature and creep rupture time increased, the depth showed roughly an increasing trend. Finally, for the Cr-depleted layer, the depth of the He environment was not different compared to that of the air. As the temperature and creep rupture time increased, the depth was enhanced on the micrographs, although the depth was difficult to quantitatively measure. The amount of carbide precipitates at the grain
Fig. 12. Comparison of SEM micrographs of the oxide layers formed during creep rupture time in (a) air and (b) He environments under an identical applied stress of 20 MPa at 950 ◦ C.
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Fig. 13. Typical SEM image observed for longitudinal section of the specimens crept in air and He environments under 20 MPa at 950 ◦ C and 35 MPa at 850 ◦ C.
boundary was much smaller than that in the air, as shown in Fig. 13. Accordingly, it is thought that the differences of the above three microstructure features caused detrimental effects on the creep resistance, even in the low oxygen-containing He agent used in this study. In addition, in this examination, it was difficult to quantitatively measure the depths of the Al-oxide intergrowth and Cr-depleted zone with creep rupture time. Thus, it is focused on the surface oxide-layer thickness, and its thickness was measured by SEM taken for the crept specimens in both environments at 950, 900, and 850 ◦ C. Typical micrographs of longitudinal section are shown in Fig. 13. The samples were polished and etched for each crept specimen. The surface oxide-layer thickness was measured for more than 20 points with an interval of ∼80 m along the longitudinal direction of samples. The average value and standard deviation were obtained.
Fig. 14 shows the measured results of the surface oxide-layer thickness with the creep rupture time at (a) 950 ◦ C, (b) 900 ◦ C, and (c) 850 ◦ C in the air and He environments. The surface oxide-layer thickness increased as the rupture time increased for both environments at the three temperatures. The He samples revealed higher thickness values compared to the air samples as the enhanced oxide-layer thickness in the He environment at 950 and 900 ◦ C reduced the effective area of the specimens carrying creep load. The deterioration of the creep resistance was due to the microstructural changes such as the thickness of the surface oxide layer, the depth of the internal oxidation, and the depth of the Crdepleted zone. These changes were attributed to the enhanced diffusion of the oxidizing agent and the gaseous reaction products along the grain boundary. It is known that the creep properties of the material will have detrimental effects when the carbide depleted zone is enhanced, and they are dependent on various
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µ
µ
µ
µ
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µ
Fig. 15. Comparisons of surface oxide-layer thickness with creep rupture time in (a) air and (b) He environments at 950, 900, and 850 ◦ C.
Fig. 14. Comparisons of surface oxide-layer thickness with creep rupture time at (a) 950 ◦ C, (b) 900 ◦ C, and (c) 850 ◦ C in air and He environments.
microstructue features such as surface cracking along the grain boundary, voids below the surface, and voids in the matrix (Dewson and Li, 2005). At 950 and 900 ◦ C, the surface oxide-layer thickness of the He samples was thinner at short rupture time than that of air, but after the time passed for 1000 h at 950 ◦ C, and 3000 h at 900 ◦ C, the thickness increased when compared to in air. This result at the short
rupture time is due to the low oxygen-containing He environment used in this study. However, at 850 ◦ C, the thickness was as low as 10 m in spite of the long-term rupture reaching 12,000 h. The thickness for both environments was almost same. The thickness changes with time variations revealed a roughly parabolic trend, as shown in Fig. 15. It is believed that the parabolic trend was determined by the diffusion rate when the supply of oxygen atom penetrates a dense oxide layer formed for long-term creep. At 950 and 900 ◦ C, the surface oxide-layer thickness was almost similar in both environments, but at 850 ◦ C, the thickness was reduced by about 50% when compared to the temperatures of 950 and 900 ◦ C. The reason why the surface oxide layer formed much deeper inside in the He than in air can be explained as following. Oxidation is a diffusion controlled process. Once the surface of a metal begins to oxidize, subsequent oxidation can take place only when oxygen ions diffuse through the oxide and react with the metal ions at the metal-oxide interface. As the oxide layer thickens, the diffusion distance increases and the rate of oxidation decreases. The density of oxide formed on Alloy 617 by annealing in air is much more than that formed in He, which tends to make the oxygen ions diffuse more difficult, while the much less density of the oxide on the surface formed in the He ambience tends to make the oxygen ions diffuse easier and then to increase the diffusion distance of the oxygen ions. From the results, it is found that the results of the surface oxide thickness were in accordance with those of the creep properties, as shown in Figs. 4–6. In particular, at 850 ◦ C, the creep and oxidation behaviors were similar regardless of both environments. Therefore,
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850 ◦ C is assumed to be a boundary temperature at which the He effects disappeared. 4. Conclusions A series of creep tests was performed at 950, 900, and 850 ◦ C in air and He environments to investigate the temperature effect on creep and oxidation behaviors of Alloy 617. The creep properties were analyzed, and oxidation microstructural features were analyzed by observing the crept specimens tested in both environments. The surface oxide-layer thickness was measured to define oxidation behavior with the rupture time. Based on the creep data and oxidation microstructural analysis, the following conclusions were drawn. At 950 and 900 ◦ C, the creep rupture time in the He environment was shorter than that in air, and the thickness of the surface oxide layer was thicker than in air. A significant deterioration of the creep resistance in He environment was due to the thicker surface oxide-layer formed by the reaction of surface oxygen and Cr diffusion in the matrix, even in a low oxygen-containing He agent. The thicker surface oxide-layer reduced the effective area carrying the creep load, and the creep resistance was reduced. On the other hand, at 850 ◦ C, the rupture time was almost the same regardless of both environments. The surface oxide-layer thickness was reduced by about 50% compared to that of 950 and 900 ◦ C. 850 ◦ C could be assumed as a boundary temperature at which the He effects disappeared. The relationships between the creep properties and surface oxide-layer thickness were in accordance in both environments.
Acknowledgments This study was supported by Nuclear Research and Development Program of the National Research Foundation of Korea (NRF) grant funded by the Korean government (MSIP) (Grant code: NRF2012MZA8A2025682). References Christ, H., Kunecke, U., Meyer, K., 1987. High temperature corrosion of the nickelbased alloy Inconel 617 in helium containing small amounts of impurities. Mater. Sci. Eng. 87, 161. Cook, R., 1984. Creep properties of Inconel-617 in air and helium at 800 to 1000 ◦ C. Nucl. Technol. 66, 283. Dewson, S., Li, X., 2005. Selection criteria for the high temperature reactor intermediate heat exchanger. In: Proceedings of ICAPP 05, Paper No. 5333, Seoul. Jang, C., Lee, D., Kim, D., 2008. Oxidation behavior of an alloy 617 in very high temperature air and helium environments. Int. J. Press. Vessels Piping 85, 368. Kim, W.G., Yin, S.N., Ryu, W.S., Chang, J.H., 2007. Analysis of the creep rupture data of alloy 617 for a high temperature gas cooled reactor. In: Proceedings of CREEP8, ASME PVP 2007-26834, Texas. Kim, W.G., Yin, S.N., Kim, Y.W., Chang, J.H., 2008. Creep characterization of a Nibased Hastelloy-X alloy by using theta projection method. Eng. Fract. Mech. 75, 4985. Kim, W.G., Yin, S.N., Kim, Y.W., 2009. Long-term creep strain-time curve modeling of alloy 617 for a VHTR intermediate heat exchanger. J. Korean Inst. Met. Mater. 47 (10), 613 (in Korean). Kim, W.G., Yin, S.N., Lee, G.G., Kim, Y.W., Kim, S.J., 2010. Creep oxidation behavior and creep strength prediction for alloy 617. Int. J. Press. Vessels Piping 87, 289. Kim, W.G., Yin, S.N., Park, J.Y., Hong, S.D., Kim, Y.W., 2012a. An improved methodology for determining tensile design strengths of alloy 617. J. Mech. Sci. Technol. 26 (2), 379. Kim, W.G., Park, J.Y., Kim, S.J., Hong, S.D., Kim, Y.W., 2012b. Reliability evaluation on creep life prediction of alloy 617 for a very high temperature reactor. J. Korean Inst. Met. Mater. 50 (10), 721 (in Korean).
Contents lists available at ScienceDirect
Nuclear Engineering and Design journal homepage: www.elsevier.com/locate/nucengdes
Temperature effect on the creep behavior of alloy 617 in air and helium environments Woo-Gon Kim a,∗ , Jae-Young Park b , Gyeong-Geun Lee a , Sung-Deok Hong a , Yong-Wan Kim a a b
Korea Atomic Energy Research Institute, Daejeon 305-353, Republic of Korea Pukyong National University, Busan 608-739, Republic of Korea
a b s t r a c t The temperature effect on creep and oxidation behaviors in air and helium (He) environments was investigated at 950, 900 and 850 ◦ C for Alloy 617, which is considered as a prime candidate material for VHTR components. Creep data were obtained with different stress levels at the three temperatures. Oxidation microstructural features such as the surface oxide layer, internal oxidation and decarburization were analyzed by observing each crept specimen. At 950 and 900 ◦ C, the creep rupture time in He environment was shorter than that in air, and the thickness of the surface oxide layer was thicker than in air. The deterioration of creep resistance in the He environment was due to a thicker oxide-layer thickness, which reduced the effective area carrying the creep load. On the other hand, at 850 ◦ C, the creep rupture time was almost the same regardless of both environments. The surface oxide-layer thickness was reduced by about 50% compared to the temperatures of 950 and 900 ◦ C. The temperature of 850 ◦ C could be assumed as a boundary temperature at which the He effects disappeared. It was found that the relationship between the creep data and surface oxide-layer thickness was in accordance with both environments. © 2013 Elsevier B.V. All rights reserved.
1. Introduction A very high temperature gas reactor (VHTR) is one of the Gen-IV reactors aiming at safe, long-lived, proliferation-resistant and economical nuclear power plants. Its high operating temperature of over 800 ◦ C enables a high energy efficiency and the production of hydrogen gas using Sulfur–Iodine process. The heat of the primary helium (He) circuit transfers to the secondary helium loop through the intermediate heat exchanger (IHX). The IHX component needs high-temperature creep resistance in the He environment, and requires good oxidation resistance, corrosion resistance, and phase stability at high temperatures. Among them, creep and oxidation properties are importantly considered, as the integrity of the components should be preserved during a prolonged period in a VHTR coolant (Kim et al., 2007, 2008, 2009, 2010, 2012a,b). Of the existing alloys, nickel-base Alloy 617 is the leading candidate for use in the next generation nuclear plant (NGNP) heat exchangers because it has the highest creep strength of solid solution alloys under consideration for temperature above 850 ◦ C. It was reported that the creep rupture time varied widely with the environments interactions between He impurities and the alloy (Jang et al., 2008; Dewson and Li, 2005; Cook, 1984). The impurities in He
∗ Corresponding author at: Korea Atomic Energy Research Institute, 989-111 Daedeok-daero, Yuseong-gu, Daejeon 305-353, Republic of Korea. Tel.: +82 42 868 2493. E-mail address: [email protected] (W.-G. Kim). 0029-5493/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.nucengdes.2013.11.050
coolant during steady-state operation of the VHTRs were reported as H2 O, H2 , CO, CO2 , CH4 , O2 and N2 . However, it has not been well established on the temperature effect in He environment to be drastically affected to the rupture time. The relationships between creep and oxidation features are not well understood. Therefore, it is necessary to investigate the temperature effect on the creep and oxidation behaviors in air and He environments of Alloy 617. In this study, the creep behaviors of Alloy 617 were comparatively investigated by creep tests with different stresses in air and He environments at 950, 900, and 850 ◦ C. The temperature effect on the creep and oxidation behaviors was evaluated, and relationships between these behaviors were found and discussed. 2. Experimental procedures Commercial grade nickel-based superalloy, Alloy 617 (Inconel 617), was used in this study. The material was a hot-rolled plate with a thickness of 15.875 mm. Creep test specimens in the air and He environments were fabricated in cylindrical form with a 30 mm gauge length and 6 mm diameter. The gage section was parallel to the longitudinal rolling direction. Circular grooves were machined at both ends beyond the shoulder region of these specimens to attach extensometers for monitoring the elongation during creep testing. The loading frames used in the creep tests had a lever arm ratio of 20:1. The overall creep rates were determined from elongations measured by extensometers equipped at the circular grooves. Creep strain data with elapsed times were taken automatically by
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3. Results and discussion 3.1. Comparison of creep behaviors in air and He environments at 950, 900 and 850 ◦ C The creep property data of Alloy 617 such as the rupture time, minimum creep rate, rupture elongation, and reduction of area were obtained in the air and He environments at 950, 900, and 850 ◦ C. The creep results in both environments were compared at the three temperatures.
Fig. 1. Creep testing apparatus in He environment.
Fig. 2. Typical creep curves obtained in air and He environments at 900 ◦ C.
a PC. K-type thermocouples were used to monitor temperatures within the gage section of the specimens. A split furnace was used to heat the specimens of Alloy 617 in the air and He environments to 950, 900, and 850 ◦ C. The pull rod and jig used for the creep tests were manufactured with a nickel-based superalloy material to sufficiently endure oxidation and thermal degradation during creep. Before creep tests in the He environment, a vacuum chamber made for the quartz tube was purged three or four times by a vacuum pump to remove some impurities in the chamber. During the creep tests, pure He gas of 99.999% was supplied to the specimens equipped in the quartz tube. Impurity concentration in pure He gas was O2 < 1.0 ppm, N2 < 5.0 ppm and H2 O < 1.0 ppm. Flow rate of the He gas during the creep test was controlled under 20 cm3 /min. The control of the other impurities such as CO2 , CO, CH4 , and H2 in He gas simulating a VHTR condition will be investigated in the further tests. Creep testing apparatus in the He environment is shown in Fig. 1. Photos showing each component such as the quart tube, He flow meter, strain indicator, specimen and jigs, furnace controller, and data acquisition PC are presented.
Fig. 3. A photo of crept specimen under 18 MPa at 950 ◦ C in air of Alloy 617.
Fig. 4. Comparison of stress vs. rupture time in the air and He environments at (a) 950 ◦ C, (b) 900 ◦ C, and (c) 850 ◦ C.
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Fig. 5. Comparison of minimum creep rate vs. stress in air and He environments at 950, 900, and 850 ◦ C. Fig. 6. Comparison of rupture time vs. minimum creep rate in air and He environments at 950, 900, and 850 ◦ C.
Fig. 2 shows typical creep curves for stresses in the air and He environments at 900 ◦ C. Alloy 617 showed little primary creep strain, and a well-defined secondary creep stage was not exhibited in the full creep curves. The onset of a tertiary creep was unclear, and a tertiary creep stage was initiated from a low strain level. There were no differences in the shapes of the creep curves between the air and He environments. Alloy 617 revealed good ductility (>30%) in spite of the long duration above 104 h at 950, 900, and 850 ◦ C.
From a series of creep tests, long-term creep data reaching 14,000 h at 950 ◦ C were achieved in the air, as shown in Fig. 3. Creep rupture occurred at three positions within the gage length. The rupture elongation was about 30%, and the reduction of area was as small as 22%. It was identified that the creep rupture was developed by the incorporation of the many small cavities formed during the creep, rather than the necking of the specimen.
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Fig. 8. Typical OM micrographs of a crept specimen under 28 MPa in air at 950 ◦ C.
Fig. 7. Comparison of creep rupture elongation vs. stress in air and He environments at 950, 900, and 850 ◦ C.
Fig. 4 shows a comparison of the log stress vs. log rupture time in the air and He environments at 950, 900 and 850 ◦ C. For 950 and 900 ◦ C, in the high stress ranges above about 30 MPa, the creep stress between the air and He environments was almost similar. However, in the low stress ranges below about 30 MPa, there were differences in the creep stress (or creep rupture time)
between the two environments. Namely, the creep stress in air was higher than that in He environment, but for 850 ◦ C, the creep data of the log–log plot was almost superimposed regardless of both environments, as shown well in the figure. Therefore, it is obvious that the creep properties in air and He environments have a temperature effect. The reason for this is closely related to the oxide layer thickness formed during the creep. The oxide-layer thicknesses in both environments are measured at the three temperatures, and the results are discussed in a later section. Fig. 5 shows a comparison of the minimum creep rate vs. stress at 950, 900, and 850 ◦ C in air and He environments. The relationships between the minimum creep rate and stress revealed a good linearity regardless of both environments at the three temperatures. It is clear that Alloy 617 followed Norton’s power law well under this creep condition because the deformation map corresponded to the power-law creep region. The creep mechanism is governed by a climb of dislocation (Kim et al., 2010). The creep rates between both environments at 950 and 900 ◦ C plot a significant difference at low stress levels. The creep rate in He was faster than that in air. However, the creep rates at 850 ◦ C were almost the same regardless of both environments. Fig. 6 shows a plot of Monkman–Grant (M–G) relationships. In the equation of M–G relationship, log tr + mε˙ m = C, the m values of the slope were obtained in the air and He environments at the three temperatures of 950, 900, and 850 ◦ C. At 950 and 900 ◦ C, it means that the creep rupture time of the He specimens was shorter than that of the air one. But at 850 ◦ C, their values were almost equal to
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Fig. 9. Precipitates formed in the grip-section region and gage-section region of the specimen crept in He environment under 20 MPa at 950 ◦ C.
m = 1.0 regardless of the environment. This means that the creep properties were not changed at 850 ◦ C under the He environment. It is obvious that the He effects on the creep properties disappeared at 850 ◦ C. This is assumed to be a temperature effect on the creep properties under this He-agent condition. Thus, the temperature of 850 ◦ C can be regarded as a boundary temperature at which the He effects disappeared. The data on the creep rupture elongation were similar in both environments, although the data were scattered, as shown in Fig. 7.
3.2. Creep fracture micrographs Fig. 8 shows typical OM micrographs at two magnifications of the crept specimen under 28 MPa in air at 950 ◦ C. Photo (a) is a low magnification of the whole gage section. While photo (b) is a high magnification showing the creep cavities formed inside the region. The creep cavities and cracks developed along the grain boundaries perpendicular to the applied stress. The cavities were widely distributed over the gage section. Creep failure finally occurred by the incorporation of these cavities (or voids) introduced from the creep damage. These cavities were initially generated by precipitates formed along the grain boundary. The precipitates become site initiating cavities. Fig. 9 shows the precipitates observed in the (a) grip section (stress free region) and (b) gage section (applied stress region) of the specimens crept in the He environment under 20 MPa (rupture time of 2020 h) at 950 ◦ C. The grip-section sample of the stress free region, taken from the grip end, precipitated through the carbides in the matrix and on the grain boundaries, as clearly shown by the carbides.
On the contrary, in the gage-section sample (near the failure region), which was applied by creep stress, the carbides were very scattered throughout matrix, and the grain boundaries were not clear. Also, the carbides were coarse and blocky, and the cavities occurred on the grain boundaries where the carbides were formed. Cylindrically-shaped coarse carbides, Crrich M23 C6 , developed mainly on the grain boundaries. In the precipitate-denuded regions, the carbides developed near the grain boundaries. According to other studies on Alloy 617, M23 C6 carbide was found to be abundant at temperatures of 649 to 1093 ◦ C. Other carbides (Cr23 C6 ), carbon nitrides [Cr Mo(C, N)], and nitride (TiN) were exhibited but rare (Kim et al., 2012a; Jang et al., 2008). After the creep failure, the fracture surface of each tested specimen, conducted under the high and low stress levels in the air and He environments at the three temperatures of 950, 900, and 850 ◦ C, was examined by scanning electron microscopy (SEM). Fig. 10 shows typical SEM micrographs taken at two magnifications showing fracture modes under 70 MPa and 35 MPa in the air and He environments at 850 ◦ C. The fracture modes are similar in both environments. The plane of the surfaces is not flat at the high stress level of 70 MPa. The fracture surface shows a dominant dimple rupture and secondary cracking with the indication of intergranular cracking. However, the fracture surface at the low stress level of 35 MPa shows a combination of dimple rupture and dominant intergranular fracture. The plane of the surfaces is very even when compared with the high stress of 70 MPa. In particular, surface oxidation is evident owing to a prolonged exposure time under 35 MPa. The intergranular cracking is initiated and developed by incorporation of minor cavities, formed on the grain boundary for prolonged creep duration.
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Fig. 10. Typical SEM micrographs of the fracture surface under the high and low stress levels of 70 and 35 MPa in air and He environments at 850 ◦ C.
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Fig. 11. Typical oxide layers formed for Alloy 617 (28 MPa in He environment at 900 ◦ C).
3.3. Oxidation features in air and He environments with temperatures Three microstructure features are investigated in this study related to oxidation behavior during creep exposure time of Alloy 617. They are the surface oxide-layer thickness, the depth of the internal oxidation and the depth of the carbide depleted zone (or Cr depleted zone). These three regions are labeled in the micrographs, as shown in Fig. 11. By quantifying the time dependence of the change in the microstructure features, the mechanism of degradation can be ascertained in the air and He environments at the three temperatures of 950, 900 and 850 ◦ C. An energy dispersive X-ray (EDX) analysis was used to identify the metallic element constituents at different locations. On the outer surface, a Cr-rich (mainly Cr2 O3 ) oxide layer was formed owing to Cr diffusion from the matrix during the creep exposure time. Just below the Cr2 O3 layer, an Al-rich (mainly Al2 O3 ) discrete internal oxide layer was formed. The internal oxidation was formed by the diffusion of oxygen through the surface oxide and matrix. Then, below the internal oxidation layer, a
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thick carbide-depleted zone (or Cr depleted zone) was formed by depletion of the carbides at the grain boundary near the surface oxides as the oxidation progressed gradually. The depletion of grain boundary carbides is caused by a reaction between the grain boundary carbides and oxygen provided by diffusion from the environment or the oxidation reaction of the surface oxide layers. Carbide dissolution occurs from chromium migration to support oxide formation on the surface, as Alloy 617 is a chromium-forming alloy. It is known that the optimal distribution of grain boundary carbide improves the creep resistance as the grain boundary carbide can prevent grain boundary migration and sliding (Christ et al., 1987). Fig. 12 shows comparisons of typical SEM micrographs of the oxide layers formed for creep failure specimens in the (a) air and (b) He environments under an identical stress of 20 MPa at 950 ◦ C. As shown in (a) and (b), different oxidation features were found between the two environments, as follows. First, for the surface oxide layer, the thickness in He was thicker than in air, although the helium gas used in this study was at high purity at 99.999% containing small amount of impurities, such as H2 O < 1.0 ppm, O2 < 1.0 ppm, N2 < 5.0 ppm. The structure was more porous and smooth than that in the air samples. The connection of the surface oxide layer in the He samples was broken by cracks. In the air samples, the connection is continuous without cracks or breaks, and the structures were very dense over the gage length. The dense oxide layer is formed owing to a high oxygen concentration in the atmosphere, and it provides a proper protection layer from surface oxidation. Then, for the internal oxide layer, the He samples were deeper in Al-oxide intergrowth, and the interconnection between the surface oxide layer and matrix was weak. The structure was also porous and bared. As the temperature and creep rupture time increased, the depth showed roughly an increasing trend. Finally, for the Cr-depleted layer, the depth of the He environment was not different compared to that of the air. As the temperature and creep rupture time increased, the depth was enhanced on the micrographs, although the depth was difficult to quantitatively measure. The amount of carbide precipitates at the grain
Fig. 12. Comparison of SEM micrographs of the oxide layers formed during creep rupture time in (a) air and (b) He environments under an identical applied stress of 20 MPa at 950 ◦ C.
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Fig. 13. Typical SEM image observed for longitudinal section of the specimens crept in air and He environments under 20 MPa at 950 ◦ C and 35 MPa at 850 ◦ C.
boundary was much smaller than that in the air, as shown in Fig. 13. Accordingly, it is thought that the differences of the above three microstructure features caused detrimental effects on the creep resistance, even in the low oxygen-containing He agent used in this study. In addition, in this examination, it was difficult to quantitatively measure the depths of the Al-oxide intergrowth and Cr-depleted zone with creep rupture time. Thus, it is focused on the surface oxide-layer thickness, and its thickness was measured by SEM taken for the crept specimens in both environments at 950, 900, and 850 ◦ C. Typical micrographs of longitudinal section are shown in Fig. 13. The samples were polished and etched for each crept specimen. The surface oxide-layer thickness was measured for more than 20 points with an interval of ∼80 m along the longitudinal direction of samples. The average value and standard deviation were obtained.
Fig. 14 shows the measured results of the surface oxide-layer thickness with the creep rupture time at (a) 950 ◦ C, (b) 900 ◦ C, and (c) 850 ◦ C in the air and He environments. The surface oxide-layer thickness increased as the rupture time increased for both environments at the three temperatures. The He samples revealed higher thickness values compared to the air samples as the enhanced oxide-layer thickness in the He environment at 950 and 900 ◦ C reduced the effective area of the specimens carrying creep load. The deterioration of the creep resistance was due to the microstructural changes such as the thickness of the surface oxide layer, the depth of the internal oxidation, and the depth of the Crdepleted zone. These changes were attributed to the enhanced diffusion of the oxidizing agent and the gaseous reaction products along the grain boundary. It is known that the creep properties of the material will have detrimental effects when the carbide depleted zone is enhanced, and they are dependent on various
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Fig. 15. Comparisons of surface oxide-layer thickness with creep rupture time in (a) air and (b) He environments at 950, 900, and 850 ◦ C.
Fig. 14. Comparisons of surface oxide-layer thickness with creep rupture time at (a) 950 ◦ C, (b) 900 ◦ C, and (c) 850 ◦ C in air and He environments.
microstructue features such as surface cracking along the grain boundary, voids below the surface, and voids in the matrix (Dewson and Li, 2005). At 950 and 900 ◦ C, the surface oxide-layer thickness of the He samples was thinner at short rupture time than that of air, but after the time passed for 1000 h at 950 ◦ C, and 3000 h at 900 ◦ C, the thickness increased when compared to in air. This result at the short
rupture time is due to the low oxygen-containing He environment used in this study. However, at 850 ◦ C, the thickness was as low as 10 m in spite of the long-term rupture reaching 12,000 h. The thickness for both environments was almost same. The thickness changes with time variations revealed a roughly parabolic trend, as shown in Fig. 15. It is believed that the parabolic trend was determined by the diffusion rate when the supply of oxygen atom penetrates a dense oxide layer formed for long-term creep. At 950 and 900 ◦ C, the surface oxide-layer thickness was almost similar in both environments, but at 850 ◦ C, the thickness was reduced by about 50% when compared to the temperatures of 950 and 900 ◦ C. The reason why the surface oxide layer formed much deeper inside in the He than in air can be explained as following. Oxidation is a diffusion controlled process. Once the surface of a metal begins to oxidize, subsequent oxidation can take place only when oxygen ions diffuse through the oxide and react with the metal ions at the metal-oxide interface. As the oxide layer thickens, the diffusion distance increases and the rate of oxidation decreases. The density of oxide formed on Alloy 617 by annealing in air is much more than that formed in He, which tends to make the oxygen ions diffuse more difficult, while the much less density of the oxide on the surface formed in the He ambience tends to make the oxygen ions diffuse easier and then to increase the diffusion distance of the oxygen ions. From the results, it is found that the results of the surface oxide thickness were in accordance with those of the creep properties, as shown in Figs. 4–6. In particular, at 850 ◦ C, the creep and oxidation behaviors were similar regardless of both environments. Therefore,
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850 ◦ C is assumed to be a boundary temperature at which the He effects disappeared. 4. Conclusions A series of creep tests was performed at 950, 900, and 850 ◦ C in air and He environments to investigate the temperature effect on creep and oxidation behaviors of Alloy 617. The creep properties were analyzed, and oxidation microstructural features were analyzed by observing the crept specimens tested in both environments. The surface oxide-layer thickness was measured to define oxidation behavior with the rupture time. Based on the creep data and oxidation microstructural analysis, the following conclusions were drawn. At 950 and 900 ◦ C, the creep rupture time in the He environment was shorter than that in air, and the thickness of the surface oxide layer was thicker than in air. A significant deterioration of the creep resistance in He environment was due to the thicker surface oxide-layer formed by the reaction of surface oxygen and Cr diffusion in the matrix, even in a low oxygen-containing He agent. The thicker surface oxide-layer reduced the effective area carrying the creep load, and the creep resistance was reduced. On the other hand, at 850 ◦ C, the rupture time was almost the same regardless of both environments. The surface oxide-layer thickness was reduced by about 50% compared to that of 950 and 900 ◦ C. 850 ◦ C could be assumed as a boundary temperature at which the He effects disappeared. The relationships between the creep properties and surface oxide-layer thickness were in accordance in both environments.
Acknowledgments This study was supported by Nuclear Research and Development Program of the National Research Foundation of Korea (NRF) grant funded by the Korean government (MSIP) (Grant code: NRF2012MZA8A2025682). References Christ, H., Kunecke, U., Meyer, K., 1987. High temperature corrosion of the nickelbased alloy Inconel 617 in helium containing small amounts of impurities. Mater. Sci. Eng. 87, 161. Cook, R., 1984. Creep properties of Inconel-617 in air and helium at 800 to 1000 ◦ C. Nucl. Technol. 66, 283. Dewson, S., Li, X., 2005. Selection criteria for the high temperature reactor intermediate heat exchanger. In: Proceedings of ICAPP 05, Paper No. 5333, Seoul. Jang, C., Lee, D., Kim, D., 2008. Oxidation behavior of an alloy 617 in very high temperature air and helium environments. Int. J. Press. Vessels Piping 85, 368. Kim, W.G., Yin, S.N., Ryu, W.S., Chang, J.H., 2007. Analysis of the creep rupture data of alloy 617 for a high temperature gas cooled reactor. In: Proceedings of CREEP8, ASME PVP 2007-26834, Texas. Kim, W.G., Yin, S.N., Kim, Y.W., Chang, J.H., 2008. Creep characterization of a Nibased Hastelloy-X alloy by using theta projection method. Eng. Fract. Mech. 75, 4985. Kim, W.G., Yin, S.N., Kim, Y.W., 2009. Long-term creep strain-time curve modeling of alloy 617 for a VHTR intermediate heat exchanger. J. Korean Inst. Met. Mater. 47 (10), 613 (in Korean). Kim, W.G., Yin, S.N., Lee, G.G., Kim, Y.W., Kim, S.J., 2010. Creep oxidation behavior and creep strength prediction for alloy 617. Int. J. Press. Vessels Piping 87, 289. Kim, W.G., Yin, S.N., Park, J.Y., Hong, S.D., Kim, Y.W., 2012a. An improved methodology for determining tensile design strengths of alloy 617. J. Mech. Sci. Technol. 26 (2), 379. Kim, W.G., Park, J.Y., Kim, S.J., Hong, S.D., Kim, Y.W., 2012b. Reliability evaluation on creep life prediction of alloy 617 for a very high temperature reactor. J. Korean Inst. Met. Mater. 50 (10), 721 (in Korean).