Microstructural investigation of Alloy 617 corroded in high-temperature helium environment

Nuclear Engineering and Design 271 (2014) 301–308

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Microstructural investigation of Alloy 617 corroded in high-temperature helium environment Gyeong-Geun Lee ∗ , Sujin Jung, Daejong Kim, Yong-Whan Jeong, Dong-Jin Kim Korea Atomic Energy Research Institute, Daejeon 305-353, Republic of Korea

a b s t r a c t The high-temperature corrosion behavior of Alloy 617, which is a candidate material for the intermediate heat exchanger (IHX) in a very high-temperature gas-cooled reactor (VHTR), was investigated. The corrosion tests were carried out at 850–950 ◦ C in impure helium environments. All specimens showed a parabolic oxidation behavior at up to 250 h, and the activation energy for the rate constants of the mass change was similar to the activation energy for Cr/Cr2 O3 oxidation in air. SEM micrographs revealed the evolution of the outer oxide layer, internal oxide, and depleted zone of the grain boundary carbides. Corrosion test in a pure helium environment showed that the minor impurity gases in helium environment affected the oxidation kinetics of Alloy 617. It is suggested that the control of minor impurities in a VHTR helium environment is necessary to apply Alloy 617 to an IHX material in a VHTR. Crown Copyright © 2013 Published by Elsevier B.V. All rights reserved.

1. Introduction A Very High Temperature Gas Reactor (VHTR) is one of the GenIV nuclear reactors that have enhanced the economics, stability, non-proliferation, and long-term operation of nuclear reactors as compared to previous generation reactors. Due to a high operating temperature of up to 950 ◦ C, a VHTR has a high thermal efficiency and can supply process heat to the manufacturing applications (Elder and Allen, 2009). The coolant of a VHTR is inert helium gas with high thermal conductivity. The high-temperature nuclear heat in the primary helium loop is transferred to the secondary helium loop through an intermediate heat exchanger (IHX). IHX should endure a high temperature of over 850 ◦ C and a gas pressure difference of up to 8 MPa during operation, and hence IHX materials require high phase stability, superior creep resistance, and better creep-fatigue properties. The Ni-base superalloy is the best candidate material for an IHX component. Wrought Ni-base superalloys such as Alloy 617, Hastelloy XR, and Haynes 230 have been considered for use in an IHX because of their formability and weldability. The United States Department of Energy (US DOE) and the Republic of Korea are considering Alloy 617 as a primary candidate material for their VHTRs (Ren and Swimdeman, 2009; Lee et al., 2009) because of its exceptional creep strength and extensive properties. Alloy 617 is a solid-solution strengthened Ni–Cr–Co–Mo alloy with a hightemperature strength and oxidation resistance (Special Metals Publication, 2005). Mo and Co are effective for solid-solution

∗ Corresponding author. Tel.: +82 42 868 4688. E-mail address: [email protected] (G.-G. Lee).

strengthening, and Al in conjunction with Cr provides oxidation resistance at high temperatures. Alloy 617 appears to have superior properties in air; however, many researchers have reported that the mechanical properties of Alloy 617 at high temperatures above 900 ◦ C decreased in the impure helium environment of a VHTR (Hosoi and Abe, 1975; Cook, 1984; Quadakkers, 1984; Christ et al., 1987; Shankar and Natesan, 2007; Totemeier and Tian, 2007; Rouillard et al., 2007; Cabet and Duprey, 2010). It is generally accepted that minor impurities such as H2 , CO, and CH4 in helium cause oxidation, decarburization, and carburization variously. A lot of researches in Korea have focused on the high-temperature corrosion behavior and mechanical property changes in a helium environment (Jo et al., 2008; Jang et al., 2008; Kim et al., 2009, 2010, 2011; Lee et al., 2011). However, there remains a need for more accurate experiments and a systematic observation of the corrosion mechanism in impure helium environments. In this study, we investigated the corrosion behavior of Alloy 617 at high-temperatures in an impure helium environment. The mass changes of the specimens were measured, and the microstructures were analyzed quantitatively. In addition, a corrosion test in a pure helium environment was conducted, and the result was compared with the impure helium results.

2. Experimental procedures The material tested was a commercial grade Alloy 617 plate from Special Metals (Huntington, West Virginia, US). The chemical composition of the plate is shown in Table 1. The plate was cut into small coupons of ∼10 mm × 10 mm × 1 mm in size using a wire-cutting technique. No thermal treatment was carried out before machining.

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Table 1 Chemical composition of Alloy 617 used in the study (wt%).

(a)

1.2 o

He 850 C o He 900 C o He 950 C

1.0

Mn

Fe

S

Si

Cu

Ni

0.08

0.11

1.49

0.001

0.06

0.08

53.16

Cr

Al

Ti

Co

Mo

P

B

22.16

1.12

0.35

11.58

9.80

0.08

0.002

0.8 0.6 0.4 0.2 0.0 0

50

100

150

200

250

300

Time, h

(b)

1.2

4 2

0.8

2

o

1.0

0.6

(Mass change) , mg /cm

The surface of specimen was polished mechanically using 1500 grit paper. The apparatus was mainly composed of a furnace component, gas supply, and an analysis component. A pre-heater and a reaction furnace were incorporated into the furnace component. The preheater contained a graphite rod to reduce the partial pressure of oxygen in impure helium gas. The reaction furnace was a tubetype furnace with a vacuum system. Each specimen was located in a small quartz tube with a diameter of 1.2 cm and a length of 10 cm, and 3–5 tubes were inserted into the reaction furnace to ensure the independent reaction volume from the flowing helium gas. The stainless steel tubes were used to reduce gas contamination in the system. The concentration of impure helium gas was controlled by changing the mixing ratio of impurity/helium mixture gases (H2 –He, CO–He, CH4 –He) using mass flow controllers (MFC). The amount of impure gases H2 , CO, CO2 , CH4 , and N2 were measured by gas chromatography (HP 7890A, Agilent Technologies, USA). The typical time interval between measuring the gas during a corrosion test was 7–15 min. The humidity in the helium gas was measured using a dew point meter (Shaw Moisture Meters, England). The test temperatures were 850, 900, and 950 ◦ C, and the temperature of the pre-heater was fixed at 900 ◦ C. The duration of the holding period at a test temperature was up to 250 h. The impure helium composition was fixed as 200 ppm H2 , 50 ppm CO, 20 ppm CH4 , and H2 O < 2 ppm, and the flow rate of the gas was 50 cm3 /min. This is widely accepted condition in the high temperature corrosion test in helium environment (Quadakkers, 1984; Christ et al., 1987; Cabet and Duprey, 2010). The specimens were heated at a rate of 10 ◦ C/min, and cooled in the furnace after a corrosion test. The mass change of the specimen was measured using a precision balance with 1 × 10−5 g accuracy. A SEM (JEOL JSM-6300, Japan) with EDS (Energy-dispersive X-ray spectroscopy) was used to observe the microstructures and analyze the composition of the specimens. An electron back-scatter diffraction (EBSD) analysis was performed on the JEOL JSM-7000F.

Mass change, mg/cm

2

C

He 850 C o He 900 C o He 950 C

y=1.10e-6x-0.092

y=5.84e-7x-0.020

0.4 0.2 y=2.86e-7x-0.015

0.0 0

2x10

5

4x10

5

5

5

6x10

6

8x10

1x10

Time, sec Fig. 1. Mass change in the impure helium environment at various temperatures: (a) mass change vs. time (hour) and (b) (mass change)2 vs. time (sec). Table 2 Parabolic oxidation rate constants of Alloy 617 in impure helium environment. Condition 850 ◦ C 2

−4

kp (mg cm

−1

s

900 ◦ C −7

2.86 × 10

)

950 ◦ C −7

1.10 × 10−6

5.84 × 10

m, mass increase in a specimen; S, surface area of a specimen; kp , oxidation rate constant; and c, constant. The oxidation rate constants calculated from the mass change of the specimens are listed in Table 2. The oxidation rate constant

3. Results and discussion o

Temperature, C

3.1. Mass changes in impure helium 950

 m 2 S

= kp · t + c

(1)

900

850

-13.0

-6

2

4

-6

1.0x10 -14.0

Kp, mg cm sec

-6

1.5x10

y = -18558x + 1.458 2 R = 0.999

-13.5

-1

2.0x10

ln(Kp)

Fig. 1 shows the mass change per unit area as a function of exposure time at 850, 900, and 950 ◦ C in an impure helium environment. As shown in Fig. 1(a), the mass change of the specimens increased as the exposure time and temperature increased, and the rate of mass changes decreased with time. Fig. 1(b) plots the square of the mass change versus time in seconds, and linear relationships are clearly shown, and this relationship is known as a parabolic oxidation of metals (Birks et al., 2006). Parabolic oxidation occurs when the metal or oxygen species diffuse through the growing oxidation layer. As the thickness of the oxidation layer increases, the diffusion path of the species also increases and the diffusion rate decreases gradually with decreasing time. The mass increase during parabolic oxidation can be represented by the following equation,

-7

-14.5

5.0x10

-15.0 -4

8.00x10

-4

8.25x10

-4

8.50x10

-4

8.75x10

-4

9.00x10

1/Temperature, 1/K Fig. 2. Calculation of the activation energy for the oxidation of Alloy 617 in impure helium environment.

G.-G. Lee et al. / Nuclear Engineering and Design 271 (2014) 301–308

303

Fig. 4. EDS analysis results of precipitates in as-received Alloy 617.

Fig. 3. SEM micrographs of as-received Alloy 617.

at 950 ◦ C was 1.10 × 10−6 mg2 cm−4 s−1 , which is quite larger than the rate constant reported from Cabet, ∼0.28 × 10−6 mg2 cm−4 s−1 (Cabet and Duprey, 2010). However, the mass change at 950 ◦ C for 250 h was 0.95 mg/cm2 , which is approximately twice the result from Cabet. It is supposed that the difference was originated from the impurity level in helium gas and the effective contact area in the reaction furnace (the number and dimension of specimens). In

addition, it appears that the existence of constant term c affected the oxidation rate constants because of the short-term exposure time of these tests. Hence, long-term tests of over 1000 h are necessary to determine the precise oxidation rate constant of the corrosion tests. The oxidation behavior by diffusion is very sensitive to temperature, and this tendency can be expressed using Arrhenius equation.

 E  a

kp = A exp −

RT

(2)

A, non-temperature sensitive constant; Ea , activation energy for oxidation; R, gas constant; and T, temperature in Kelvin. Fig. 2 shows the relation between the inverse of temperature versus log kp from Table 2. The calculated Ea from our results was

Fig. 5. Microstructures of Alloy 617 corroded at 950 ◦ C in impure helium environment for (a) 25 h, (b) 50 h, (c) 100 h, and (d) 250 h.

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154 kJ/mol. Smeltzer et al. reported that the activation energy of the Cr/Cr2 O3 oxidation in air was 157 kJ/mol, and the related oxidation mechanism was grain boundary diffusion at similar temperatures (Smeltzer and Young, 1975). This is not completely clear because oxidation and carbide-depletion might occur simultaneously during corrosion of the specimens. However, it seems that the mass change during corrosion test is mainly caused by the oxidation of Cr in Alloy 617.

3.2. Microstructure observation The SEM micrographs obtained from the planes of an asreceived specimen are given in Fig. 3. The average grain size was above 100 ␮m, which was similar to the grain size from the information of the supplier (Special Metals Publication, 2005). This grain size was reported to provide the best creep-rupture strength and bend ductility at room temperature. There were bands of precipitates aligned in the rolling direction in the matrix, and thin precipitates were located along the grain boundaries. Fig. 4 shows a back-scattered SEM (BS SEM) image of the asreceived specimen, and it was found that there were two types of precipitates, bright and dark. EDS revealed that the bright precipitates were heavy Mo-rich M6 C, and the dark precipitates were light Cr-rich M23 C6 (Fig. 4(b)). This result is consistent with the literature (Totemeier and Tian, 2007; Jang et al., 2008). Fig. 5 presents the cross-sectional micrographs of the specimens corroded at 950 ◦ C. Because the exposure time was rather short, it

Fig. 6. BS SEM image and EDS results of the specimen tested for 250 h at 950 ◦ C in impure helium.

Fig. 7. EBSD results of (a) surface and (b) cross-section of Cr-rich outer oxide.

G.-G. Lee et al. / Nuclear Engineering and Design 271 (2014) 301–308

was difficult to observe the grain growth. It seems that the carbides precipitated along the grain boundaries and twins in the matrix. The bands of carbides remained stable during the corrosion test. The results of the corrosion tests at 900 ◦ C and 850 ◦ C were similar to that of 950 ◦ C. Fig. 6 shows the magnified cross-sectional micrograph and EDS results of the specimen corroded at 950 ◦ C for 250 h. The outer oxide layer has a smooth surface and consistent thickness, and metal islands were isolated in the layer. Below the outer oxide layer, internal oxides with dark gray intruded the matrix, and some internal oxides grew heavily along the grain boundaries. The M23 C6 and M6 C carbides along the grain boundaries were removed from the surface (denoted as carbide-depleted zone). As shown in Fig. 6(b), the EDS results elucidated that the outer oxide layer was composed of Cr-oxide with Ti-oxide, and the internal oxides were Al-oxide. To identify the morphology of the surface outer oxide layer, an electron back-scatter diffraction (EBSD) analysis was performed, and the results are given in Fig. 7. The grains in the outer oxide layer were small, and the size was under ∼0.5 ␮m. The grain directions were distributed around (1 0 2) direction. A cross-sectional EBSD micrograph shows that normal grain growth occurred. This result is comparable with the oxidation behavior in air, and further analysis is planned by our group. Fig. 8 shows the BS SEM micrographs of the specimens corroded at 950 ◦ C. The outer oxide layer, internal oxide, and carbidedepleted zone are observed in all specimens. As the exposure time increases, the thickness of the outer oxide layer and the depth of the internal oxide increased simultaneously. In addition, the carbidedepleted zone depth increased with time. An ascending tendency

305

Table 3 Tentative parabolic rate constants from the microstructure observation of Alloy 617 corroded in impure helium environment. Condition 850 ◦ C 2

−1

Outer oxide thickness (␮m s ) Internal oxide depth (␮m2 s−1 ) Decarburized zone depth (␮m2 s−1 )

900 ◦ C −6

1.92 × 10 4.58 × 10−5 8.43 × 10−5

950 ◦ C −6

3.72 × 10 1.21 × 10−4 2.30 × 10−4

1.07 × 10−5 2.87 × 10−4 8.60 × 10−4

with increasing time was more pronounced as the test temperature increased. The thickness and depth of the layers with various conditions are plotted in Fig. 9. In these results, the representative value of the internal oxide layer depth was determined from the average deeply penetrated depth of the internal oxides at the grain boundaries. The thickness changes showed a parabolic relationship. This relationship can be expressed following equations similar to a mass change equation. x2 = 2kt · t + c

(3)

x, thickness of the layer. The tentative values of kt at various conditions were calculated and are listed in Table 3. Note that the internal oxide layer depth and carbide-depleted depth are measured by observing the grain boundary region, so it may not be the reprehensive value of the matrix. This result suggests that the outer oxide layer thickness, internal oxide depth, and carbide-depleted zone depth may grow to ∼116 ␮m, ∼600 ␮m and ∼1000 ␮m, respectively, if Alloy 617

Fig. 8. BS SEM images of the cross-section of the corroded specimens. The captions in the figures show the temperature and test time of the specimen.

G.-G. Lee et al. / Nuclear Engineering and Design 271 (2014) 301–308

10

1.2 1.0

o

85 0 C o 90 0 C o 95 0 C

8

Thickness, µm

(a) o

Mass change, mg/cm

(a)

2

306

6

4 2

He 950 C o pure He 950 C

0.8 0.6 0.4 0.2 0.0

0 0

50

10 0

15 0

20 0

25 0

0

30 0

50

100

20

10

50

10 0

15 0

20 0

25 0

30 0

50

4

1.0

30

20

10 0 0

50

2

0.8 0.6

He 950 C o Pure He 950 C y=1.10e-6x-0.092

0.4 0.2

y=1.76e-8x-0.001

0.0 -0.2 0.0

5

2.0x10

5

4.0x10

5

6.0x10

5

8.0x10

6

1.0x10

Fig. 10. Comparison of mass change between impure helium environment and pure helium environment at 950 ◦ C: (a) mass change vs. time (hour) and (b) (mass change)2 vs. time (sec).

o

850 C o 900 C o 950 C

40

300

Time, sec

Time, h

Depth, µm

1.2

2

30

0

250

o

0

(c)

(b) o

85 0 C o 90 0 C o 95 0 C

(Mass change) , mg /cm

Penetration Depth, µm

50 40

200

Time, h

Time, h

(b)

150

10 0

150

20 0

25 0

300

Time, h Fig. 9. Changes in (a) outer oxide layer thickness, (b) internal oxide depth, and (c) carbide-depleted zone depth.

oxidizes at 950 ◦ C in the impure helium environment for twenty years. The removal of carbides at the grain boundaries may cause a critical decrease in the creep properties because the grain boundary carbides act as obstacles for grain boundary sliding. Hence, this should be verified to apply Alloy 617 to long-term operation in an impure helium environment. Note that the tentative rate constants were acquired from relatively short-term results; long-term corrosion tests at high temperatures are needed to establish a precise estimate. 3.3. Comparison with pure helium Among the impurity gases in a VHTR, H2 O and CO are the source of oxygen in the oxidation of Alloy 617. To identify the effect of impurities on the oxidation, a corrosion test was carried out in

pure helium environment. The purity of the as-received helium was 99.999%, and the H2 O content was measured as a few ppm by a dew point meter. CO gas was not detected by GC. As pure helium gas passed through the pre-heater, which was maintained at 900 ◦ C, there was ∼2 ppm CO in the pure helium. It is supposed that the oxygen in helium gas were converted into CO by the reaction with graphite in the pre-heater. Fig. 10 plots the mass change of the specimen as a function of exposure time at 950 ◦ C, which was corroded in the pure helium with ∼2 ppm CO. The mass increase was reduced such that it was only 13% of the mass increase in the impure helium environment. The oxidation rate was drastically reduced to 1.6% of the impure helium result. Fig. 11 shows a BS image of the microstructure of the specimens. The outer Cr-oxide layer could not be identified, and the internal Aloxide was not developed widely except some penetration into the grain boundaries. A carbide-depleted zone was found, but the depth decreased compared to the impure helium results. It was calculated that the formation of Al-oxide is more favorable than that of Croxide in a low-oxygen environment (Kim et al., 2011). Hence, it seems that internal Al-oxide was developed before Cr-oxide growth in a pure helium condition. As a preliminary experiment on the effect of CO in pure helium for the corrosion of Alloy 617, corrosion tests were carried out at various CO concentrations. Fig. 12 shows the mass change of specimens with various CO concentrations in helium gas during a 250 h period at 950 ◦ C. The change in mass increased linearly with CO concentration, and with low CO-containing helium was rather high. In other words, the slope of the line was low and the y-intercept was high. It is thought that small amounts (∼2 ppm) of H2 O and O2

G.-G. Lee et al. / Nuclear Engineering and Design 271 (2014) 301–308

307

Fig. 11. BS SEM images of the cross section of the specimens corroded in the pure helium environment at 950 ◦ C for (a) 25 h, (b) 50 h, (c) 100 h, and 250 h.

the critical concentration. In our results, however, it was difficult to see an apparent transition between the oxidation and reduction in the mass changes. Further analysis is planned to elucidate the CO effect on the corrosion of Alloy 617. These results from the pure helium and the CO containing helium experiments indicate that the impurity level of helium gas affected the oxidation behavior drastically. The microstructural change might cause the degradation of Alloy 617. Therefore, longterm and systematic research on the impurity in helium gas is needed to apply Alloy 617 to an IHX in a VHTR.

1.2 o

Mass change, mg/cm

2

CO + pure He, 950 C, 250 h 1.0 Microclimate condition limit

0.8 0.6

reduction

oxidation

0.4 0.2

4. Conclusion 0.0 0

20

40

60

80

100

120

CO concentration in He, ppm Fig. 12. Mass change with CO concentration in helium environment. The experiments are carried out at 950 ◦ C for 250 h.

in CO-containing helium were oxidized first, which caused rather high value of the mass change. It was reported that CO critical concentration under a “microclimate condition” was about 30 ppm at 950 ◦ C (Quadakkers, 1984; Christ et al., 1987; Rouillard et al., 2007; Kim et al., 2010). MO + C = CO + M

(4)

As the CO content is higher than the critical concentration, the metal species and CO react to form oxide (oxidation). On the contrary, a reduction of oxide occurs when CO content is lower than

A high-temperature corrosion test of Alloy 617 in an impure helium environment of a VHTR was carried out, and the oxidation behavior was investigated by the measuring the mass change and observing microstructure. The mass change of the specimens showed a pronounced parabolic oxidation. The activation energy for the oxidation which was calculated from the oxidation rate at 850–950 ◦ C was similar to the activation energy of Cr/Cr2 O3 through grain boundary diffusion. All specimens had similar surface microstructures involving the outer Cr-oxide layers, internal Al-oxides, and carbide-depleted zone along the grain boundaries. A normal grain growth occurred in the outer oxide layer. The corrosion test in pure helium yielded a reduced mass increase. The outer Cr-oxide was difficult to observe, and the depth of the internal oxide and carbide-depleted zone were reduced compared to the impure helium results. The impurity in helium affected the corrosion behavior of Alloy 617 and may cause a decrease in the

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mechanical properties. Therefore, the control of minor impurity gases in VHTR helium is necessary for the application of Alloy 617 to the IHX material of a VHTR. References Birks, N., Meier, G.H., Pettit, F.S., 2006. Introduction to the High-Temperature Oxidation of Metals, 2nd ed. Cambridge University Press, UK, pp. 49. Cabet, C., Duprey, B., 2010. Comparison of the long term high temperature corrosion resistance of candidate alloys for VHTR heat exchangers. In: Proc. of HTR 2010, 18–20, October, Prague, Czech Republic, p. 54. Christ, H.-J., Kunecke, U., Meyer, K., Sockel, G., 1987. High temperature corrosion of the nickel-based Alloy Inconel 617 in helium containing small amounts of impurities. Mater. Sci. Eng. 87, 161. Cook, R.H., 1984. Creep properties of Inconel-617 in air and helium at 800 to 1000 ◦ C. Nucl. Technol. 66, 283. Elder, R., Allen, R., 2009. Nuclear heat for hydrogen production: coupling a very high/high temperature reactor to a hydrogen production plant. Prog. Nucl. Energy 51, 500. Hosoi, Y., Abe, S., 1975. The effect of helium environment on the creep rupture properties of Inconel 617 at 1000 ◦ C. Metall. Trans. A 6A, 1171. Jang, C., Lee, D., Kim, D., 2008. Oxidation behaviour of an Alloy 617 in very hightemperature air and helium environments. Int. J. Press. Vessels Pip. 85, 368. Jo, T.S., Kim, S.-H., Kim, D.-G., Park, J.Y., Kim, Y.D., 2008. Thermal degradation behavior of Inconel 617 Alloy. Met. Mater. Int. 14, 739. Kim, D., Jang, C., Ryu, W.S., 2009. Oxidation characteristics and oxide layer evolution of Alloy 617 and Haynes 230 at 900 ◦ C and 1100 ◦ C. Oxid. Met. 71, 271.

Kim, D.-J., Lee, G.-G., Kim, S.W., Kim, H.P., 2010. The use of thermodynamics and phase equilibra for prediction of the behavior of high temperature corrosion of Alloy 617 in impure helium environment. Corros. Sci. Technol. 9, 164. Kim, D.-J., Lee, G.-G., Jeong, S.J., Kim, W.-G., Park, J.Y., 2011. Investigation on material degradation of Alloy 617 in high temperature impure helium coolant. Nucl. Eng. Technol. 43, 429. Lee, W.J., Kim, Y.W., Chang, J., 2009. Perspectives of nuclear heat and hydrogen. Nucl. Eng. Technol. 41, 412. Lee, G.-G., Jung, S., Kim, D., Kim, W.-G., Park, J.Y., Kim, D.-J., 2011. Microstructural investigation of Alloy 617 creep-ruptured in pure helium environment at 950 ◦ C. Korean J. Mater. Res. 21, 596. Quadakkers, M.R., 1984. Corrosion of high temperature alloys in the primary circuit helium of high temperature gas cooled reactors. Part II: Experimental results. Werkstoffe und Korrosion 36, 335. Ren, W., Swimdeman, R., 2009. A review paper on aging effects in Alloy 617 for Gen IV nuclear reactor applications. J. Press. Vessel Technol. 131, 024002. Rouillard, F., Cabet, C., Wolski, K., Terlain, A., Tabarant, M., Pijolat, M., Valdivieso, F., 2007. High temperature corrosion of a nickel base alloy by helium impurities. J. Nucl. Mater. 362, 248. Shankar, P.S., Natesan, K., 2007. Effect of trace impurities in helium on the creep behavior of Alloy 617 for very high temperature reactor applications. J. Nucl. Mater. 366, 28. Smeltzer, W.W., Young, D.J., 1975. Oxidation properties of transition metals. Prog. Solid State Chem. 10, 17. Special Metals Publication, Number SMC-029, INCONEL Alloy 617, (2005). Totemeier, T.C., Tian, H., 2007. Creep-fatigue-environment interactions in INCONEL 617. Mater. Sci. Eng. A 468–470, 81.