Creep-fatigue of High Temperature Materials for VHTR: Effect of ...

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or length of the bulk cracking did not appear to be related to location relative to the primary crack. Hence, the cracking features (number, average length, cumulated length of the cracks) were calculated for the all of the areas combined. Figure 10 shows the results graphically as a function of hold time. As mentioned previously, there was no cracking observed in the continuous cycle fatigue specimens. It was also evidenced that increasing hold times induced relatively similar amounts of grain boundary cracking. Additionally, whatever the environment, air or impure helium, the material exhibited a similar degree of bulk grain boundary cracking. Fig. 9. Schematic of the statistical method used to quantify damage in the LCF and creep-fatigue deformed specimens. The stress axis is horizontal and in the plane of the page. It is believed that creep-fatigue interaction may be caused by bulk cracking accelerating the surface fatigue crack propagation. Equivalent creep strain accumulated for a 180s hold and a 1800s hold (because of the rapid relaxation) would produce similar amount of grain boundary cracking, resulting in the same reduction of the creep-fatigue life. Although the main factor in the drop in the fatigue life when a hold is introduced may be the creep damage promoting intergranular crack propagation, oxidation may also play a role. IV.D. Role of oxidation One faces a challenge in discussing the role of environment at elevated temperature. In an ideal case, comparable creep-fatigue testing performed under an inert atmosphere would be used as a baseline, exemplifying the effect of oxidation on fatigue life. However, creep-fatigue in inert conditions is not practically achievable at high temperature. 320 a. b. c. Number of cracks Average length (μm) Total length (μm) 110 100 90 80 70 60 50 40 35 30 25 20 15 10 5 Proceedings of ICAPP 2011 Nice, France, May 2-5, 2011 Paper 11284 0.3% - air 0.6% - air 0.3% - helium 0 0 500 1000 1500 2000 �� Hold Time (sec) 0.3% - air 0.6% - air 0.3% - helium 30 0 500 1000 Hold �� time (sec) 1500 2000 2000 1500 1000 500 0.3% - air 0.6% - air 0.3% - helium 0 0 500 1000 1500 2000 �� Hold time (sec) Fig. 10. Grain boundary cracking analysis for LCF and creep-fatigue testing in air and in helium; a) total number of cracks, b) crack average length, and c) total length.
Exposure at 950°C in vacuum and in (never completely) pure helium or argon is accompanied by decarburization. Dissolution of intergranular carbides to an increasing depth fosters grain boundary sliding and the formation of cavities and wedge-cracks in the carbide-free region 3,18 . These metallurgical instabilities contribute to the fatigue crack initiation and/or growth, and the fatigue life may be significantly lower than observed in air 3 . Therefore, characterization of the cycled alloy microstructure on the surface and in the surrounding of the cracks will be used to get an insight of the possible role of oxidation in the fatigue crack initiation and growth. Close examination of cross sections of the creepfatigue tested specimen revealed that surface grain boundaries were either cracked or showed precipitation of finger-like alumina (see Figure 7a). Intergranular alumina precipitates are typical of Alloy 617 oxidized in air or in an impure but oxidizing helium (as is the process gas in the present study): the oxygen potential at the scale/alloy interface, set by the dissociation of chromia, is very low but high enough for aluminum as well as silicon to be oxidized 8 . Precipitation is preferred at grain boundaries which are short diffusion pathways for oxygen. It is believed that the alumina veins may act as preferential sites for crack initiation. Once initiated the cracks will open, exposing new metal surfaces to the gas phase and oxidation of the crack edge would occur with growth of a chromiumrich oxide. Eventually, the chromium oxide may fully fill the crack, forming a large region filled with chromium oxide. Following an equivalent process, formation of intergranular alumina was observed at grain boundaries intersecting with main oxidized cracks, including ahead of the crack tip, as shown in Figure 7b. Cracks are likely to propagate preferentially through these brittle alumina precipitates and the crack growth may be accelerated. Therefore, internal oxidation of alumina first below the external chromia scale and then ahead of the crack tip may increase both crack initiation and propagation rate. This hypothesis is consistent with the observed intergranular cracking and may account for at least a portion of the reduction in the fatigue life with a tensile hold IV. CONCLUSIONS This paper reviews two previous publications by Carroll and co-workers 10,11 on creep-fatigue testing of Alloy 617 in air and in a VHTR simulated impure helium at 950°C at total strain ranges of 0.3% and 0.6%. This work supports acceptation of Alloy 617 in the nuclear section (section III) of the ASME code for application at high temperature as well as licensing of VHTR systems with IHX. The creep-fatigue behavior of Alloy 617 at 950°C may be influenced by both the environment and the creep 321 Proceedings of ICAPP 2011 Nice, France, May 2-5, 2011 Paper 11284 deformation that occurs during the tensile hold. The shift in crack initiation mode from transgranular during fatigue deformation to intergranular during creep-fatigue deformation likely results from the environmental influence. Oxidation may also play a role in determining the crack propagation mode, tending the creep-fatigue crack propagation to intergranular as opposed to transgranular. However, longer hold times did not result in the further decrease in cycle life suggesting that the environmental influence on creep-fatigue crack propagation is not the controlling factor. Measurements of the oxidation phenomena are planned for a more quantitative approach. Constant accumulated creep damage despite an increasing hold time agrees well with the experimental tendency of cycles to failure saturating. The creep damage in the alloy bulk, in the form of grain boundary cracking, may have accelerated the crack propagation. The amount of grain boundary cracking was similar among all of the hold times due to the rapid saturation during the stress relaxation. Creep-fatigue testing will be performed under another helium environment producing different microstructure and corrosion products to determine any change in the deformation mechanisms. ACKNOWLEDGMENTS The authors would like to acknowledge Joel Simpson, DC Haggard, Dave Swank, Randy Lloyd, Tammy Trowbridge, Todd Morris and Barry Rabin for conducting the experiments and the microscopy. This work was supported through the U.S. Department of Energy Nuclear Energy. REFERENCES 1. H. BURLET et al., "Evaluation of Nickel-Based Materials for VHTR Heat Exchanger," Proceedings of the Structural Materials for Innovative Nuclear Systems (SMINS) Workshop, OECD Publishing, London, 79 (2008) 2. J. CORUM and J.J. Blass, “Rules for Design of Alloy 617 Nuclear Components to Very High temperature,” Proc. Fatigue, Fracture and Risk PVP Conference, 215, 147 (1191). 3. K.B.S. RAO, H. Schiffers, H. Schuster and H. Nickel, "Creep-Fatigue Interaction of Inconel 617 at 950°C in Simulated Nuclear Reactor Helium," Metallurgical Transactions A, 19A, 359 (1988).
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