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

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- An area free of grain boundary carbides has evolved in the subsurface. Such a carbide-depletion was also evidenced in the region surrounding the primary creep-fatigue cracks. Carbide dissolution was related to the diffusion of chromium to the surface (of the specimen or of the crack) to form the chromia scale. Chromium removal then may have destabilized the grain boundary Cr-rich carbides in the subsurface area. III.F. Bulk evolution For pure fatigue specimens (and oxidized coupons without loading as well), the material bulk was clean from cracking. On the contrary, cracks were observed in the bulk of the creep-fatigue tested specimens. Examples of these cracks are shown in the images in Figure 8. a. b. 50 µm 50 µm Fig. 8. Grain boundary damage observed in the bulk material of specimens deformed at a 0.3% total strain range with a) a 180 s hold in impure helium and b) a 600 s hold in air. The stress axis is horizontal and in the plane of the page. The cracks were relatively thin and typically followed grain boundaries perpendicular to the stress axis. Their length varied, from several micrometers to several grain diameters in length. In some cases, there were boundaries with multiple small cracks. Although this type of damage was randomly distributed in the bulk, most of the grain boundaries were completely free of cracking and these intact grain boundaries did not exhibit cavitation and 318 Proceedings of ICAPP 2011 Nice, France, May 2-5, 2011 Paper 11284 instead microscopically appeared free of grain boundary damage (at magnifications of up to 100,000 times). The bulk cracks were absent of any oxide and thus it can be concluded that they were not surface connected and exposed to the environment. IV. DISCUSSION IV.A. Effect of an air or helium environment The cycle lives of continuous cycle and creep-fatigue of Alloy 617 in air and in VHTR simulated helium were within the same scatter band at 950°C, as shown in Figure 3. This tendency is rather consistent with the small number of previous studies on the influence in fatigue of air compared to impure helium. Nagato and coworkers 14 observed that the fatigue life of Hastelloy X, a nickel base alloy rich in Cr, was longer in impure helium than in air but the effect of environment was less remarkable at temperatures above 800°C and for creep-fatigue. At temperatures of 750°, 850° and 950°C, Meurer and coworkers 5 observed a slight increase in the fatigue life of Alloy 617 tested in impure helium compared to air at a total strain of 0.3%; this difference was not evidenced at higher strain ranges. Strizak and coworkers 15 observed that a helium environment was not detrimental to fatigue life of Alloy 617 at temperatures of up to 704°C and Rao and coworkers 3 concluded that at 950°C the influence of environment, if any, is strongly reduced for longer hold times. In addition to an equivalent fatigue life, the tensile/compressive peak stress versus cycle, the hysteresis loop, and the stress relaxation during the peak tensile hold were similar under both atmospheres (see Figure 4). No difference in the crack initiation behavior, propagation mode, or oxidation of the cracks was evidenced between the air and impure helium cycled specimens. Finally, metallurgical characterization of the specimens cyclically deformed in air and in helium failed to reveal any significant differences in the oxidation products or surface morphology. The corrosion product development and the surface evolution are consistent with the general corrosion behavior of Alloy 617 under a high temperature oxidizing environment as is air or controlled VHTR helium coolant (helium with significant water vapor and carbon monoxide partial pressures at 950°C). In an atmosphere with a sufficient oxidation potential, Alloy 617 forms a chromia surface scale and aluminum oxidizes internally 9,13 . These similarities in the overall behavior of the specimens cycled in air and in simulated impure VHTR helium allow for the joint discussion of the results for both environments without specific reference to the testing atmosphere.
IV.B. Basis for creep-fatigue interaction Creep-fatigue interaction is a combination of creep damage and fatigue damage which accelerates the deformation process. Typically three types of failure are defined: fatigue-dominated, creep-dominated, and creepfatigue interaction 16-18 . Fatigue dominated failures are typically observed at lower to intermediate temperatures and transgranular crack initiation and propagation is observed. On the contrary, creep dominated failures occur at high temperatures and intergranular crack initiation and propagation with extensive creep cavitation occurs. Creepfatigue interaction is illustrated as mixed-mode crack propagation and the presence of creep cavitation on the grain boundaries in the bulk material. Cavity formation and fatigue crack initiation and propagation independently develop, then interaction occurs if one mode accelerates the other damage process. The creep and fatigue failure modes interact through cavitation accelerating the crack initiation or propagation process or fatigue deformation enhancing cavitation 18 . Identifying the dominant failure mode and the influence of environment is important for life prediction and modeling efforts. On the one hand, the creep-fatigue specimens did not fail in a fatigue-dominated manner at 950°C. The primary crack in the continuously cycled specimens initiated and propagated in a transgranular manner (fatigue dominated manner). Oxidized grain boundaries were occasionally present but did not result in specimen failure. For all cases in creep-fatigue, several longer cracks likely initiated and did propagate intergranularly. Rao and coworkers 3 also reported transgranular cracking of Alloy 617 at 950°C in fatigue cycled specimens and the addition of a 60s tensile hold resulted in the cracking initiating transgranularly and propagating by a mixed mode. Besides, at lower strain rates in a helium environment: approximately 10 -4 /s, the crack propagation mode switched to intergranular. On the other hand, massive cavitation which would have caused creep dominated failure was not observed in any of the creep-fatigue specimens deformed at 950°C with holds as long as 1800s. Grain boundary cavities were depicted for many alloys in the regime of creep-fatigue interaction 3,16,18-20 . A detailed assessment correlated the size and number of grain boundary cavities and the accumulated creep damage for 316 stainless steel 16 . It was also found that a critical amount of deformation was required before creep cavities nucleated. Recent work by Lillo and co-workers 21 suggested that significant creep cavitation did not occur in Alloy 617 until greater than 10% creep strain during creep deformation in the temperature range from 900°C to 1000°C. Inelastic strain was evaluated by integrating the stress relaxation data from Figure 5 [this calculation was done by S. Sham from Oak Ridge National Laboratory]. The total inelastic strain was 319 Proceedings of ICAPP 2011 Nice, France, May 2-5, 2011 Paper 11284 calculated by integrating the stress relaxation curve at a mid-life cycle. The accumulated creep strain was estimated to be less than 2.5% and 0.5% for respectively the 0.3% and the 0.6% total strain tests,. Therefore, creep cavities are not expected to be formed in the present testing of Alloy 617 at 950°C. IV.C. Effect of a hold period in tensile strain Similar to what has been observed for stainless steels and nickel base alloys, the introduction of hold times at peak tensile strain in continuous cycle fatigue reduces the fatigue life of Alloy 617 (Figure 3). For example at a 0.3% total strain, the introduction of a 180s hold period at peak tensile strain decreased the life by approximately a factor of 2. A greater decrease in cycles to failure at lower strain ranges was also observed, consistent with the results on Nimonic PE-16, a Ni-Cr superalloy 19 . The drop in cycle life is likely associated with the change in the cracking mode. The shift in crack propagation mode from transgranular to intergranular may be a result of grain boundary damage from creep processes that accelerates intergranular crack propagation (see paragraph IV.B) but may be also influenced by the environment accelerating grain boundary crack propagation and promoting intergranular failure. Interestingly, the increasing duration of the peak tensile hold at 0.3% total strain did not further decrease the number of cycles to failure; that is to say, the 180s hold creep-fatigue life was similar to that of the 1800s hold. The amount of creep damage, estimated by the accumulated inelastic strain, is relatively similar with increasing hold periods as the stresses during the tensile hold were almost completely relaxed by the end of the first 180s. This suggests creep-fatigue interaction in the current testing. Although creep intergranular cavities were not observed, bulk cracking was evidenced in all of the creepfatigue deformed specimens as illustrated in Figure 8. Because intergranular crack was not found in the continuously cycle fatigue specimens (or in unloaded material), it was assumed that bulk cracking was caused by creep deformation. The amount of bulk material cracking was quantified to determine if a relationship existed between the duration of the hold time and the amount of grain boundary damage. The total number of cracks, total length of cracks, and the average length of the cracks were determined for a (0.5mm)² area at five locations, as shown schematically in Figure 9. The areas analyzed were 1mm in each direction (in the plane of the page) from the primary crack tip and 3mm in each direction along the stress axis (in the plane of the page) from the primary crack tip. For all specimens, the cracking was statistically similar among the five areas and no general trends were observed. In particular, the amount
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