The effect of thermomechanical processing on the creep behavior of Alloy 690

Transcription

1 Materials Science and Engineering A 473 (2008) The effect of thermomechanical processing on the creep behavior of Alloy 690 C.J. Boehlert Department of Chemical Engineering and Materials Science, Michigan State University, East Lansing, MI 48824, USA Received 22 February 2007; received in revised form 17 March 2007; accepted 20 March 2007 Abstract The effect of thermomechanical processing on the microstructure and elevated-temperature creep behavior of Alloy 690 was investigated. Commercially available sheet was subjected to four cycles of cold rolling to 25% deformation followed by annealing at 1000 C for 1 h. Both the resultant microstructure and the original microstructure were characterized using electron backscattered diffraction. The thermomechanically processed microstructure exhibited a slightly lower fraction of twins and a smaller average grain size than the original microstructure. Tensile creep experiments were performed in an open-air environment at temperatures between 650 and 690 C and stresses between 75 and 172 MPa. The measured creep stress exponents (4 5) activation energies ( kj/mol) suggested that dislocation creep with lattice self-diffusion was dominant. The thermomechanically processed microstructure exhibited significantly worse creep resistance than the original as-processed microstructure. Thus, cyclic strain and annealing processing, which has been shown to improve the ductility-dip cracking susceptibility of Alloy 690, is not recommended for enhancing the creep resistance Elsevier B.V. All rights reserved. Keywords: Nickel-based alloy; Electron backscattered diffraction; Creep; Microstructure 1. Introduction Alloy 690, a commercially available nickel-based alloy with a nominal composition close to Ni 30Cr 10Fe (wt.%), is attractive for pressurized-water nuclear reactor components because of its superior corrosion resistance [1 3]. Due primarily to its intergranular stress corrosion cracking resistance (IGSCC), Alloy 690 is intended to replace Alloy 600 (Ni 16Cr 9Fe (wt.%)) as a steam generator tube material in pressurized-water reactors [4]. One potential means to improve the IGSCC resistance of Alloy 690 is through thermomechanical processing (TMP) treatments which alter the grain boundary character distribution (GBCD). Strain-recrystallization-based TMP treatments have resulted in improved IGSCC resistance of pure Ni and Ni-based alloys [5 7]. However, strain-recrystallizationbased TMP investigations of Alloy 690 have been limited [8,9]. Xia et al. [8] have evaluated the effects of thermomechanical processing on the distributions of twin boundaries in Alloy 690. They found that the strain and annealing processes significantly influenced the distributions of twins [8]. They performed cold address: boehlert@egr.msu.edu. rolling between 5 and 50% followed by annealing at 1100 C for 5 min. With small strains (5% cold rolled material), the twin boundaries were parts of clusters, and the overall fraction of special boundaries was However, with larger amounts of cold rolling deformation almost no twin clusters existed, and the overall special boundary fraction was 0.47 [8]. Dave et al. [9] evaluated the effect of TMP on the microstructure and ductility-dip cracking susceptibility of Alloy 690. In their work, the as-received wrought mill microstructure was subjected to a repeated cycle of 25% cold rolling followed by an anneal at 1000 C for 1 h. This cycle was repeated four times and the total reduction of the initial sheet was approximately 67%. The TMP material exhibited a slightly greater percentage of special boundaries (50 55%) than the as-received material (40% special boundaries), and there were regions in which the random boundary network was effectively disrupted. The strainrecrystallization processed material exhibited a higher ductility recovery temperature and a higher minimum ductility than the material that did not undergo cyclic strain-recrystallization processing. Thus the additional TMP treatment had a beneficial impact on the alloy s resistance to cracking and thereby improved its fracture behavior. A correlation between intact random boundary networks and intergranular brittle fracture modes /$ see front matter 2007 Elsevier B.V. All rights reserved. doi: /j.msea

2 234 C.J. Boehlert / Materials Science and Engineering A 473 (2008) was observed, and in regions where this random boundary network had been disrupted, transgranular ductile fracture occurred. Their work suggested that the GBCD and, more specifically, the topological connectivity of random boundaries, had an effect on material resistance to ductility-dip cracking, although there were other intervening microstructural factors mentioned in their study. A complete understanding of the physical mechanisms responsible for the elevated-temperature creep behavior and associated microstructure property relationships of Alloy 690 is lacking. In particular, it has yet to be established if the GBCD, which has been shown to have a significant influence on the mechanical deformation behavior, including creep resistance, of pure Ni and Ni-based superalloy systems [10 14], has an effect on the creep behavior. This work was intended to evaluate processing microstructure property relationships of Alloy 690. In particular the effect of TMP on the microstructure (GBCD, grain size, etc.) and high-temperature creep behavior was evaluated. The material was chosen because it is a commercially used alloy and, as such, the aim of this study was to show relevance to actual engineering materials in use. 2. Experimental procedure The as-processed (AP) Alloy 690 sheet was mill annealed at 1066 C. To produce the TMP sheet, the AP sheet material was subjected to the following strain annealing sequence: cold rolling to 25% deformation followed by a solution treatment at 1000 C for 1 h followed by air cooling. This sequence was repeated four times. This sequence was chosen based on identical TMP treatments performed by Dave et al. [9]. The original thickness of the AP sheet was approximately 8 mm while the final thickness of the TMP sheet was 2.7 mm. This constituted a total reduction of the AP sheet s thickness of approximately 66%. Bulk chemical analysis was performed using inductively coupled plasma optical emission spectroscopy and inert gas fluorescence. Each AP and TMP sheet material was sectioned and metallographically polished to prepare it for imaging. Spatially resolved electron backscattered diffraction (EBSD) orientation maps were obtained from polished sections using a FEI XL-30 Field Emission Gun scanning electron microscope (SEM). The EBSD hardware and software were manufactured by EDAX-TSL, Inc. (Draper, UT, USA). The specimens were ground mechanically by 15, 6, and 1 m diamond suspension for 10 min, respectively, and then polished by 0.06 m colloidal silica for 60 min. The GBCD was characterized using the orientation maps to determine the orientation relationships between grains. For each map, more than 500 grain boundaries were analyzed using a step size of between 0.5 and 2 m. Low-angle boundaries (LABs) were defined as those boundaries containing misorientations between 2 and 15. Brandon s criteria [15] were used to distinguish between general high-angle boundaries (GHABs) and coincident site lattice boundaries (CSLBs). The reported fractions of GHABs, LABs, CSLBs, and twins ( 3) were the averaged values taken from several orientation maps, performed on the cross-sections, rolling faces, or longitudinal sections of the sheets. Grain size was determined using the mean line intercept method [16,17]. Blanks from the AP and TMP sheet materials were machined, using either electrodischarge machining or a mill, into a flat dogbone geometry used for tensile and tensile creep specimens. Open-air tensile creep experiments were performed on vertical Applied Test System, Incorporated (Butler, PA, USA) load frames with a 20:1 lever-arm ratio. Applied Test System single zone furnaces were used to heat the specimens to within ±3 Cof the target temperature. Specimen temperatures were monitored by three chromel alumel type K thermocouples located within the specimen s reduced section. Creep strain was monitored during the tests using a linear variable differential transformer that was connected to a 25.4 mm gage length Applied Test System high-temperature extensometer. The extensometer was attached directly to the gage section of each sample. The testing temperatures and stresses ranged between 650 and 690 C and MPa, respectively. All creep specimens were loaded parallel to the rolling direction, and the experiments were conducted such that the specimens were soaked at the creep temperature for at least 60 min prior to applying load in order to minimize the thermal stresses. After the creep strain had proceeded well into the secondary regime, either the load or temperature was changed or the creep test was discontinued. The tested specimens were cooled under load to minimize recovery of the deformed structures. Room-temperature (RT) tensile tests were performed at a strain rate of 10 3 s 1 using an Instron 4206 tensile testing machine. Strain was measured during the tensile tests with an extensometer attached directly to the gage section of the sample. 3. Results and discussion 3.1. Microstructure Table 1 lists the measured alloy composition. Fig. 1 illustrates backscattered electron SEM images of the AP microstructures. Fig. 2 illustrates backscattered electron SEM images of the TMP microstructures. The TMP material exhibited an average equiaxed grain size of 8.9 m and the AP material exhibited an average grain size of 16.1 m. It is noted that the grain boundaries were intact and uncracked in both the AP and TMP microstructures. The AP microstructure was somewhat banded which has also been observed in previous studies of this alloy [9]. Banding was also exhibited in the TMP microstructures. Fine M 23 C 6 carbides were observed to decorate the grain boundaries in both microstructures, see Figs. 1b and 2a, while larger carbides were distributed throughout the microstructures. Table 1 Measured Composition of the IN 690 alloy in wt.% Ni Cr Fe Mn Si C Si Al Mo Co O N

3 C.J. Boehlert / Materials Science and Engineering A 473 (2008) Fig. 1. Backscattered electron SEM images (a) and (b) of the as-processed (AP) microstructure. Fig. 2. Backscattered electron SEM images (a) and (b) of the thermomechanically-processed (TMP) microstructure. Fig. 3 illustrates an EBSD orientation map for the AP microstructure, which exhibited an equiaxed microstructure where the grain orientations were distributed fairly evenly. Pole figure analysis indicated that the microstructure was not strongly textured. Fig. 4 illustrates an EBSD orientation map for the TMP microstructure. The GHAB fractions were the majority in both microstructures. The volume fraction of twins observed was 0.10 and 0.05 for the AP and TMP microstructures, respectively. The volume fraction of LABs observed was 0.25 and 0.30 for the AP and TMP microstructures, respectively. Thus, the TMP material exhibited a slightly lower fraction of twins and a slightly higher fraction of LABs compared to the AP material. If we consider that the combination of the LAB and CSLB fractions (< 29) constitute the overall fraction of special boundaries, then the maximum special boundary fraction was approximately 0.35 for both the AP and TMP microstructures. This value is significantly less than the maximum special boundary fraction value, , measured by Dave et al. [9] Tensile and creep behavior Representative RT tensile curves are illustrated in Fig. 5. The AP microstructure exhibited a slightly greater yield strength (364 MPa) than the TMP microstructure (344 MPa). The ultimate tensile strengths for both microstructures were close to 725 MPa and significant work hardening was exhibited. The creep strain-life history resembled that for most metals exhibiting three stages of creep: primary, secondary, and tertiary [18,19]. Fig. 6 illustrates strain versus time plots for each of the materials at T = 650 C and σ = 172 MPa. Once the creep stress was reached there appeared to be a short incubation period before a positive creep strain was achieved. This was common for both the AP and TMP conditions. The AP material exhibited significantly greater creep resistance than the TMP material. No grain boundary cracking was observed even for samples deformed to over 15% strain. Fig. 7a illustrates the minimum creep strain rate versus applied stress behavior at T = 650 C, while Fig. 7b illustrates the minimum creep strain versus inverse temperature response at σ = 125 MPa.

4 236 C.J. Boehlert / Materials Science and Engineering A 473 (2008) Fig. 5. RT stress vs. strain behavior for as-processed (AP) and thermomechanically processed (TMP) specimens. Fig. 3. EBSD orientation map for the as-processed (AP) microstructure. The TMP condition exhibited lower creep resistance than the AP material at all the applied stress and temperatures evaluated. In particular, the minimum creep rate was approximately 0.5 orders of magnitude faster for the TMP microstructure compared with the AP condition. The measured creep stress exponents were between 4 and 5 at T = 650 C, and the Q app values were between 320 kj/mol (TMP) and 368 kj/mol (AP) at σ = 125 MPa. Such creep parameters are similar to those measured for other similar nickel-based superalloys [14] and suggest a dislocation climb mechanism [20,21] is active. Based on the measured exponents and activation energies, dislocation creep with lattice self-diffusion was the suggested dominant creep mechanism over the entire applied stress range examined (75 MPa < σ < 172 MPa). It is felt that the smaller grain size and the lower twin fractions could have contributed to the worse creep resistance exhibited by the TMP material. The presence of the grain boundary carbides may have also contributed to the creep behavior discrepancy as Fig. 4. EBSD orientation map for the thermomechanically processed (TMP) microstructure. Fig. 6. Creep strain vs. time plots for as-processed (AP) and thermomechanically processed (TMP) specimens at T = 650 C and σ = 172 MPa.

5 C.J. Boehlert / Materials Science and Engineering A 473 (2008) microstructure exhibited significantly greater creep resistance than the TMP microstructure. For the applied stresses and temperature evaluated, the measured creep parameters suggested that dislocation creep was the dominant creep deformation mechanism. Overall, the TMP sequence degraded the creep resistance, refined the equiaxed grain size and promoted a lower fraction of twins in the microstructure. Thus TMP involving cyclic cold rolling and annealing is not suggested for improving the creep resistance of Alloy 690. Acknowledgements This work was supported by the National Science Foundation through grant DMR The author is grateful to Mr. Nathan Eisinger (Special Metals Corporation, Huntington, WV, USA) for directing the alloy processing effort. References Fig. 7. Creep plots used to determine the creep parameters n and Q app : (a) minimum creep rate vs. applied stress at T = 650 C and (b) ln minimum creep rate vs. 1/T at σ = 125 MPa. the TMP material exhibited more grain boundary area, due to its fine grain size, than the AP material. 4. Summary This work evaluated the effect of TMP on the microstructure and creep behavior of Alloy 690. The AP microstructure exhibited a slightly larger volume fraction of twins than the TMP microstructure, yet the overall fraction of special boundaries was similar for both microstructures. The AP material exhibited a grain size almost twice that of the TMP material. The AP [1] R.S. Dutta, R. Tewari, Br. Corros. J. 3 (34) (1999) [2] A.J. Sedricks, J.W. Schultz, M.A. Cordovi, Boshoku Gijutsu 28 (1979) [3] C. Cheung, U. Erb, G. Palumbo, Mater. Sci. Eng. 185A (1994) [4] M. Thuvander, K. Stiller, Mater. Sci. Eng. 281A (2000) [5] G. Palumbo, K.T. Aust, Acta. Metall. Mater. 11 (38) (1990) [6] G. Palumbo, U.S. Patent 5,817,193 (1998). [7] B. Alexandreanu, B.M. Capell, G. Was, Mater. Sci. Eng. 300A (2001) [8] S. Xia, B.X. Zhou, W.J. Chen, W.G. Wang, Scripta Mater. 54 (2006) [9] V.R. Dave, M.J. Cola, M. Kumar, A.J. Schwartz, G.N.A. Hussen, Welding J. (January 2004) 1-S 5-S (American Welding Society and Welding Research Council). [10] E.M. Lehockey, G. Palumbo, Mater. Sci. Eng. 237A (1997) [11] G.S. Was, V. Thaveeprungsriporn, D.C. Crawford, J. Met. (1998) [12] V. Thaveeprungsriporn, G. Was, Metall. Mater. Trans. A 28A (1997) [13] G. Palumbo, K.T. Aust, in: D. Wolf, S. Yip (Eds.), Special Properties of Grain Boundaries, Materials Interfaces: Atomic Level Structure and Properties, Chapman and Hall, NY, 1989, pp [14] C.J. Boehlert, D.S. Dickmann, N.C. Eisinger, Metall. Mater. Trans. A 37A (1) (2006) [15] D.G. Brandon, Acta Metall. 14 (1966) [16] J.E. Hilliard, Met. Prog. 78 (1964) [17] Standard Test Methods for Determining Average Grain Size, ASTM Designation E112-96e3, American Society for Testing and Materials, West Conshohocken, PA. [18] R.W. Evans, B. Wilshire, Creep of Metals and Alloys, The Institute of Metals, New York, NY, [19] R.W. Hertzberg, Deformation and Fracture Mechanics of Engineering Materials, fourth ed., John Wiley and Sons, New York, NY, [20] T.G. Langdon, P. Yavari, Acta Metall. 30 (1982) [21] J. Weertman, Trans. Am. Soc. Met. 61 (1968)