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Microstructural response on the cracking resistance of alloy 600
Materials Letters 61 (2007) 274 – 277 www.elsevier.com/locate/matlet
Microstructural response on the cracking resistance of alloy 600 A. Aguilar a,c , J.L. Albarran a,⁎, H.F. Lopez b , L. Martinez a a
b
Centro de Ciencias Fisicas, UNAM. P.O. Box 48-3, Cuernavaca, Mor. C.P. 62251, Mexico Materials Department, University of Wisconsin–Milwaukee, P.O. Box 784, Milwaukee, Wisconsin 53201, USA c Facultad de Quimica, UNAM, Ciudad Universitaria, D.F., C.P. 04510, Mexico Received 6 April 2005; accepted 14 April 2006 Available online 12 May 2006
Abstract Precipitation of chromium rich carbides promotes the development of a Cr-depleted zone which in turn provided a weak path for the intergranular crack propagation. The role of low temperature anneals on the intergranular cracking resistance (IGC) of alloy 600 was investigated using modified wedge opening loading specimens heat treated at 930, 800 and 600 °C and exposed to high purity water pressurized with hydrogen at 300 °C. Mill annealing at 930 °C did lead to IGC susceptible microstructures. In this condition the alloy 600 exhibited the least crack growth rates (da / dt) of the order of 1.86 × 10− 12 m/s and characterized the substantial work hardening ahead of the crack front. In contrast, annealing at 600 °C (HT600) resulted in increasing IGC susceptibilities. Under these conditions, crack growth rates, da / dt, as high as 7.10 × 10− 10 m/s were found (HT600). Accordingly, significant interactions between the slip bands and the crack path lead to crack bifurcation into the slip planes and cavity formation. © 2006 Elsevier B.V. All rights reserved. Keywords: Nickel alloys; Microstructure; Heat treatment; Grain boundaries; Mechanical properties; Deformation; Fracture
1. Introduction Nickel-based alloy 600 (UNS N06600), in the mill annealed condition, can be susceptible of undergoing intergranular cracking (IGC) in pressurized water reactors (PWR) [1]. Accordingly, the number of reported failures has prompted extensive research efforts into the properties of alloy 600 and into the active IGC mechanisms [2]. The thermomechanical history resulting from mill processing, as well as the heat treatments, has a significant effect on the resultant microstructural development, and hence in the exhibited IGC susceptibility. It has been suggested that the IGC susceptibility arises from a low temperature final anneal (≤ 950 °C) during processing [2]. The grain boundary (gb) features influenced by the heat treatment in turn can account for the exhibited corrosion behavior, as chromium depletion and impurity segregation are known to occur preferentially along the gbs. In addition, the type and distribution of precipitated carbides are expected to play a role on the IGC resistance of alloy 600. Moreover, increasing hydrogen overpressures in high temperature deaerated steam ⁎ Corresponding author. Tel.: +52 777 329 1785; fax: +52 777 329 1775. E-mail address: [email protected] (L. Martinez). 0167-577X/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2006.04.086
have been found to increase the failure susceptibility, although the effect is rather small [3]. More recently, it has been reported that the IGC susceptibility in alloy 600 becomes maximum when approximately 3 ppm of hydrogen is dissolved in a high temperature steam above 350 °C [1]. In addition, recent work has indicated that the resistance to IGSCC can be improved by the precipitation of gb carbides. Nevertheless, there is not enough evidence to conclusively demonstrate that carbide precipitation alone is responsible for the improved IGSCC resistance [4]. Also, theoretical simulations of the processes on the microscopic level in order to model the energy of desorption for hydrogen traps induced by particles precipitation are needed. According to Bruemmer [2], gb carbides can act as stress relievers which Table 1 Heat treatments Specimen
Temperature (°C)
Mill annealed 930 (M.A.) HT600 600 HT800 800
Time (min)
Cooling conditions
60
Air
30 30
Water Water
Cooling rates (°/s) 0.1987 50.74 50.74
A. Aguilar et al. / Materials Letters 61 (2007) 274–277 Table 2 Loading conditions and measured crack growth rates in alloy 600 Condition
KI Crack (MPa m1/2) size (m)
Incubation time (h)
Testing time (h)
(da / dt) (m/s)
As-received M.A. HT600 HT800
49.49 46.51 35.18 34.78
72 720 168 720
384 720 240 720
1.72 ⁎ E− 11 1.86 ⁎ E− 12 7.01 ⁎ E− 10 2.27 ⁎ E− 11
0.0017 0.0016 0.0029 0.0015
effectively blunt the crack tip leading to an overall reduction in crack growth rates. The aim of this work is to further investigate the effect of low temperature anneals on the exhibited microstructure, and hence on the IGC susceptibility of alloy 600. 2. Experimental The chemical composition of alloy 600 (UNS N0600) was: 0.08% C, 15.7% Cr, 6.9% Fe, 0.2% Mn, 0.007% Si, 0.05% Cu, and 77.06% Ni. From this plate, self-loaded modified wedgeopening–loading (MWOL) linear elastic fracture mechanics specimens were machined following the procedure proposed by Novak and Rolfe [5]. Prior to pre-cracking, the specimens were heat treated according to the conditions given in Table 1. Notice that in the milled annealed (M.A.) condition the alloy was annealed for 60 min at 930 °C, but for only 30 min for the low
275
temperature anneals (HT800 and HT600). Cooling was in air (mill anneal) or water (HT800 and HT600). Before exposure to a high temperature steam, the specimens were fatigue pre-cracked using a universal testing machine (Instron 4206). The MWOL specimens were then loaded at stress levels showed in Table 2, using the back-face strain technique proposed by Deans and Richards [6]. This was followed by enclosing the specimens in an instrumented autoclave containing 360 ml of deaerated distilled water pressurized with hydrogen. The autoclave was heated at 300 °C and a total pressure of 200 kPa (steam + hydrogen). Metallographic characterization was carried out in all of the specimens before and after exposure to the simulated PWR environment. A solution containing 2.5 ml HNO3, 10 ml HCl and 7.5 ml glycerol reagent was employed to develop the microstructure. Optical, scanning and transmission electron microscopies were used for microstructural characterization. In the case of TEM observations, thin foils were made of samples sliced using a diamond saw, mechanically grinded and electrolytically thinned using a solution of 8% perchloric acid in ethanol with a current of 125 mA. 3. Results and discussion In general, it was evident that all the specimens exhibited a relatively high density of gb precipitates. Fig. 1.a to d shows the microstructures
Fig. 1. Precipitation found by TEM in the samples: a) as-received, b) MA, c) HT600 and d) HT800.
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Fig. 2. a) Cracking path in the as-received sample, and b) crack path observed for the MA sample.
obtained by TEM. In the as-received condition there was a wide dispersion of Cr23C6 and Cr7C3 carbides within the matrix showing a mean particle diameter of 5.3894 μm and a density of 8.858 ⁎ 10− 6% area. In addition, the gbs were decorated with discontinuous rows of precipitates as seen in Fig. 1.a. Diffraction patterns of these particles indicated that they were a mixture of Cr23C6 and Cr7C3. S. M. Payne and P. C. McIntyre [7] have found that a mixture of gb precipitates can increase the susceptibility to intergranular cracking. Microstructural evaluations indicated that in the M.A. condition, the precipitate density seems to decrease, with the gbs containing discontinuous rows of precipitates. In this case, it was found that the matrix contained a mixture of Cr23C6 and Cr7C3 precipitates, while the dominant precipitate at the gbs was Cr23C6, as shown in Fig. 1.b. C.M. Younes et al. [8] found that sulfur segregation increases the susceptibility to the intergranular corrosion due to the formation of sulfur rich chromium precipitates. In the HT600 specimen, a high precipitation density in the matrix was found to occur in addition to semi-continuous rows of gb precipitates as shown in Fig. 1.c. In particular, the gb and matrix precipitates were identified as Cr23C6 and Cr7C3. The specimen heat treated at 800 °C (HT800) exhibited a less-copious matrix precipitation, but the gbs were decorated with discontinuous precipitates as shown in Fig. 1.d. The discontinuous gb precipitates were identified as a mixture of Cr7C3 and Cr23C6 phases. After loading the M-WOL specimens were exposed to the high temperature steam environment. The total crack lengths and crack growth rates were determined and they are given in Table 2. Notice that in the as-received condition, the crack growth rate is not very different from that exhibited by the HT800 specimen, even though their inherent strengths are vastly different. It was found that in the as-received condition the crack tended to be arrested by the grain boundaries as seen in Fig. 2.a. In addition, fine intergranular voids along the gbs were observed probably due to the
interaction between plastic deformation lines with the intergranular crack. In contrast, from Table 2 it can be observed that the HT600 specimen exhibits the higher cracking susceptibility as indicated by a higher crack growth rate than the one found in the M.A. specimen. The M.A. specimen exhibited the lowest crack growth rate. Apparently, the relatively large applied stress intensity results in significant strain hardening as a result of appreciable matrix precipitate strengthening, that produces relatively slow crack rates. This in agreement with the observations of G. Sui et al. [9], who found that only short cracks are produced after long exposure times in the mill annealed condition in a plastically strained alloy 600. It was found that in the M.A. condition the crack path was predominantly intergranular. In particular, in front of the crack tip voids were found near intergranular precipitates (see Fig. 2.b). The crack path of the HT600 specimen involved cavity coalescence (Fig. 3.a), as well as crack branching. For the sample HT800 the crack path was predominantly transgranular as shown in Fig. 3.b. Notice from this figure the evidence of significant plastic deformation denoted by slip bands interacting with the advancing crack, as is also observed in the development of microcavities, crack branching and crack tip blunting. This in agreement with Bruemmer [2], who found crack tip blunting and crack growth across activated slip planes in front of the crack tip, such as the one shown in Fig. 3.b. Moreover, gb cavitation can be induced by dislocation pile up at gbs [10]. Similar observations have been reported by D. Delafosse and T. Magnin [11], who found that hydrogen diffusion in front of a crack tip can induce microfracture along the slip planes. In the HT800 specimens, the improvements in the IGC resistance can be attributed to the development of a semi-continuous distribution of gb chromium carbides of the type Cr7C3 mixed with Cr23C6 [8]. In contrast, in the HT600 specimens, the relatively poor cracking resistance can be associated with chromium depleted regions near the
Fig. 3. a) Cracking path produced by cavity coalescence and crack branching of the sample HT600, and b) crack path following the slip lines in sample HT800.
A. Aguilar et al. / Materials Letters 61 (2007) 274–277
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gbs making it a weak path for crack propagation. This agrees with recent work by H. Sahlaoui et al. [12], who found that in stainless steel, gb precipitation of Cr23C6 increases the susceptibility to intergranular corrosion mainly due to chromium depletion near the gbs.
CONACyT for financial support Grants C27435-A and 42454. Dr. H.F. Lopez wishes to acknowledge the support provided by The National Science Foundation under contract NSF-INT9314175.
4. Conclusions
References
In the mill annealing condition the crack path was intergranular with crack growth rates (da / dt) of the order of 1.86 × 10− 12 m/s. In the HT800 samples, the crack path was transgranular, and it was also characterized by significant plastic deformation interactions ahead of the crack tip. Moreover, it was characterized by appreciable interactions with slip bands, as well as gb cavitation at the slip bands–crack intersections. However, annealing at 600 °C for 30 min resulted in poor IGC resistance. Under these conditions, crack growth rates, da / dt = 7.10 × 10− 10 m/s, were found with the crack path exhibiting significant branching which is associated to a higher density of carbide type (Cr23C6) in gbs.
[1] Nobuo Totsuka, Yoshito Nishikawa, Nobuo Nakajima, Corrosion (2002) 1 (Paper No. 02523). [2] S.M. Bruemmer, L.A. Charlot, C.H. Henager Jr., Corros. Sci. (Nov. 1988) 782. [3] C.M. Brown and W.J. Mills, Effect of water on mechanical properties and stress corrosion behavior of alloy 600, alloy 690, EN82H welds, and EN52 welds. Corrosion Engineering Section, vol. 55 no. 2 pp. 173, 174. [4] G.S. Was, K. Lian, Corrosion 54–59 (1998) 675. [5] S.R. Novak, S.T. Rolfe, Corrosion 4 (1969) 701. [6] W.F. Deans, C.E. Richards, J. Test. Eval. 7 (1979) 147. [7] S.M. Payne, P. McIntyre, Corrosion 44 (5) (May 1988). [8] C.M. Younes, F.H. Morrissey, G.C. Allen, P. McIntyre, Br. Corros. J. 32–33 (1997) 185. [9] G. Sui, J.M. Titchmarsh, G.B. Heys, J. Congleton, Corros. Sci. 39-3 (1997) 565. [10] M. Uhlemann, B.G. Pound, Corros. Sci. 40-4/5 (1998) 645. [11] D. Delafosse, T. Magnin, Eng. Fract. Mech. 68-6 (2001) 693. [12] H. Sahlaoui, K. Makhlouf, H. Sidhom, J. Philibert, Mater. Sci. Eng., A Struct. Mater.: Prop. Microstruct. Process. 372 (2004) 98.
Acknowledgments The authors wish to acknowledge Osvaldo Flores, Rene Guardian and Anselmo Gonzalez for their technical support and
Microstructural response on the cracking resistance of alloy 600 A. Aguilar a,c , J.L. Albarran a,⁎, H.F. Lopez b , L. Martinez a a
b
Centro de Ciencias Fisicas, UNAM. P.O. Box 48-3, Cuernavaca, Mor. C.P. 62251, Mexico Materials Department, University of Wisconsin–Milwaukee, P.O. Box 784, Milwaukee, Wisconsin 53201, USA c Facultad de Quimica, UNAM, Ciudad Universitaria, D.F., C.P. 04510, Mexico Received 6 April 2005; accepted 14 April 2006 Available online 12 May 2006
Abstract Precipitation of chromium rich carbides promotes the development of a Cr-depleted zone which in turn provided a weak path for the intergranular crack propagation. The role of low temperature anneals on the intergranular cracking resistance (IGC) of alloy 600 was investigated using modified wedge opening loading specimens heat treated at 930, 800 and 600 °C and exposed to high purity water pressurized with hydrogen at 300 °C. Mill annealing at 930 °C did lead to IGC susceptible microstructures. In this condition the alloy 600 exhibited the least crack growth rates (da / dt) of the order of 1.86 × 10− 12 m/s and characterized the substantial work hardening ahead of the crack front. In contrast, annealing at 600 °C (HT600) resulted in increasing IGC susceptibilities. Under these conditions, crack growth rates, da / dt, as high as 7.10 × 10− 10 m/s were found (HT600). Accordingly, significant interactions between the slip bands and the crack path lead to crack bifurcation into the slip planes and cavity formation. © 2006 Elsevier B.V. All rights reserved. Keywords: Nickel alloys; Microstructure; Heat treatment; Grain boundaries; Mechanical properties; Deformation; Fracture
1. Introduction Nickel-based alloy 600 (UNS N06600), in the mill annealed condition, can be susceptible of undergoing intergranular cracking (IGC) in pressurized water reactors (PWR) [1]. Accordingly, the number of reported failures has prompted extensive research efforts into the properties of alloy 600 and into the active IGC mechanisms [2]. The thermomechanical history resulting from mill processing, as well as the heat treatments, has a significant effect on the resultant microstructural development, and hence in the exhibited IGC susceptibility. It has been suggested that the IGC susceptibility arises from a low temperature final anneal (≤ 950 °C) during processing [2]. The grain boundary (gb) features influenced by the heat treatment in turn can account for the exhibited corrosion behavior, as chromium depletion and impurity segregation are known to occur preferentially along the gbs. In addition, the type and distribution of precipitated carbides are expected to play a role on the IGC resistance of alloy 600. Moreover, increasing hydrogen overpressures in high temperature deaerated steam ⁎ Corresponding author. Tel.: +52 777 329 1785; fax: +52 777 329 1775. E-mail address: [email protected] (L. Martinez). 0167-577X/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2006.04.086
have been found to increase the failure susceptibility, although the effect is rather small [3]. More recently, it has been reported that the IGC susceptibility in alloy 600 becomes maximum when approximately 3 ppm of hydrogen is dissolved in a high temperature steam above 350 °C [1]. In addition, recent work has indicated that the resistance to IGSCC can be improved by the precipitation of gb carbides. Nevertheless, there is not enough evidence to conclusively demonstrate that carbide precipitation alone is responsible for the improved IGSCC resistance [4]. Also, theoretical simulations of the processes on the microscopic level in order to model the energy of desorption for hydrogen traps induced by particles precipitation are needed. According to Bruemmer [2], gb carbides can act as stress relievers which Table 1 Heat treatments Specimen
Temperature (°C)
Mill annealed 930 (M.A.) HT600 600 HT800 800
Time (min)
Cooling conditions
60
Air
30 30
Water Water
Cooling rates (°/s) 0.1987 50.74 50.74
A. Aguilar et al. / Materials Letters 61 (2007) 274–277 Table 2 Loading conditions and measured crack growth rates in alloy 600 Condition
KI Crack (MPa m1/2) size (m)
Incubation time (h)
Testing time (h)
(da / dt) (m/s)
As-received M.A. HT600 HT800
49.49 46.51 35.18 34.78
72 720 168 720
384 720 240 720
1.72 ⁎ E− 11 1.86 ⁎ E− 12 7.01 ⁎ E− 10 2.27 ⁎ E− 11
0.0017 0.0016 0.0029 0.0015
effectively blunt the crack tip leading to an overall reduction in crack growth rates. The aim of this work is to further investigate the effect of low temperature anneals on the exhibited microstructure, and hence on the IGC susceptibility of alloy 600. 2. Experimental The chemical composition of alloy 600 (UNS N0600) was: 0.08% C, 15.7% Cr, 6.9% Fe, 0.2% Mn, 0.007% Si, 0.05% Cu, and 77.06% Ni. From this plate, self-loaded modified wedgeopening–loading (MWOL) linear elastic fracture mechanics specimens were machined following the procedure proposed by Novak and Rolfe [5]. Prior to pre-cracking, the specimens were heat treated according to the conditions given in Table 1. Notice that in the milled annealed (M.A.) condition the alloy was annealed for 60 min at 930 °C, but for only 30 min for the low
275
temperature anneals (HT800 and HT600). Cooling was in air (mill anneal) or water (HT800 and HT600). Before exposure to a high temperature steam, the specimens were fatigue pre-cracked using a universal testing machine (Instron 4206). The MWOL specimens were then loaded at stress levels showed in Table 2, using the back-face strain technique proposed by Deans and Richards [6]. This was followed by enclosing the specimens in an instrumented autoclave containing 360 ml of deaerated distilled water pressurized with hydrogen. The autoclave was heated at 300 °C and a total pressure of 200 kPa (steam + hydrogen). Metallographic characterization was carried out in all of the specimens before and after exposure to the simulated PWR environment. A solution containing 2.5 ml HNO3, 10 ml HCl and 7.5 ml glycerol reagent was employed to develop the microstructure. Optical, scanning and transmission electron microscopies were used for microstructural characterization. In the case of TEM observations, thin foils were made of samples sliced using a diamond saw, mechanically grinded and electrolytically thinned using a solution of 8% perchloric acid in ethanol with a current of 125 mA. 3. Results and discussion In general, it was evident that all the specimens exhibited a relatively high density of gb precipitates. Fig. 1.a to d shows the microstructures
Fig. 1. Precipitation found by TEM in the samples: a) as-received, b) MA, c) HT600 and d) HT800.
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A. Aguilar et al. / Materials Letters 61 (2007) 274–277
Fig. 2. a) Cracking path in the as-received sample, and b) crack path observed for the MA sample.
obtained by TEM. In the as-received condition there was a wide dispersion of Cr23C6 and Cr7C3 carbides within the matrix showing a mean particle diameter of 5.3894 μm and a density of 8.858 ⁎ 10− 6% area. In addition, the gbs were decorated with discontinuous rows of precipitates as seen in Fig. 1.a. Diffraction patterns of these particles indicated that they were a mixture of Cr23C6 and Cr7C3. S. M. Payne and P. C. McIntyre [7] have found that a mixture of gb precipitates can increase the susceptibility to intergranular cracking. Microstructural evaluations indicated that in the M.A. condition, the precipitate density seems to decrease, with the gbs containing discontinuous rows of precipitates. In this case, it was found that the matrix contained a mixture of Cr23C6 and Cr7C3 precipitates, while the dominant precipitate at the gbs was Cr23C6, as shown in Fig. 1.b. C.M. Younes et al. [8] found that sulfur segregation increases the susceptibility to the intergranular corrosion due to the formation of sulfur rich chromium precipitates. In the HT600 specimen, a high precipitation density in the matrix was found to occur in addition to semi-continuous rows of gb precipitates as shown in Fig. 1.c. In particular, the gb and matrix precipitates were identified as Cr23C6 and Cr7C3. The specimen heat treated at 800 °C (HT800) exhibited a less-copious matrix precipitation, but the gbs were decorated with discontinuous precipitates as shown in Fig. 1.d. The discontinuous gb precipitates were identified as a mixture of Cr7C3 and Cr23C6 phases. After loading the M-WOL specimens were exposed to the high temperature steam environment. The total crack lengths and crack growth rates were determined and they are given in Table 2. Notice that in the as-received condition, the crack growth rate is not very different from that exhibited by the HT800 specimen, even though their inherent strengths are vastly different. It was found that in the as-received condition the crack tended to be arrested by the grain boundaries as seen in Fig. 2.a. In addition, fine intergranular voids along the gbs were observed probably due to the
interaction between plastic deformation lines with the intergranular crack. In contrast, from Table 2 it can be observed that the HT600 specimen exhibits the higher cracking susceptibility as indicated by a higher crack growth rate than the one found in the M.A. specimen. The M.A. specimen exhibited the lowest crack growth rate. Apparently, the relatively large applied stress intensity results in significant strain hardening as a result of appreciable matrix precipitate strengthening, that produces relatively slow crack rates. This in agreement with the observations of G. Sui et al. [9], who found that only short cracks are produced after long exposure times in the mill annealed condition in a plastically strained alloy 600. It was found that in the M.A. condition the crack path was predominantly intergranular. In particular, in front of the crack tip voids were found near intergranular precipitates (see Fig. 2.b). The crack path of the HT600 specimen involved cavity coalescence (Fig. 3.a), as well as crack branching. For the sample HT800 the crack path was predominantly transgranular as shown in Fig. 3.b. Notice from this figure the evidence of significant plastic deformation denoted by slip bands interacting with the advancing crack, as is also observed in the development of microcavities, crack branching and crack tip blunting. This in agreement with Bruemmer [2], who found crack tip blunting and crack growth across activated slip planes in front of the crack tip, such as the one shown in Fig. 3.b. Moreover, gb cavitation can be induced by dislocation pile up at gbs [10]. Similar observations have been reported by D. Delafosse and T. Magnin [11], who found that hydrogen diffusion in front of a crack tip can induce microfracture along the slip planes. In the HT800 specimens, the improvements in the IGC resistance can be attributed to the development of a semi-continuous distribution of gb chromium carbides of the type Cr7C3 mixed with Cr23C6 [8]. In contrast, in the HT600 specimens, the relatively poor cracking resistance can be associated with chromium depleted regions near the
Fig. 3. a) Cracking path produced by cavity coalescence and crack branching of the sample HT600, and b) crack path following the slip lines in sample HT800.
A. Aguilar et al. / Materials Letters 61 (2007) 274–277
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gbs making it a weak path for crack propagation. This agrees with recent work by H. Sahlaoui et al. [12], who found that in stainless steel, gb precipitation of Cr23C6 increases the susceptibility to intergranular corrosion mainly due to chromium depletion near the gbs.
CONACyT for financial support Grants C27435-A and 42454. Dr. H.F. Lopez wishes to acknowledge the support provided by The National Science Foundation under contract NSF-INT9314175.
4. Conclusions
References
In the mill annealing condition the crack path was intergranular with crack growth rates (da / dt) of the order of 1.86 × 10− 12 m/s. In the HT800 samples, the crack path was transgranular, and it was also characterized by significant plastic deformation interactions ahead of the crack tip. Moreover, it was characterized by appreciable interactions with slip bands, as well as gb cavitation at the slip bands–crack intersections. However, annealing at 600 °C for 30 min resulted in poor IGC resistance. Under these conditions, crack growth rates, da / dt = 7.10 × 10− 10 m/s, were found with the crack path exhibiting significant branching which is associated to a higher density of carbide type (Cr23C6) in gbs.
[1] Nobuo Totsuka, Yoshito Nishikawa, Nobuo Nakajima, Corrosion (2002) 1 (Paper No. 02523). [2] S.M. Bruemmer, L.A. Charlot, C.H. Henager Jr., Corros. Sci. (Nov. 1988) 782. [3] C.M. Brown and W.J. Mills, Effect of water on mechanical properties and stress corrosion behavior of alloy 600, alloy 690, EN82H welds, and EN52 welds. Corrosion Engineering Section, vol. 55 no. 2 pp. 173, 174. [4] G.S. Was, K. Lian, Corrosion 54–59 (1998) 675. [5] S.R. Novak, S.T. Rolfe, Corrosion 4 (1969) 701. [6] W.F. Deans, C.E. Richards, J. Test. Eval. 7 (1979) 147. [7] S.M. Payne, P. McIntyre, Corrosion 44 (5) (May 1988). [8] C.M. Younes, F.H. Morrissey, G.C. Allen, P. McIntyre, Br. Corros. J. 32–33 (1997) 185. [9] G. Sui, J.M. Titchmarsh, G.B. Heys, J. Congleton, Corros. Sci. 39-3 (1997) 565. [10] M. Uhlemann, B.G. Pound, Corros. Sci. 40-4/5 (1998) 645. [11] D. Delafosse, T. Magnin, Eng. Fract. Mech. 68-6 (2001) 693. [12] H. Sahlaoui, K. Makhlouf, H. Sidhom, J. Philibert, Mater. Sci. Eng., A Struct. Mater.: Prop. Microstruct. Process. 372 (2004) 98.
Acknowledgments The authors wish to acknowledge Osvaldo Flores, Rene Guardian and Anselmo Gonzalez for their technical support and