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Causes of breakdown of creep strength in 9Cr–1.8W–0.5Mo–VNb steel
Materials Science and Engineering A 428 (2006) 270–275
Causes of breakdown of creep strength in 9Cr–1.8W–0.5Mo–VNb steel Jae Seung Lee a,1 , Hassan Ghassemi Armaki a , Kouichi Maruyama b,∗ , Taro Muraki c , Hitoshi Asahi c b
a Graduate School of Engineering, Tohoku University, 6-6-02 Aobayama, Sendai 980-8579, Japan Graduate School of Environmental Studies, Tohoku University, 6-6-02 Aobayama, Sendai 980-8579, Japan c Steel Research Laboratories, Nippon Steel Corporation, Futtsu 293-8511, Japan
Received 7 March 2006; received in revised form 8 May 2006; accepted 9 May 2006
Abstract Premature breakdown of creep strength is a serious problem to be solved in long-term creep of advanced high Cr ferritic steels. The material studied was ASTM grade 92 steel crept at 550–650 ◦ C for up to 63 151 h. Stress exponent for rupture life decreases from 17 in short-term creep to 8 in long-term creep, confirming the breakdown in the steel. The steel shows ductile to brittle transition with increasing rupture life, and the breakdown accords with the onset of brittle intergranular fracture. Creep cavities are nucleated at coarse precipitates of Laves phase along grain boundaries. These findings suggest the following story of the breakdown of creep strength. Laves phase precipitates and grows during creep exposure. Coarsening of Laves phase particles over a critical size triggers the cavity formation and the consequent brittle intergranular fracture. The brittle fracture causes the breakdown. The coarsening of Laves phase can be detected non-destructively by means of hardness testing of the steel exposed to elevated temperature without stress. © 2006 Elsevier B.V. All rights reserved. Keywords: Grade 92 steel; Premature breakdown of creep strength; Ductile to brittle transition; Coarsening of Laves phase; Non-destructive evaluation
1. Introduction Long-term creep properties of structural materials to be used in high temperature plants are evaluated from short-term rupture data with the aid of time–temperature-parameter (TTP) methods [1]. However, the activation energy for rupture life of advanced 9–12Cr ferritic steels often decreases in long-term creep, and the conventional TTP methods overestimate long-term rupture life of such steels [2]. This is called “premature failure” [3–5]. The premature failure brings about an unexpected shutdown of plants. In 2004 such a premature failure of an advanced high Cr ferritic steel took place in a power plant in Japan, and it is an urgent engineering issue how to detect and prevent the premature failure. If we can postpone the decrease of activation energy in long-term creep of a steel, we can substantially extend service life of the steel. The decrease in activation energy is always accompanied by a decrease in stress exponent for rupture life. This aspect of the change in rupture behavior is named ∗
Corresponding author. Tel.: +81 22 795 7324; fax: +81 22 795 7324. E-mail address: [email protected] (K. Maruyama). 1 Present address:Wire Rod Research Group, POSCO, Pohang 790-785, Republic of Korea. 0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.05.010
as “premature breakdown of creep strength” [6,7]. Thorough understanding of causes of the changes in activation energy and stress exponent is necessary for detecting and preventing the premature failure and the premature breakdown. Several explanations have been proposed on the premature failure and breakdown. MX carboniterides are one of the most effective strengthening elements in the advanced high Cr ferritic steels. However, Z phase is formed at the expense of MX carbonitrides during high temperature exposure. Z phase grows rapidly and cannot contribute to strengthening. The premature breakdown has often been explained by this story of the disappearance of MX carboniterides [6–8]. These researches have studied on creep of ferritic rotor steels tempered at a low temperature (below 700 ◦ C). In such steels, M2 X carboniterides are formed, thereby promoting the formation of Z phase. However, M2 X phase is not formed in ferritic boiler steels tempered at a temperature close to 800 ◦ C [7]. Semba and Abe [9] have found a similar breakdown of creep strength in a high Cr ferritic steel in which Z phase cannot be formed because of the absence of nitrogen atom. Kushima et al. [10] have found a recovered zone along prior austenite grain boundaries in Mod. 9Cr–1Mo steel, and have proposed the local recovery as a cause of the premature failure reported in the steel. However, sufficient proof supporting
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their proposal has not been given yet. W is one of the most effective solid solution elements in advanced high Cr ferritic steels. The matrix contains supersaturated W atoms in the early stage of creep. They are consumed to from Fe2 W Laves phase during creep exposure. The premature breakdown of creep strength has been explained also by the loss of solid solution hardening [11]. Transition from ductile transgranular fracture to brittle intergranular fracture often occurs in long-term creep of advanced high Cr ferritic steels tempered close to 800 ◦ C for boiler application. It has been well demonstrated in the fracture mechanism maps [12] that the transition in fracture mechanism brings about decreases in activation energy and stress exponent for rupture life, and therefore results in premature failure and breakdown. Maruyama et al. [2,4] have explained the premature failure of 11Cr–2W–0.4Mo–1Cu–VNb steel (grade 122) by the transition in fracture mode. However, the ductile to brittle transition does not occur in ferritic rotor steel tempered at a low temperature [4]. The causes of the premature failure and the premature breakdown of creep strength have not been fully understood yet. The major cause may be different depending on alloy composition and heat treatment condition. 9Cr–1.8W–0.5Mo–VNb steel is an advanced high Cr ferritic steel used extensively in ultra super critical power plants over the world. In this paper, the cause of the premature breakdown in grade 92 steel is studied paying special attention to precipitation and growth of Laves phase. Non-destructive evaluation by means of hardness testing is employed for detecting the breakdown. 2. Experimental procedure
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cles (dark contrast) from Laves phase (bright contrast). This technique enables quantitative characterization of precipitation and coarsening of Laves phase during creep. Size distribution of Laves phase precipitates was determined through measurement of 300–1000 particles. 3. Mechanism of premature breakdown 3.1. Creep properties Creep rupture lives of the grade 92 steel studied are plotted in Fig. 1(a) as a function of creep stress. The following Arrhenius type equation describes stress σ and temperature T dependence of creep rupture life tr : Q −n tr = t0 σ exp (1) RT where t0 is a constant, n the stress exponent, Q the apparent activation energy, R the universal gas constant, and T is the absolute temperature. There are two regions with different values of Q and n in the figure, and the dash-dot line divides the two regions: Q = 667 kJ/mol and n = 17 in the short-term and high stress region H, and Q = 624 kJ/mol and n = 8.4 in the long-term and low stress region L. The values of Q and n decrease in the long-term region, confirming the breakdown of creep strength in the steel. However, the difference in Q values between the two regions is not so significant as compared to grade 122 steel [2]. This is due to the lower Cr concentration of grade 92 steel. Reduction of area measured after creep rupture is plotted in Fig. 1(b) as a function of rupture life. The rupture ductility drops beyond a critical point, demonstrating ductile to brittle
Chemical composition of the steel studied is listed in Table 1. The steel is called ASTM grade 92, and was normalized at 1050 ◦ C for 1 h, followed by tempering at 780 ◦ C for 3 h, to have a tempered martensitic lath structure. Creep specimens, 6 mm in diameter and 30 mm in gage length, were tested at 550, 600 and 650 ◦ C under constant load in air. The longest rupture life was 63 151 h. Microstructures of crept specimens were analyzed by field emission scanning electron microscopy (FESEM) and field emission transmission electron microscopy (FETEM). Segregation of W was measured by FETEM equipped with energy dispersive X-ray spectroscopy (EDS). The specimens for SEM analysis were etched with a mixed solution of 5 ml hydrochloric acid, 1 g picric acid and 150 ml ethanol. Back scattered electron (BSE) images of SEM were used to distinguish M23 C6 partiTable 1 Chemical composition of the steel studied (mass%) C Si Mn Cr W Mo V Nb N
0.10 0.20 0.44 9.2 1.82 0.45 0.20 0.05 0.05
Fig. 1. (a) Stress vs. rupture life plot and (b) reduction of area after creep rupture of 9Cr–1.8W–0.5Mo–VNb steel studied. The dash-dot line divides the data points into two regions H and L with different stress exponents. Fractographs corresponding to the solid symbols are shown in Fig. 2.
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Fig. 3. Fracture mechanism map of the steel studied. The dashed line is the boundary between transgranular and intergranular fracture modes. The intergranular fracture field is further divided into wedge cracking and isolated creep cavities.
Fig. 2. Fracture surfaces after (a) 150.2 h creep at 550 ◦ C and 270 MPa (ductile fracture), and (b) 26 783 h at 650 ◦ C and 80 MPa (brittle fracture).
high density of Laves phase particles are present. This type is labeled as “wedge crack” in Fig. 3. The other type is “isolated cavity” formed independently at coarse Laves phase particles on prior austenite grain boundaries (arrows in Fig. 4(c)). This type is usually observed after long-term creep in the intergranular fracture region as shown in Fig. 3. Though M23 C6 carbide and Lave phase particles can be potential nucleation sites of creep cavities, coarse Laves phase particles are the preferential site in the present steel. 3.3. Evolution of Laves phase particles
550 ◦ C
transition of fracture mode. The dotted line of was predicted from the measurements at 600 and 650 ◦ C. The dashed curve in Fig. 1(a) indicates the points of ductile to brittle transition. The transition points accord fairly well with the boundary between regions H and L.
Evolution process of Laves phase is depicted in Fig. 5. The figure includes the measurements in gage and grip portions, but there is not obvious difference in particle size between the two portions. The following equation represents the coarsening of particles, namely Ostwald ripening [13,14]:
3.2. Fracture mode
d m − d0m = Kd t
Fig. 2 exhibits the representative fractographs taken after (a) ductile and (b) brittle fracture. The corresponding data points are indicated in Fig. 1(b) with the solid symbols. The dimple pattern in Fig. 2(a) is typical of ductile transgranular fracture. The fracture appearance seen in Fig. 2(b) is obviously different from Fig. 2(a). The absence of dimple pattern suggests brittle intergranular fracture, though morphology of the fracture surface is not clear due to oxidation. Results of the examination of creep fracture modes are summarized in Fig. 3. There are two types of fracture modes; transgranular and intergranular fracture. The intergranular fracture region is further divided into two domains according to the location of cavity formation. Secondary electron (SE) and back scatted electron (BSE) images of the same area are shown in Fig. 4. Laves phase with a higher average atomic number is easily identified by its bright contrast in the BSE image. The other particles seen in the micrographs are mostly M23 C6 carbides. At high stresses in the intergranular fracture region (Fig. 4(a) and (b)), cavities are formed at the triple grain junction at which a
where d0 and d are the average diameter of particles measured before and after high temperature exposure, respectively, t the exposure time, and Kd is a material constant proportional to diffusivity. When d0 is negligibly small, Eq. (2) can be simplified to Qd m d = K0 t exp − (3) RT
(2)
where Qd is the activation energy for particle coarsening, and K0 is a constant. The exponent m is characteristic of the coarsening mechanisms of precipitates; for example, m = 3 for the coarsening controlled by lattice diffusion, and 4 for grain boundary diffusion [13]. It is evident in Fig. 5 that there are two stages F and C with different exponents m on the log d versus log t curves. Apparent activation energies Qd and exponent m obtained are Qd = 348 kJ/mol and m = 7 in stage F, and Qd = 342 kJ/mol and m = 4 in stage C. Amount of W in precipitates has been reported as a function of exposure time in Ref. [15] on grade 92 steel. It
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Fig. 4. Cavity formation in specimens crept ((a) and (b)) for 26 875 h at 600 ◦ C (wedge crack), and ((c) and (d)) for 26 783 h at 650 ◦ C (isolated cavities); (a) and (c) are secondary electron images, and (b) and (d) are back scattered electron images of the same location. The arrows indicate cavities.
increases with increasing exposure time, and saturates at an equilibrium value. The upward arrows in Fig. 5 indicate the points at which 95% of W has precipitated. In stage C, W precipitation has finished and the basic assumption of Ostwald ripening is fulfilled. The Qd value of 342 kJ/mol is close to that for lattice diffusion of W in 9Cr steel [16]. Therefore, it can be concluded that lattice diffusion controls the growth of Laves phase in stage C, though m = 4 is larger than the expected value of 3 for this mechanism. In stage F, nucleation and coarsening of Laves phase proceeds simultaneously. Since the W precipitation delays at lower temperature [15], the apparent activation energy for coarsening, 348 kJ/mol in stage F, should be larger than the true value. This in turn suggests grain boundaries as diffusion path in this stage.
Fig. 5. Coarsening of Laves phase in gage and grip portions. The dash-dot line is the boundary between regions F and C with different coarsening characteristics.
Segregation of W was analyzed by EDS attached to FETEM in a grip portion after 646 h and 8955 h exposure at 650 ◦ C, and the results are given in Fig. 6. The two exposure durations are in stage F (646 h), at which the supersaturated W is still present, and stage C (8955 h), after finishing the precipitation of Laves phase, respectively. In stage F, the W peaks measured at grain boundary is obviously higher than those from the matrix. This result suggests that a part of supersaturated W segregates to grain boundaries. The segregated W is consumed to form precipitates of Laves phase on grain boundaries. The path of W flow for the precipitation can be grain boundaries as speculated. The W peaks almost disappear in stage C after finishing the precipitation. In
Fig. 6. EDS analyses of matrix, Laves phase and grain boundary after 646 h exposure at 650 ◦ C (stage F in Fig. 5), and of grain boundary after 8955 h at 650 ◦ C (stage C in Fig. 5).
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this stage, Ostwald ripening of Laves phase probably proceeds through lattice diffusion, as suggested by the activation energy for the ripening. 3.4. Cause of fracture mechanism change The stress exponent n for rupture life of the grade 92 steel decreases in long-term region L. The dash-dot line in Fig. 1 is the boundary between the two regions H and L. The breakdown of creep strength in region L is typical of advanced high Cr ferritic steels containing W [2,4]. Transition from ductile transgranular fracture to brittle intergranular fracture occurs in the grade 92 steel. The dashed line in Fig. 1(a) represents the boundary of the ductile to brittle transition. The two boundaries of the breakdown and the fracture mode change coincide fairly well with each other, thereby indicating a close correlation between the change in fracture mechanism and the breakdown of creep strength. Similar good correlations have been reported on grade 122 steel [2] and austenitic stainless steel [17]. The fracture mechanism maps published by Ashby et al. [12] have clearly demonstrated the breakdown at the boundaries between different fracture mechanisms. Formation and coalescence of creep cavities along grain boundaries brings about the brittle intergranular fracture. The cavities are usually attached to coarse particles of Laves phase (see Fig. 4), suggesting that the precipitation and coarsening of Laves phase are responsible for the intergranular fracture. The downward arrows in Fig. 5 indicate the points of ductile to brittle transition. The points correspond to an average diameter of Laves phase of 130 nm (127 nm at 600 ◦ C and 140 nm at 650 ◦ C). This fact together with the nucleation site of cavities suggests that coarsening of Laves phase particles over a critical size triggers the brittle grain boundary fracture in the steel studied. Creep cavities are nucleated on coarse particles of Laves phase in brittle intergranular fracture region of grade 122 steel also [2]. This agreement suggests that the major cause of the ductile to brittle transition is the same in grade 92 and 122 steels for boiler application. 4. Detection of onset of intergranular fracture Advanced high Cr ferritic steels with W have been widely used in electric power plants [18]. Accurate evaluation of their creep lives is necessary for safe and efficient operation of those power plants, and detection of the breakdown of creep strength is of crucial importance. Coarsening of Laves phase particles over a critical size triggers the transition from ductile to brittle fracture, and the transition of fracture mechanism brings about the breakdown. Therefore, measurements of cavity density and particles size of Laves phase are essential for the detection. The measurements can be done with replica sampling. However, it is more practical if the transition can be detected more easily, for example, by hardness testing. Fig. 7 shows the changes of Vickers hardness as a function of exposure time. The hardness was measured in gage and grip portions after creep rupture. The solid symbols indicate the presence of Laves phase particles after the exposure. The high Cr ferritic
Fig. 7. Change of Vickers hardness in (a) gage and (b) grip portions of the steel studied. Downward arrows: Laves phase have grown to 130 nm. Upward arrows: Boundary of ductile to brittle transition. Open arrows: 0.5 mass% W has precipitated.
steel has a tempered martensitic lath structure, and recovery of the dislocation substructure primarily brings about the decrease in hardness after creep deformation in the gage portion [19]. However, the softening curve does not provide any signal of the breakdown and the ductile to brittle transition. As seen in Fig. 7(b), the hardness changes even in the grip portion, and the curves reveal two stages. The hardness remains almost constant during the short-term exposure, and decreases gradually after the long-term exposure. The solid downward arrows in Fig. 7(b) indicate the points at which Laves phase particles have grown to the average size of 130 nm. The upward arrows correspond to the ductile to brittle transition. These points well accord with the drop of hardness in grip. So, the drop of hardness can be a useful signal of the particle coarsening and the consequent ductile to brittle transition. The drop of hardness can reasonably be correlated to coarsening of Laves phases as follows. The steel was tempered at 780 ◦ C for 3 h. In grip portion without creep deformation, recovery of the tempered martensitic lath structure does not proceed evidently, since the recovery is a strain assisted process [4]. Loss of solid solution hardening and/or precipitation hardening should be the cause of the hardness decrease in the grip portion. Coarsening of M23 C6 carbides and MX carbonitrides starts at the beginning of creep, and cannot explain the late onset of the decrease in hardness after 103 h at 650 ◦ C. Laves phase precipitates during the high temperature exposure at the expense of W in the matrix and on grain boundaries. Removal of W should result in decrease of hardness. The downward open arrows in Fig. 7(b) indicates the points at which 0.5 mass% W has precipitated [15]. However, the drop of hardness does not start at those points. Precipitation hardening due to fine Laves phase particles can compensate the loss of solid solution hardening in the early stage of Laves phase precipitation as reported
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in literatures [20]. This is why the hardness remains constant after the precipitation of Laves phase. Laves phase grows faster than M23 C6 and MX particles. The hardness starts to drop when Laves phase particles have grown and failed to compensate the loss of solution hardening. Therefore, the hardness drop in grip portion coincides with the coarsening of Laves phase particles. 5. Summary 1. Stress exponent for rupture life decreases from 17 to 8.4 in long-term creep, confirming the typical breakdown of creep strength in ASTM grade 92 steel. Transition from ductile transgranular fracture to brittle intergranular fracture is the major cause of the breakdown. 2. Most of cavities are nucleated at coarse Laves phase particles on grain boundaries. 3. Cavity formation at coarse particles of Laves phase triggers the brittle intergranular fracture. The ductile to brittle transition occurred when Laves phase has reached an average diameter of 130 nm. 4. One can non-destructively detect the coarsening of Laves phase particles and the consequent brittle fracture by means of hardness testing during high temperature exposure without stress. References [1] R. Viswanathan, Damage Mechanisms and Life Assessment of High Temperature Components, ASM International, Metals Park, OH, 1989 (Chapter 3).
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[2] K. Maruyama, J.S. Lee, in: I.A. Shibli, et al. (Eds.), Creep and Fracture in High Temperature Components—Design and Life Assessment Issues, DEStech Publications, Lancaster, PA, 2005, p. 372. [3] H. Cerjak, P. Hofer, B. Schaffernak, ISIJ Int. 39 (1999) 874. [4] K. Maruyama, K. Sawada, J. Koike, ISIJ Int. 41 (2001) 641. [5] K. Kimura, H. Kushima, K. Sawada, in: A. Strang, et al. (Eds.), Parsons 2003 Engineering Issues in Turbine Machinery, Power Plant and Renewables, Institute of Materials, Minerals and Mining, London, 2003, p. 443. [6] H. Danielson, J. Hald, Proceedings of the fourth EPRI International Conference on Advances in Materials Technology for Fossil Power Plants, EPRI, Palo Alto, CA, 2004 (CD-ROM). [7] R. Agamennone, W. Blum, C. Gupta, J.K. Chakravartty, Acta Mater. (2006), in press. [8] A. Strang, V. Vodarek, Mater. Sci. Technol. 12 (1996) 552. [9] H. Semba, F. Abe, Proceedings of the Symposium on USC Steels for Fossil Power Plants 2005, Beijing, China, 2005, p. 157. [10] H. Kushima, K. Kimura, F. Abe, J. Iron Steel Inst. Jpn. 85 (1999) 841. [11] T. Muraki, Y. Hasegawa, M. Ohgami, Key Eng. Mater. 171–174 (2000) 499. [12] M.F. Ashby, C. Gandihi, D.M.R. Taplin, Acta Metall. 27 (1979) 699. [13] I.M. Lifshitz, V.V. Slyozov, J. Phys. Chem. Solids 19 (1961) 35. [14] C. Wagner, Z. Electrochem. 65 (1961) 581. [15] J. Hald, in: E. Metcalfe (Ed.), Proceedings of the EPRI/National Power Conference on the New Steels for Advanced Plant up to 620 ◦ C, EPRI, Palo Alto, CA, 1995, p. 152. [16] J. Cermak, J. Ruzickova, A. Pokorna, Acta Mater. 12 (1996) 5007. [17] H. Nakakuki, K. Maruyama, H. Oikawa, K. Yagi, J. Iron Steel Inst. Jpn. 81 (1995) 220. [18] F. Masuyama, in: J. Lecomte-Beckers, et al. (Eds.), Materials for Advanced Power Engineering 1998, Forschungszentrum Julich GmbH, Julich, 1998, p. 1807. [19] K. Sawada, K. Maruyama, Y. Hasegawa, T. Muraki, Key Eng. Mater. 171–174 (2000) 109. [20] F. Abe, Mater. Sci. Eng. A 319–321 (2001) 770.
Causes of breakdown of creep strength in 9Cr–1.8W–0.5Mo–VNb steel Jae Seung Lee a,1 , Hassan Ghassemi Armaki a , Kouichi Maruyama b,∗ , Taro Muraki c , Hitoshi Asahi c b
a Graduate School of Engineering, Tohoku University, 6-6-02 Aobayama, Sendai 980-8579, Japan Graduate School of Environmental Studies, Tohoku University, 6-6-02 Aobayama, Sendai 980-8579, Japan c Steel Research Laboratories, Nippon Steel Corporation, Futtsu 293-8511, Japan
Received 7 March 2006; received in revised form 8 May 2006; accepted 9 May 2006
Abstract Premature breakdown of creep strength is a serious problem to be solved in long-term creep of advanced high Cr ferritic steels. The material studied was ASTM grade 92 steel crept at 550–650 ◦ C for up to 63 151 h. Stress exponent for rupture life decreases from 17 in short-term creep to 8 in long-term creep, confirming the breakdown in the steel. The steel shows ductile to brittle transition with increasing rupture life, and the breakdown accords with the onset of brittle intergranular fracture. Creep cavities are nucleated at coarse precipitates of Laves phase along grain boundaries. These findings suggest the following story of the breakdown of creep strength. Laves phase precipitates and grows during creep exposure. Coarsening of Laves phase particles over a critical size triggers the cavity formation and the consequent brittle intergranular fracture. The brittle fracture causes the breakdown. The coarsening of Laves phase can be detected non-destructively by means of hardness testing of the steel exposed to elevated temperature without stress. © 2006 Elsevier B.V. All rights reserved. Keywords: Grade 92 steel; Premature breakdown of creep strength; Ductile to brittle transition; Coarsening of Laves phase; Non-destructive evaluation
1. Introduction Long-term creep properties of structural materials to be used in high temperature plants are evaluated from short-term rupture data with the aid of time–temperature-parameter (TTP) methods [1]. However, the activation energy for rupture life of advanced 9–12Cr ferritic steels often decreases in long-term creep, and the conventional TTP methods overestimate long-term rupture life of such steels [2]. This is called “premature failure” [3–5]. The premature failure brings about an unexpected shutdown of plants. In 2004 such a premature failure of an advanced high Cr ferritic steel took place in a power plant in Japan, and it is an urgent engineering issue how to detect and prevent the premature failure. If we can postpone the decrease of activation energy in long-term creep of a steel, we can substantially extend service life of the steel. The decrease in activation energy is always accompanied by a decrease in stress exponent for rupture life. This aspect of the change in rupture behavior is named ∗
Corresponding author. Tel.: +81 22 795 7324; fax: +81 22 795 7324. E-mail address: [email protected] (K. Maruyama). 1 Present address:Wire Rod Research Group, POSCO, Pohang 790-785, Republic of Korea. 0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.05.010
as “premature breakdown of creep strength” [6,7]. Thorough understanding of causes of the changes in activation energy and stress exponent is necessary for detecting and preventing the premature failure and the premature breakdown. Several explanations have been proposed on the premature failure and breakdown. MX carboniterides are one of the most effective strengthening elements in the advanced high Cr ferritic steels. However, Z phase is formed at the expense of MX carbonitrides during high temperature exposure. Z phase grows rapidly and cannot contribute to strengthening. The premature breakdown has often been explained by this story of the disappearance of MX carboniterides [6–8]. These researches have studied on creep of ferritic rotor steels tempered at a low temperature (below 700 ◦ C). In such steels, M2 X carboniterides are formed, thereby promoting the formation of Z phase. However, M2 X phase is not formed in ferritic boiler steels tempered at a temperature close to 800 ◦ C [7]. Semba and Abe [9] have found a similar breakdown of creep strength in a high Cr ferritic steel in which Z phase cannot be formed because of the absence of nitrogen atom. Kushima et al. [10] have found a recovered zone along prior austenite grain boundaries in Mod. 9Cr–1Mo steel, and have proposed the local recovery as a cause of the premature failure reported in the steel. However, sufficient proof supporting
J.S. Lee et al. / Materials Science and Engineering A 428 (2006) 270–275
their proposal has not been given yet. W is one of the most effective solid solution elements in advanced high Cr ferritic steels. The matrix contains supersaturated W atoms in the early stage of creep. They are consumed to from Fe2 W Laves phase during creep exposure. The premature breakdown of creep strength has been explained also by the loss of solid solution hardening [11]. Transition from ductile transgranular fracture to brittle intergranular fracture often occurs in long-term creep of advanced high Cr ferritic steels tempered close to 800 ◦ C for boiler application. It has been well demonstrated in the fracture mechanism maps [12] that the transition in fracture mechanism brings about decreases in activation energy and stress exponent for rupture life, and therefore results in premature failure and breakdown. Maruyama et al. [2,4] have explained the premature failure of 11Cr–2W–0.4Mo–1Cu–VNb steel (grade 122) by the transition in fracture mode. However, the ductile to brittle transition does not occur in ferritic rotor steel tempered at a low temperature [4]. The causes of the premature failure and the premature breakdown of creep strength have not been fully understood yet. The major cause may be different depending on alloy composition and heat treatment condition. 9Cr–1.8W–0.5Mo–VNb steel is an advanced high Cr ferritic steel used extensively in ultra super critical power plants over the world. In this paper, the cause of the premature breakdown in grade 92 steel is studied paying special attention to precipitation and growth of Laves phase. Non-destructive evaluation by means of hardness testing is employed for detecting the breakdown. 2. Experimental procedure
271
cles (dark contrast) from Laves phase (bright contrast). This technique enables quantitative characterization of precipitation and coarsening of Laves phase during creep. Size distribution of Laves phase precipitates was determined through measurement of 300–1000 particles. 3. Mechanism of premature breakdown 3.1. Creep properties Creep rupture lives of the grade 92 steel studied are plotted in Fig. 1(a) as a function of creep stress. The following Arrhenius type equation describes stress σ and temperature T dependence of creep rupture life tr : Q −n tr = t0 σ exp (1) RT where t0 is a constant, n the stress exponent, Q the apparent activation energy, R the universal gas constant, and T is the absolute temperature. There are two regions with different values of Q and n in the figure, and the dash-dot line divides the two regions: Q = 667 kJ/mol and n = 17 in the short-term and high stress region H, and Q = 624 kJ/mol and n = 8.4 in the long-term and low stress region L. The values of Q and n decrease in the long-term region, confirming the breakdown of creep strength in the steel. However, the difference in Q values between the two regions is not so significant as compared to grade 122 steel [2]. This is due to the lower Cr concentration of grade 92 steel. Reduction of area measured after creep rupture is plotted in Fig. 1(b) as a function of rupture life. The rupture ductility drops beyond a critical point, demonstrating ductile to brittle
Chemical composition of the steel studied is listed in Table 1. The steel is called ASTM grade 92, and was normalized at 1050 ◦ C for 1 h, followed by tempering at 780 ◦ C for 3 h, to have a tempered martensitic lath structure. Creep specimens, 6 mm in diameter and 30 mm in gage length, were tested at 550, 600 and 650 ◦ C under constant load in air. The longest rupture life was 63 151 h. Microstructures of crept specimens were analyzed by field emission scanning electron microscopy (FESEM) and field emission transmission electron microscopy (FETEM). Segregation of W was measured by FETEM equipped with energy dispersive X-ray spectroscopy (EDS). The specimens for SEM analysis were etched with a mixed solution of 5 ml hydrochloric acid, 1 g picric acid and 150 ml ethanol. Back scattered electron (BSE) images of SEM were used to distinguish M23 C6 partiTable 1 Chemical composition of the steel studied (mass%) C Si Mn Cr W Mo V Nb N
0.10 0.20 0.44 9.2 1.82 0.45 0.20 0.05 0.05
Fig. 1. (a) Stress vs. rupture life plot and (b) reduction of area after creep rupture of 9Cr–1.8W–0.5Mo–VNb steel studied. The dash-dot line divides the data points into two regions H and L with different stress exponents. Fractographs corresponding to the solid symbols are shown in Fig. 2.
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Fig. 3. Fracture mechanism map of the steel studied. The dashed line is the boundary between transgranular and intergranular fracture modes. The intergranular fracture field is further divided into wedge cracking and isolated creep cavities.
Fig. 2. Fracture surfaces after (a) 150.2 h creep at 550 ◦ C and 270 MPa (ductile fracture), and (b) 26 783 h at 650 ◦ C and 80 MPa (brittle fracture).
high density of Laves phase particles are present. This type is labeled as “wedge crack” in Fig. 3. The other type is “isolated cavity” formed independently at coarse Laves phase particles on prior austenite grain boundaries (arrows in Fig. 4(c)). This type is usually observed after long-term creep in the intergranular fracture region as shown in Fig. 3. Though M23 C6 carbide and Lave phase particles can be potential nucleation sites of creep cavities, coarse Laves phase particles are the preferential site in the present steel. 3.3. Evolution of Laves phase particles
550 ◦ C
transition of fracture mode. The dotted line of was predicted from the measurements at 600 and 650 ◦ C. The dashed curve in Fig. 1(a) indicates the points of ductile to brittle transition. The transition points accord fairly well with the boundary between regions H and L.
Evolution process of Laves phase is depicted in Fig. 5. The figure includes the measurements in gage and grip portions, but there is not obvious difference in particle size between the two portions. The following equation represents the coarsening of particles, namely Ostwald ripening [13,14]:
3.2. Fracture mode
d m − d0m = Kd t
Fig. 2 exhibits the representative fractographs taken after (a) ductile and (b) brittle fracture. The corresponding data points are indicated in Fig. 1(b) with the solid symbols. The dimple pattern in Fig. 2(a) is typical of ductile transgranular fracture. The fracture appearance seen in Fig. 2(b) is obviously different from Fig. 2(a). The absence of dimple pattern suggests brittle intergranular fracture, though morphology of the fracture surface is not clear due to oxidation. Results of the examination of creep fracture modes are summarized in Fig. 3. There are two types of fracture modes; transgranular and intergranular fracture. The intergranular fracture region is further divided into two domains according to the location of cavity formation. Secondary electron (SE) and back scatted electron (BSE) images of the same area are shown in Fig. 4. Laves phase with a higher average atomic number is easily identified by its bright contrast in the BSE image. The other particles seen in the micrographs are mostly M23 C6 carbides. At high stresses in the intergranular fracture region (Fig. 4(a) and (b)), cavities are formed at the triple grain junction at which a
where d0 and d are the average diameter of particles measured before and after high temperature exposure, respectively, t the exposure time, and Kd is a material constant proportional to diffusivity. When d0 is negligibly small, Eq. (2) can be simplified to Qd m d = K0 t exp − (3) RT
(2)
where Qd is the activation energy for particle coarsening, and K0 is a constant. The exponent m is characteristic of the coarsening mechanisms of precipitates; for example, m = 3 for the coarsening controlled by lattice diffusion, and 4 for grain boundary diffusion [13]. It is evident in Fig. 5 that there are two stages F and C with different exponents m on the log d versus log t curves. Apparent activation energies Qd and exponent m obtained are Qd = 348 kJ/mol and m = 7 in stage F, and Qd = 342 kJ/mol and m = 4 in stage C. Amount of W in precipitates has been reported as a function of exposure time in Ref. [15] on grade 92 steel. It
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Fig. 4. Cavity formation in specimens crept ((a) and (b)) for 26 875 h at 600 ◦ C (wedge crack), and ((c) and (d)) for 26 783 h at 650 ◦ C (isolated cavities); (a) and (c) are secondary electron images, and (b) and (d) are back scattered electron images of the same location. The arrows indicate cavities.
increases with increasing exposure time, and saturates at an equilibrium value. The upward arrows in Fig. 5 indicate the points at which 95% of W has precipitated. In stage C, W precipitation has finished and the basic assumption of Ostwald ripening is fulfilled. The Qd value of 342 kJ/mol is close to that for lattice diffusion of W in 9Cr steel [16]. Therefore, it can be concluded that lattice diffusion controls the growth of Laves phase in stage C, though m = 4 is larger than the expected value of 3 for this mechanism. In stage F, nucleation and coarsening of Laves phase proceeds simultaneously. Since the W precipitation delays at lower temperature [15], the apparent activation energy for coarsening, 348 kJ/mol in stage F, should be larger than the true value. This in turn suggests grain boundaries as diffusion path in this stage.
Fig. 5. Coarsening of Laves phase in gage and grip portions. The dash-dot line is the boundary between regions F and C with different coarsening characteristics.
Segregation of W was analyzed by EDS attached to FETEM in a grip portion after 646 h and 8955 h exposure at 650 ◦ C, and the results are given in Fig. 6. The two exposure durations are in stage F (646 h), at which the supersaturated W is still present, and stage C (8955 h), after finishing the precipitation of Laves phase, respectively. In stage F, the W peaks measured at grain boundary is obviously higher than those from the matrix. This result suggests that a part of supersaturated W segregates to grain boundaries. The segregated W is consumed to form precipitates of Laves phase on grain boundaries. The path of W flow for the precipitation can be grain boundaries as speculated. The W peaks almost disappear in stage C after finishing the precipitation. In
Fig. 6. EDS analyses of matrix, Laves phase and grain boundary after 646 h exposure at 650 ◦ C (stage F in Fig. 5), and of grain boundary after 8955 h at 650 ◦ C (stage C in Fig. 5).
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this stage, Ostwald ripening of Laves phase probably proceeds through lattice diffusion, as suggested by the activation energy for the ripening. 3.4. Cause of fracture mechanism change The stress exponent n for rupture life of the grade 92 steel decreases in long-term region L. The dash-dot line in Fig. 1 is the boundary between the two regions H and L. The breakdown of creep strength in region L is typical of advanced high Cr ferritic steels containing W [2,4]. Transition from ductile transgranular fracture to brittle intergranular fracture occurs in the grade 92 steel. The dashed line in Fig. 1(a) represents the boundary of the ductile to brittle transition. The two boundaries of the breakdown and the fracture mode change coincide fairly well with each other, thereby indicating a close correlation between the change in fracture mechanism and the breakdown of creep strength. Similar good correlations have been reported on grade 122 steel [2] and austenitic stainless steel [17]. The fracture mechanism maps published by Ashby et al. [12] have clearly demonstrated the breakdown at the boundaries between different fracture mechanisms. Formation and coalescence of creep cavities along grain boundaries brings about the brittle intergranular fracture. The cavities are usually attached to coarse particles of Laves phase (see Fig. 4), suggesting that the precipitation and coarsening of Laves phase are responsible for the intergranular fracture. The downward arrows in Fig. 5 indicate the points of ductile to brittle transition. The points correspond to an average diameter of Laves phase of 130 nm (127 nm at 600 ◦ C and 140 nm at 650 ◦ C). This fact together with the nucleation site of cavities suggests that coarsening of Laves phase particles over a critical size triggers the brittle grain boundary fracture in the steel studied. Creep cavities are nucleated on coarse particles of Laves phase in brittle intergranular fracture region of grade 122 steel also [2]. This agreement suggests that the major cause of the ductile to brittle transition is the same in grade 92 and 122 steels for boiler application. 4. Detection of onset of intergranular fracture Advanced high Cr ferritic steels with W have been widely used in electric power plants [18]. Accurate evaluation of their creep lives is necessary for safe and efficient operation of those power plants, and detection of the breakdown of creep strength is of crucial importance. Coarsening of Laves phase particles over a critical size triggers the transition from ductile to brittle fracture, and the transition of fracture mechanism brings about the breakdown. Therefore, measurements of cavity density and particles size of Laves phase are essential for the detection. The measurements can be done with replica sampling. However, it is more practical if the transition can be detected more easily, for example, by hardness testing. Fig. 7 shows the changes of Vickers hardness as a function of exposure time. The hardness was measured in gage and grip portions after creep rupture. The solid symbols indicate the presence of Laves phase particles after the exposure. The high Cr ferritic
Fig. 7. Change of Vickers hardness in (a) gage and (b) grip portions of the steel studied. Downward arrows: Laves phase have grown to 130 nm. Upward arrows: Boundary of ductile to brittle transition. Open arrows: 0.5 mass% W has precipitated.
steel has a tempered martensitic lath structure, and recovery of the dislocation substructure primarily brings about the decrease in hardness after creep deformation in the gage portion [19]. However, the softening curve does not provide any signal of the breakdown and the ductile to brittle transition. As seen in Fig. 7(b), the hardness changes even in the grip portion, and the curves reveal two stages. The hardness remains almost constant during the short-term exposure, and decreases gradually after the long-term exposure. The solid downward arrows in Fig. 7(b) indicate the points at which Laves phase particles have grown to the average size of 130 nm. The upward arrows correspond to the ductile to brittle transition. These points well accord with the drop of hardness in grip. So, the drop of hardness can be a useful signal of the particle coarsening and the consequent ductile to brittle transition. The drop of hardness can reasonably be correlated to coarsening of Laves phases as follows. The steel was tempered at 780 ◦ C for 3 h. In grip portion without creep deformation, recovery of the tempered martensitic lath structure does not proceed evidently, since the recovery is a strain assisted process [4]. Loss of solid solution hardening and/or precipitation hardening should be the cause of the hardness decrease in the grip portion. Coarsening of M23 C6 carbides and MX carbonitrides starts at the beginning of creep, and cannot explain the late onset of the decrease in hardness after 103 h at 650 ◦ C. Laves phase precipitates during the high temperature exposure at the expense of W in the matrix and on grain boundaries. Removal of W should result in decrease of hardness. The downward open arrows in Fig. 7(b) indicates the points at which 0.5 mass% W has precipitated [15]. However, the drop of hardness does not start at those points. Precipitation hardening due to fine Laves phase particles can compensate the loss of solid solution hardening in the early stage of Laves phase precipitation as reported
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in literatures [20]. This is why the hardness remains constant after the precipitation of Laves phase. Laves phase grows faster than M23 C6 and MX particles. The hardness starts to drop when Laves phase particles have grown and failed to compensate the loss of solution hardening. Therefore, the hardness drop in grip portion coincides with the coarsening of Laves phase particles. 5. Summary 1. Stress exponent for rupture life decreases from 17 to 8.4 in long-term creep, confirming the typical breakdown of creep strength in ASTM grade 92 steel. Transition from ductile transgranular fracture to brittle intergranular fracture is the major cause of the breakdown. 2. Most of cavities are nucleated at coarse Laves phase particles on grain boundaries. 3. Cavity formation at coarse particles of Laves phase triggers the brittle intergranular fracture. The ductile to brittle transition occurred when Laves phase has reached an average diameter of 130 nm. 4. One can non-destructively detect the coarsening of Laves phase particles and the consequent brittle fracture by means of hardness testing during high temperature exposure without stress. References [1] R. Viswanathan, Damage Mechanisms and Life Assessment of High Temperature Components, ASM International, Metals Park, OH, 1989 (Chapter 3).
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[2] K. Maruyama, J.S. Lee, in: I.A. Shibli, et al. (Eds.), Creep and Fracture in High Temperature Components—Design and Life Assessment Issues, DEStech Publications, Lancaster, PA, 2005, p. 372. [3] H. Cerjak, P. Hofer, B. Schaffernak, ISIJ Int. 39 (1999) 874. [4] K. Maruyama, K. Sawada, J. Koike, ISIJ Int. 41 (2001) 641. [5] K. Kimura, H. Kushima, K. Sawada, in: A. Strang, et al. (Eds.), Parsons 2003 Engineering Issues in Turbine Machinery, Power Plant and Renewables, Institute of Materials, Minerals and Mining, London, 2003, p. 443. [6] H. Danielson, J. Hald, Proceedings of the fourth EPRI International Conference on Advances in Materials Technology for Fossil Power Plants, EPRI, Palo Alto, CA, 2004 (CD-ROM). [7] R. Agamennone, W. Blum, C. Gupta, J.K. Chakravartty, Acta Mater. (2006), in press. [8] A. Strang, V. Vodarek, Mater. Sci. Technol. 12 (1996) 552. [9] H. Semba, F. Abe, Proceedings of the Symposium on USC Steels for Fossil Power Plants 2005, Beijing, China, 2005, p. 157. [10] H. Kushima, K. Kimura, F. Abe, J. Iron Steel Inst. Jpn. 85 (1999) 841. [11] T. Muraki, Y. Hasegawa, M. Ohgami, Key Eng. Mater. 171–174 (2000) 499. [12] M.F. Ashby, C. Gandihi, D.M.R. Taplin, Acta Metall. 27 (1979) 699. [13] I.M. Lifshitz, V.V. Slyozov, J. Phys. Chem. Solids 19 (1961) 35. [14] C. Wagner, Z. Electrochem. 65 (1961) 581. [15] J. Hald, in: E. Metcalfe (Ed.), Proceedings of the EPRI/National Power Conference on the New Steels for Advanced Plant up to 620 ◦ C, EPRI, Palo Alto, CA, 1995, p. 152. [16] J. Cermak, J. Ruzickova, A. Pokorna, Acta Mater. 12 (1996) 5007. [17] H. Nakakuki, K. Maruyama, H. Oikawa, K. Yagi, J. Iron Steel Inst. Jpn. 81 (1995) 220. [18] F. Masuyama, in: J. Lecomte-Beckers, et al. (Eds.), Materials for Advanced Power Engineering 1998, Forschungszentrum Julich GmbH, Julich, 1998, p. 1807. [19] K. Sawada, K. Maruyama, Y. Hasegawa, T. Muraki, Key Eng. Mater. 171–174 (2000) 109. [20] F. Abe, Mater. Sci. Eng. A 319–321 (2001) 770.