K. B. YOO, 1) J. H. KIM 2) and N. H. HEO 1) mechanism is clarified.

Transcription

1 , pp Impurities Segregation to Grain Boundary Carbide Interfaces and Grain Boundaries and the Mechanism of Elevated Temperature Intergranular Cracking in Heat-resistant Steel K. B. YOO, 1) J. H. KIM 2) and N. H. HEO 1) 1) KEPCO Research Institute, 65 Munji-Ro, Yusung-Gu, Daejeon Republic of Korea. 2) Department of Mechanical Design Engineering, Chungnam National University, Daeduk Science Town, Daejeon Republic of Korea. (Received on June 29, 2010; accepted on July 29, 2010) It is shown that the time to elevated temperature intergranular failure in a heat-resistant steel is expressed by t t 0 s n exp(q/rt) where n is the stress exponent, Q the activation enthalpy, and t 0 the proportional constant. It is also shown that the segregation concentration of impurities is markedly higher in the dimples of grain boundary area than at the smooth grain boundaries and so the dimples observed usually at reheat or stress relief cracked surfaces are not the micro-ductile fracture areas but the carbide/matrix interfaces at the grain boundaries. Finally, it is shown that the elevated temperature intergranular cracking results from the smooth grain boundary cracking following the cracking of the carbide/matrix interfaces at the grain boundaries. KEY WORDS: AES; iron alloys; grain boundary embrittlement; interface segregation; stress rupture. 1. Introduction Recently, a premature failure of W-modified 9Cr ferritic steel tubes or pipes took place in a power plant in Japan and the fracture type was intergranular. 1) Also, the similar failure has frequently occurred in a W-modified 2.25Cr steel of several power plants in Korea since Some explanations for the premature failure are based on the formation of Laves phase 1) at grain boundaries, the formation of recovery zone along prior austenite grain boundaries 2) and the formation of Z phase which consumes MX carbides. 3,4) It has been also proposed that the elevated temperature intergranular cracking in heat-resistant steels occurs through the development of closely spaced voids at the grain boundaries oriented normal to the tensile stress. 5 11) Although several authors have analyzed the growth of a regular array of voids on grain boundaries normal to the tensile stress and a great deal has been known about the phenomenological effects of tensile stress on grain boundary segregation and low ductility intergranular fracture, how the uniform array of voids originates and what the nature of low ductility intergranular fracture induced by the tensile stress is are still obscure ) In this study, the phenomenon of elevated temperature intergranular cracking occurring in the W-modified heat-resistant ferritic steels is reassessed through simple rupture tests and the cracking mechanism is clarified. 2. Experimental A 2.25Cr1.5W heat-resistant steel ingot of 6 kg was prepared through vacuum induction melting process. The chemical composition of the steel is shown in Table 1. The ingot was homogenized at C for 1 h and forged to 12 mm thick plates. Unnotched and cylindrical rupture test specimens with a thread-head of 11 mm in diameter and a gauge of 15 mm in length and 6 mm in diameter were machined from the plates. The specimens were solutiontreated at C for 1 h under a vacuum of about 1 Pa and water-quenched for obtaining a martensitic structure. Rupture tests were carried out using conventional creep testers in the temperature range of 550 to 700 C and in the tensile stress range of 75 to 500 MPa. The specimens to which the N-type thermocouple was attached were heated to a test temperature at C/h and then a tensile stress was imposed on the specimens with no soaking at the temperature. Fracture surfaces were examined by a scanning electron microscope (SEM). The segregation behavior was investigated with Auger electron spectroscopy (AES). AES specimens machined from the ruptured specimens were fractured after chilling for about 30 min with liquid nitrogen in a vacuum of about Pa or higher to minimize the post-fracture Table 1. The chemical composition of the steel (wt%). 1702

2 contamination. Grain boundary facets or carbide/matrix interfaces more than ten points were investigated. Peak-topeak height ratios, I/I Fe, were obtained from each differential Auger spectra and then averaged. The used peaks were Fe 703, P 120, S 152, C 271, N 379, Cr 529 and W Results and Discussion Figure 1 shows changes in time to failure with tensile stress and temperature. As shown in Figs. 1(a1) and 1(b1), the overall time to failure increased with decreasing stress and temperature. Except for the results for 75 and 100 MPa at 700 C and 400 MPa at 675 C which showed a reduction in area (RA) higher than 30%, most of fractured specimens showed nearly zero RA. As shown in Fig. 1(a2) produced from Fig. 1(a1), the time to failure in a log scale was inversely proportional to the tensile stress in a log scale, and the gradient was the same, regardless of temperature. In Fig. 1(b2) which is obtained from Fig. 1(b1), the time to failure in a log scale was proportional to Q/RT, and the gradient was also the same, regardless of tensile stress. As a result, the time to failure, t, can be expressed by t 0 s n exp(q/rt) where t 0 is a proportional constant, n the stress exponent, Q the activation enthalpy and R the universal gas constant. As shown in Table 2, n and Q were evaluated from Figs. 1(a2) and 1(b2), and the averaged t 0 was obtained using the evaluated n and Q and the experimental results of Figs. 1(a1) and 1(a2). That is, the time to failure decreases as n increases and Q decreases. As shown in Fig. 2, the calculated results can be obtained using the constant values of Table 1. Here, the lines denote the calculated results and the symbols the experimental ones of Figs. 1(a1) and 1(b1). The calculated results were consistent with the experimental ones, except for the conditions with a considerable ductility. Figure 3 shows fracture surfaces of the ruptured specimen of the 350 MPa-575 C condition of Fig. 1(b1) and the corresponding AES specimen and the obtained AES spectra. As shown in Figs. 3(a) and 3(b), the fracture surfaces of the ruptured specimen and the AES specimen were similar. The fracture surface were composed of smooth grain boundaries and dimples which have been observed generally at reheat or stress relief cracked surfaces and considered as micro-ductile fracture areas. 15) Fine particles were also observed in the dimples of Fig. 3(b). As shown in Fig. Table 2. Evaluated values of constants in t t 0 s n exp(q/rt). Fig. 1. Changes in time to failure with tensile stress and temperature: (a1) and (b1) the experimental results; (a2) and (b2) produced from (a1) and (b1). Fig. 2. The consistency between experimental and calculated results: (a) changes in time to failure with tensile stress and (b) changes in time to failure with temperature. 1703

3 ISIJ International, Vol. 50 (2010), No. 11 Fig. 3. Fracture surfaces of the ruptured specimen of the 350 MPa-575 C condition of Fig. 1(b1) (a), fracture surfaces of the corresponding AES specimen (b) and the obtained AES spectra (c). In Fig. 3(c), the square and the circle denote the dimples and the smooth grain boundaries, respectively. Fig. 4. The difference in segregation concentration of the impurities between the carbide/matrix interfaces at the grain boundaries and the smooth grain boundaries: (a) 300 and (b) 500 MPa. 3(c), the strong phosphorus segregation was observed in the dimples. The segregation concentration of phosphorus was markedly higher in the dimples than at the smooth grain boundaries. In the other dimples, the noticeable sulfur segregation was additionally observed together with the strong phosphorus segregation. Such a prominent segregation behavior of the impurities was also reported in a nickel-based superalloy,16) although changes in segregation concentration of the impurities with temperature and tensile test were not investigated in the research but the segregation behavior was only obtained from one rupture test condition. The markedly higher segregation concentration of phosphorus would not be obtained from the dimples if they were the micro-ductile fracture areas around the grain boundaries. The strong sulfur and chromium peaks in the dimples might arise from MnS particles and chromium carbides, but it is not clear in the present study. Changes in segregation concentration of phosphorus and sulfur with temperature and tensile stress are summarized in Fig. 4. Irrespective of temperature, the AES specimens showed a nearly saturated phosphorus segregation concentration in the dimples (Pdimple) much higher than that at the smooth grain boundaries (PG.B.). The phosphorus concentration in the dimples was higher under 500 MPa than 300 MPa, while the segregation concentration at the grain boundaries was similar. The phosphorus at the smooth grain boundaries and in the dimples increased slightly with decreasing temperature. On the other hand, the segregation concentration of sulfur under 300 MPa increased with increasing temperature, but the segregation concentration under 500 MPa was overall negligible. From the prominent segregation concentration of the impurities in the dimples of Figs. 3 and 4, the dimples observed in the present study or at intergranular reheat cracked surfaces are not the micro-ductile fracture areas15) but the carbide/matrix interfaces formed at the grain boundaries. 1704

4 Fig. 5. The dependence of theoretical matrix, yield and grain boundary strengths on temperature (a) and a schematic diagram for understanding the mechanism of elevated temperature intergranular cracking (b). Considering the free surface in which the distance between atoms in the broken surface and the broken surface pair is infinite, the surface energy is much higher than the grain boundary energy ) This is because the grain boundary may be treated as a kind of interface at which the average distance between the outermost atoms of two grains is longer than the equilibrium atomic distance. Therefore, the more active segregation of impurities occurs at the free surface. Furthermore, because the incoherent carbide/matrix interface at the grain boundaries may be considered to be a quasi-free surface at which the atomic bond between the outermost atoms of the carbide surface and the neighboring grain is relatively weak, the active segregation of the impurities can occur also at the carbide/matrix interfaces of Figs. 3 and 4. In order to understand the mechanism of elevated temperature intergranular cracking, the dependence of theoretical matrix, yield and grain boundary strengths on temperature needs to be introduced. Here, the theoretical matrix strength is defined as a strength that fractures a metal by breaking bonds between atoms which face each other across the plane of fracture without plastic flow of any kind. When a pure body-centered cubic metal is heated from 0 K, the theoretical matrix strength s th (T) decreases gradually as the equilibrium atomic distance increases. 20) The dependence of the grain boundary strength s gb (T) on temperature should also show a similar behavior to that of s th (T). However, the grain boundary strength s gb (T) is lower than s th (T), due to the longer distance between the outermost atoms of two grains at the grain boundaries than the equilibrium atomic distance. Figure 5(a) shows an oversimplified schematic diagram for the dependence of the theoretical matrix and grain boundary strengths on temperature. As shown also in Fig. 5(a), the discrepancies between the theoretical strength and the observed yield strength s y (T) in the regions I, II and III are mainly related to the easiness in movement of dislocations, 21 24) the decrease in shear modulus with increasing temperature 25) and the micro-creep, 24) respectively. Because there exist two equicohesive temperatures 26,27) T e1 and T e2 (T e1 T e2 ) where the yield strength and the grain boundary strength are the same, intergranular cracking can occur only in the regions W 1 below T e1 and W 2 above T e2. The increase in yield strength and the decrease in grain boundary strength arising from the segregation of impurities to the grain boundaries, which result in an increase in T e1 and a concurrent decrease in T e2, raise the susceptibility to the intergranular cracking. Finally, the mechanism of elevated temperature intergranular cracking in the present heat-resistant ferritic alloy may be understood with Fig. 5(b) which is constructed on the basis of the present experimental results and Fig. 5(a). Here, the elevated temperature is fixed at T 0 and the tensile stress is given to s 0. Based on the present study, the grain boundary area is composed of the smooth grain boundaries and the carbide/matrix interfaces that an active segregation of the impurities is observed. During the rupture test, the yield strength continues to decrease with increasing time due to the tempering phenomenon. The strengths of the carbide/matrix interfaces and the grain boundaries also decrease due to the embrittlement by the impurities segregation. Because the carbide/matrix interface strength decreases more rapidly due to the active segregation of phosphorus to the carbide/matrix interfaces, the carbide/matrix interfaces should be first fractured when the carbide/matrix interface strength s c at T 0 drops to s 0. In this case, the tensile stress that the remaining grain boundaries experience after cracking of the carbide/matrix interfaces will increase from s 0 to s 0 due to the decrease in grain boundary area. The subsequent cracking at the carbide-free grain boundaries is possible if the equicohesive temperature T e2 should be lower than T 0 and s 0 is lower than the yield strength. Under this condition, the intergranular cracking occurs eventually when the grain boundary strength s gb drops to the stress s Conclusions The time to elevated temperature intergranular failure is expressed by t t 0 s n exp(q/rt) where t 0 is the proportional constant of (s), n the stress exponent of 4.5, Q the activation enthalpy of 84.6 kcal/mol and R the universal gas constant. The dimples observed usually at reheat or stress relief cracked surfaces are not the micro-ductile fracture areas but the grain boundary carbide interfaces. Due to the stronger segregation behavior of impurities to the carbide/matrix interfaces at the grain boundaries, the carbide/matrix interfaces are first fractured when the interface strength drops to the given tensile stress. The final intergranular cracking occurs when the grain boundary strength drops to the tensile stress increased after the carbide/matrix interface cracking. REFERENCES 1) T. Muraki, Y. Hasegawa and M. Ohgami: Key Eng. Mater., (2000), ) H. Kushima, K. Kimura and F. Abe: Testu-to-Hagané, 85 (1999), ) H. Danielson and J. Hald: Proc. of the Fourth EPRI Int. Conf. on Advances in Materials Technology for Fossil Power Plants, EPRI, Palo Alto, CA, (2007), ) A. Strang, J. Cawley and G. W. Greenwood: Microstructure of High Temperature Materials, No. 2, The Institute of Materials, (1998), ) C. A. Hippsley, J. F. Knott and B. C. Edwards: Acta Metall., 28 (1980), ) D. Hull and D. E. Rimmer: Philos. Mag. A, 4 (1959), ) M. V. Speight and J. E. Harris: Met. Sci., 1 (1967), 83. 8) R. Raj and M. F. Ashby: Acta Metall., 23 (1975), ) T. J. Chuang, K. I. Kagawa, J. R. Rice and L. B. Sills: Acta Metall., 27 (1979),

5 10) M. W. D. Van Der Brug, E. Van Der Giessen and R. C. Brouwer: Acta Metall., 44 (1996), ) R. Raj: Acta Metall., 26 (1978), ) J. Shin and C. J. McMahon: Met. Sci., 32 (1984), ) R. D. K. Misra: Acta Metall., 44 (1996), ) D. Bika, J. Pfaendtner, M. Menyhard and C. J. McMahon: Acta Metall., 43 (1995), ) N. Bailey: Weldability of Ferritic Steels, The Materials Information Society, Materials Park, Ohio 44073, USA, (1994), ) J. X. Dong, X. B. Liu, X. S. Xie and R. G. Thompson: Acta Metall. Sinica, 10 (1997), ) E. D. Hondros and M. P. Seah: Int. Met. Rev., 222 (1977), ) C. Lea and M. P. Seah: Philos. Mag., 35 (1977), ) M. P. Seah: Acta Metall., 28 (1980), ) N. T. Barrett: The Principles of Engineering Materials, Prentice-Hall INC., Eaglewood Cliffs, New Jersey, (1973), ) J. Weertman and J. R. Weertman: Elementary Dislocation Theory, MaCmillan Company Collier-MaCmillan Limited, London, (1971). 22) W. C. Leslie: Metall. Trans., 3 (1972), 5. 23) H. Conrad: Iron and Its Dilute Solid Solutions, Interscience, New York, (1961), ) W. C. Leslie: The Physical Metallurgy of Steels, McGraw-hill International Book Company, International Student Edition, (1982). 25) G. R. Speich, A. S. Schwoeble and W. C. Leslie: Metall. Trans., 3 (1972), ) Z. Jeffries: Trans. AIME, 60 (1919), ) N. J. Grant: Utilization of Heat Resistant Alloys, ASM, Ohio, (1954),