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Microstructural evolution and change in hardness in type 304H stainless steel during long-term creep
Materials Science and Engineering A319– 321 (2001) 788– 791 www.elsevier.com/locate/msea
Microstructural evolution and change in hardness in type 304H stainless steel during long-term creep H. Tanaka *, M. Murata, F. Abe, H. Irie Frontier Research Center for Structural Materials, National Research Institute for Metals, 1 -2 -1 Sengen, Tsukuba, Ibaraki 305 -0047, Japan
Abstract The microstructural evolution and the change in hardness have been investigated for 18Cr– 8Ni (type 304H) stainless steel during long-term creep. Creep and creep-rupture tests were carried out at temperatures between 550 (823) and 750°C (1023 K) for up to 180 000 h. The hardening behavior during creep depends on the stress level, as well as the precipitation of M23C6 carbides and s phase. At a high stress of 177 MPa, the hardening during creep is much larger than the age hardening, indicating that the hardening during creep is mainly caused by the strain hardening due to creep deformation. At a later stage of creep, the softening occurs due to the recovery of excess dislocations, which becomes more significant with decreasing stress and increasing test duration. The strain hardening disappears with decreasing stress level and increasing test duration. At a low stress of 61 MPa or less, the hardening during creep can be approximately given by the age hardening under no stress, except for the final stage of creep. © 2001 Elsevier Science B.V. All rights reserved. Keywords: Microstructural evolution; Type 304H; Long-term creep
1. Introduction Austenitic stainless steels, such as type 304 (l8Cr – 8Ni), 316 (l8Cr – l2Ni – Mo), 321 (18Cr – 10Ni –Ti) and 347 (18Cr –12Ni – Nb) steels, are widely used as hightemperature components, such as boilers, superheaters and chemical reactors, which require good mechanical properties and corrosion resistance at temperatures up to 650 –700°C. It is well known that long-term exposure of the steels to stress at these temperatures can cause microstructure evolution and creep deformation leading ultimately to creep-rupture. Of the steels, the type 304 steel exhibits the simplest microstructure where only M23C6 carbides and s phase precipitate at high temperatures, while the microstructure of the other steels are complicated [1]. Several investigations have dealt with the microstructure evolution in the type 304 steel during thermal aging and creep [2,3]. But these studies are limited to specimens tested in periods not exceeding 60 000 h, while high-temperature components are usually operating for long periods exceeding 100 000 h. At * Corresponding author. Tel.: + 81-298-592219; fax: +81-298592201. E-mail address: [email protected] (H. Tanaka).
present, little is known about the microstructure evolution during long-term creep for up to 100 000 h or more. The present authors and co-workers have investigated comprehensively the microstructure evolution during creep and its effect on long-term creep strength for type 304H steel, where the term H means high carbon concentration, using specimens tested in the NRIM Creep data sheet project [1,4 –6]. A number of micrographs for the type 304H steel were recently published in ‘Metallographic Atlas of Long-term Crept Materials’ [7]. In the present research, the relationship between the microstructure evolution and the change in hardness has been investigated for the type 304H steel during creep for up to 180 000 h at high temperature. The measurement of hardness is frequently used for estimating materials degradation and remaining life for high-temperature components.
2. Experimental procedure The material used was the commercially produced heat ABE of 18Cr –8Ni steel (304H TB) in the NRIM Creep Data Sheet No. 4B [8], where the creep rupture data were presented for nine heats of 18Cr –8Ni steel.
0921-5093/01/$ - see front matter © 2001 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 0 1 2 - 7
H. Tanaka et al. / Materials Science and Engineering A319–321 (2001) 788–791
The chemical compositions of the steel used were 0.07C, 0.55Si, l.46Mn, 0.3P, 0.006S, 9.57Ni, 18.95Cr, 0.04Mo, 0.07Cu, 0.062Ti, 0.014Al, 0.0018B, 0.0278 (mass percentage) N and Fe, balance. The creep specimens, having a geometry of 6 mm in diameter and 30 mm in gauge length, were taken longitudinally from the middle of the wall thickness of the as-received boiler tube. The tube had a size of 50.8 mm in outer diameter and 8 mm in wall thickness and was already solution annealed. Creep and creep rupture tests were carried out for up to about 180 000 h at temperatures between 550 (823) and 750°C (1023 K). The longitudinal crosssection of the gauge and head (or grip) portions of the specimens was observed using optical, scanning and transmission electron microscopes (SEM and TEM). The head portion can be regarded as an unstressed portion. The Vickers hardness was measured at a load of 49 N (5 kgf mm − 2). The indentation size at this load
789
was 200–240 mm which corresponds to two or three times the average grain size.
3. Results and discussion
3.1. Creep rupture strength Fig. 1 shows the stress versus time to rupture data. The solid curves are drawn on the based of the timetemperature parameter method of Manson-Haferd [8]. The creep fracture modes for the 304H steel are divided into regimes of transgranular fracture (denoted by T) and three types of intergranular fracture; the wedgetype cracking (W), the creep cavitation associated with M23C6 carbides at grain boundaries (C) and the s/ matrix interface cracking along grain boundaries (s) [4,5]. The present results suggest that the creep fracture modes at long times above about 10 000 h are closely connected with the precipitation behavior of M23C6 carbides and s phase.
3.2. Change in hardness during aging and creep
Fig. 1. Stress versus time to rupture for the heat ABE of type 304H steel.
Fig. 2. Vickers hardness of (a) head and (b) gauge portions of creep-ruptured specimens, as a function of duration of creep rupture testing.
Fig. 2 shows the Vickers hardness for both the head and gauge portions of creep-ruptured specimens as a function of duration of creep rupture testing. The Vickers hardness in the as-received condition was 160. The specimen head portion under no stress exhibits the two step age hardening, short-term age hardening at times less than 1000 h and long-term age hardening at times above 10 000 h. The short-term age hardening occurs substantially at short times less than 100 h and the hardening DH = H−H0, where H and H0 are the hardness at time t and in the as-received condition, respectively, is relatively small 10–15. The long-term age hardening becomes more significant with time above 10 000 h at around 600°C and is much larger than the short-term age hardening. Fig. 3 shows the microstructure evolution with time in the specimen head portion at 650°C. The TEM observations show that the precipitation of M23C6 carbides started to occur from short times less than 1 h but that of the s phase needed long times above about 10 000 h at around 600–650°C. Therefore, the short-term and long-term age hardening in Fig. 2a results from the precipitation of M23C6 carbides and s phase, respectively. In the gauge portion of creep-ruptured specimens, the hardness decreases with time at temperatures except for 600°C where there is a tendency to increase again at long times above 10 000 h. The change in hardness with time for the gauge portion results from the change in dislocation density produced by creep deformation as well as the precipitation hardening due to the M23C6 carbides and s phase described above. It should be noted that the solid lines in Fig. 2b are connecting the
790
H. Tanaka et al. / Materials Science and Engineering A319–321 (2001) 788–791
Fig. 3. Change in TEM microstructure during thermal aging at 650°C. (a) As-received, (b) 71.9 h, (c) 13 718.5 h and (d) 100 491.4 h.
suggests that the solid lines in Fig. 2b cannot represent the test duration dependence alone but that they also involve the effect of stress level. In order to exclude a possible influence of stress level, the change in hardness was measured as a function of time during creep at a constant load condition. The creep tests were carried out at 650°C and at three different stress levels, 177, 118 and 61 MPa at which the time to rupture tr was 71.9, 2621.3 and 100 491.4 h, respectively. Fig. 4 shows the respective creep curves. The creep tests were interrupted at times shown by the arrows in Fig. 4 and then the hardness was measured. Fig. 5 shows the change in hardness during creep as functions of time (Fig. 5a) and normalized time t/tr (Fig. 5b). In this figure, the hardness in the specimen head portion under no stress is also shown by the dotted line for comparison. At a high stress of 177 MPa, the hardness increases for up to 50 h, corresponding to t/tr= 0.7, then decreases slightly and again increases just before creep-rupture. The hardness in the specimen gauge portion is much larger than that in the specimen head portion under no stress over a whole range of test duration, indicating that the hardening during creep is mainly caused by the strain hardening. The precipitation of M23C6 carbides occurs during creep but the precipitation hardening DH due to M23C6 carbides is relatively small as described above. The TEM observations showed that the dislocation density increased for up to 50 h (t/tr= 0.7) at which the accumulated creep strain is 0.16 as can be seen from
Fig. 4. Creep curves of type 304h steel at 177, 118 and 61 MPa at 650°C.
data points for the creep-ruptured specimens which were tested at different stress levels as shown in Fig. 1. In general, resultant dislocation density and resultant dislocation arrangements in the specimens are strongly influenced by stress level in creep rupture testing. This
Fig. 5. Vickers hardness during creep as functions of (a) time and (b) normalized time t/tr at 650°C.
H. Tanaka et al. / Materials Science and Engineering A319–321 (2001) 788–791
Fig. 4. At 65 h (t/tr =0.9), the matrix having low density of dislocations was surrounded by walls of high dislocation density, indicating that rearrangement of excess dislocations had occurred during creep. In the final stage of creep above 65 h, an increase in dislocation density was observed again in the matrix. At 118 MPa, there is not a large difference in hardness between the specimen gauge and head portions for up to about 1000 h but the hardness in the specimen gauge portion increases rapidly with time above 1000 h followed by softening after reaching a maximum at about t/tr= 0.8. At short times less than 1000 h, the hardening during creep is determined mainly by the precipitation hardening due to the M23C6 carbides but the strain hardening is not large, because the creep strain is relatively small 0.04 or less as can be seen in Fig. 4. The rapid increase in hardness above about 1000 h correlates with the rapid increase in creep strain, indicating strain hardening. The softening in the final stage of creep is considered to result from the recovery or rearrangement of excess dislocations and the coarsening of precipitates. At a low stress of 61 MPa, the hardening during creep is approximately the same as the age hardening for almost whole range of test duration, except for the final stage of creep where the softening occurs as at 118 MPa. This suggests that the hardening during creep is determined substantially by the precipitation hardening due to the M23C6 carbides at short times and due to the s phase at long times above 10 000 h. The strain hardening is much smaller at 61 MPa than at 177 and 118 MPa. The effects of recovery or rearrangement of excess dislocations and of coarsening of precipitates become more significant with decreasing stress, because the test duration becomes longer and the creep deformation rate becomes lower with decreasing stress. The present results indicate that the difference in hardness between the specimen gauge and head portions disappears with decreasing stress level and that the hardening during creep under stresses less than 61 MPa can be approximately given by the age hardening under no stress, except for the final stage of creep. In Fig. 2b, the hardness in the creep-ruptured specimens decreases with increasing temperature except for 600°C. This is also correlated with an increased effect of recovery of excess dislocations with increasing temperature, depending on self-diffusion rates. The increase in hardness at long times above 10 000 h at 600°C is caused by the large precipitation hardening due to the s phase.
3.3. Application to e6aluation of long-term operating components Most of high-temperature plants using type 304 austenitic steel are being operated for long times greater
791
than 100 000 h, indicating a low stress condition presumably much less than 61 MPa at 650°C. Therefore, for estimating materials degradation or remaining life for type 304 austenitic steel components, the hardness can be approximated by the age hardening under no stress and the results shown in Fig. 2a can be used as reference or standard data. By comparing the hardness of the components with the data shown in Fig. 2a and by observing microstructure, we can also estimate the operating conditions, such as operating temperature and time, for the components.
4. Conclusions During thermal aging under no stress, the precipitation of M23C6 carbides and s phase causes the age hardening at short times below 1000 h and at long times above about 10 000 h, respectively, at around 650°C. The hardening behavior during creep at 650°C depends on stress levels as well as the precipitation of M23C6 carbides and s phase. At a high stress of 177 MPa, the hardness increases for up to t/tr=0.7, then decreases slightly and again increases just before creeprupture. The hardening during creep is much larger than the age hardening, indicating that the hardening during creep is mainly caused by the strain hardening. The strain hardening disappears with decreasing stress level and increasing test duration. For estimating materials degradation or remaining life for type 304 austenitic steel components, which are usually operated under stresses presumably less than 61 MPa, the hardness can be approximated by the age hardening under no stress, except for the final stage just before creep-rupture.
References [1] H. Tanaka, M. Murata, F. Abe, K. Yagi, Mater. Sci. Eng. A234 – 236 (1997) 1049. [2] Y. Minami, H. Kimura, Y Ihara, Mater. Sci. Technol. 2 (1986) 795. [3] V.A. Biss, V.K. Sikka, Metall. Trans. 12A (1981) 1360. [4] N. Shinya, J. Kyono, H. Tanaka, M. Murata, S. Yokoi, TetsuTo-Hagane 69 (1983) 1668 in Japanese. [5] H. Tanaka, M. Murata, M. Kaise, N. Shinya, Tetsu-To-Hagane 74 (1988) 2009 in Japanese. [6] M. Murata, H. Tanaka, E. Abe, H. Irie, Proceedings of the Eighth. International Conference on Creep and Fracture of Engineering Materials and Structures, Key Eng. Mater. 171 –174 (1999) 513. [7] National Research Institute for Metals Creep Data Sheet, Metallographic Atlas of Long-Term Crept Materials, National Research Institute for Metals, Japan, No. M-1 (1999). [8] National Research Institute for Metals Creep Data Sheet, National Research Institute for Metals, Japan, No. 4B (1986).
Microstructural evolution and change in hardness in type 304H stainless steel during long-term creep H. Tanaka *, M. Murata, F. Abe, H. Irie Frontier Research Center for Structural Materials, National Research Institute for Metals, 1 -2 -1 Sengen, Tsukuba, Ibaraki 305 -0047, Japan
Abstract The microstructural evolution and the change in hardness have been investigated for 18Cr– 8Ni (type 304H) stainless steel during long-term creep. Creep and creep-rupture tests were carried out at temperatures between 550 (823) and 750°C (1023 K) for up to 180 000 h. The hardening behavior during creep depends on the stress level, as well as the precipitation of M23C6 carbides and s phase. At a high stress of 177 MPa, the hardening during creep is much larger than the age hardening, indicating that the hardening during creep is mainly caused by the strain hardening due to creep deformation. At a later stage of creep, the softening occurs due to the recovery of excess dislocations, which becomes more significant with decreasing stress and increasing test duration. The strain hardening disappears with decreasing stress level and increasing test duration. At a low stress of 61 MPa or less, the hardening during creep can be approximately given by the age hardening under no stress, except for the final stage of creep. © 2001 Elsevier Science B.V. All rights reserved. Keywords: Microstructural evolution; Type 304H; Long-term creep
1. Introduction Austenitic stainless steels, such as type 304 (l8Cr – 8Ni), 316 (l8Cr – l2Ni – Mo), 321 (18Cr – 10Ni –Ti) and 347 (18Cr –12Ni – Nb) steels, are widely used as hightemperature components, such as boilers, superheaters and chemical reactors, which require good mechanical properties and corrosion resistance at temperatures up to 650 –700°C. It is well known that long-term exposure of the steels to stress at these temperatures can cause microstructure evolution and creep deformation leading ultimately to creep-rupture. Of the steels, the type 304 steel exhibits the simplest microstructure where only M23C6 carbides and s phase precipitate at high temperatures, while the microstructure of the other steels are complicated [1]. Several investigations have dealt with the microstructure evolution in the type 304 steel during thermal aging and creep [2,3]. But these studies are limited to specimens tested in periods not exceeding 60 000 h, while high-temperature components are usually operating for long periods exceeding 100 000 h. At * Corresponding author. Tel.: + 81-298-592219; fax: +81-298592201. E-mail address: [email protected] (H. Tanaka).
present, little is known about the microstructure evolution during long-term creep for up to 100 000 h or more. The present authors and co-workers have investigated comprehensively the microstructure evolution during creep and its effect on long-term creep strength for type 304H steel, where the term H means high carbon concentration, using specimens tested in the NRIM Creep data sheet project [1,4 –6]. A number of micrographs for the type 304H steel were recently published in ‘Metallographic Atlas of Long-term Crept Materials’ [7]. In the present research, the relationship between the microstructure evolution and the change in hardness has been investigated for the type 304H steel during creep for up to 180 000 h at high temperature. The measurement of hardness is frequently used for estimating materials degradation and remaining life for high-temperature components.
2. Experimental procedure The material used was the commercially produced heat ABE of 18Cr –8Ni steel (304H TB) in the NRIM Creep Data Sheet No. 4B [8], where the creep rupture data were presented for nine heats of 18Cr –8Ni steel.
0921-5093/01/$ - see front matter © 2001 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 0 1 2 - 7
H. Tanaka et al. / Materials Science and Engineering A319–321 (2001) 788–791
The chemical compositions of the steel used were 0.07C, 0.55Si, l.46Mn, 0.3P, 0.006S, 9.57Ni, 18.95Cr, 0.04Mo, 0.07Cu, 0.062Ti, 0.014Al, 0.0018B, 0.0278 (mass percentage) N and Fe, balance. The creep specimens, having a geometry of 6 mm in diameter and 30 mm in gauge length, were taken longitudinally from the middle of the wall thickness of the as-received boiler tube. The tube had a size of 50.8 mm in outer diameter and 8 mm in wall thickness and was already solution annealed. Creep and creep rupture tests were carried out for up to about 180 000 h at temperatures between 550 (823) and 750°C (1023 K). The longitudinal crosssection of the gauge and head (or grip) portions of the specimens was observed using optical, scanning and transmission electron microscopes (SEM and TEM). The head portion can be regarded as an unstressed portion. The Vickers hardness was measured at a load of 49 N (5 kgf mm − 2). The indentation size at this load
789
was 200–240 mm which corresponds to two or three times the average grain size.
3. Results and discussion
3.1. Creep rupture strength Fig. 1 shows the stress versus time to rupture data. The solid curves are drawn on the based of the timetemperature parameter method of Manson-Haferd [8]. The creep fracture modes for the 304H steel are divided into regimes of transgranular fracture (denoted by T) and three types of intergranular fracture; the wedgetype cracking (W), the creep cavitation associated with M23C6 carbides at grain boundaries (C) and the s/ matrix interface cracking along grain boundaries (s) [4,5]. The present results suggest that the creep fracture modes at long times above about 10 000 h are closely connected with the precipitation behavior of M23C6 carbides and s phase.
3.2. Change in hardness during aging and creep
Fig. 1. Stress versus time to rupture for the heat ABE of type 304H steel.
Fig. 2. Vickers hardness of (a) head and (b) gauge portions of creep-ruptured specimens, as a function of duration of creep rupture testing.
Fig. 2 shows the Vickers hardness for both the head and gauge portions of creep-ruptured specimens as a function of duration of creep rupture testing. The Vickers hardness in the as-received condition was 160. The specimen head portion under no stress exhibits the two step age hardening, short-term age hardening at times less than 1000 h and long-term age hardening at times above 10 000 h. The short-term age hardening occurs substantially at short times less than 100 h and the hardening DH = H−H0, where H and H0 are the hardness at time t and in the as-received condition, respectively, is relatively small 10–15. The long-term age hardening becomes more significant with time above 10 000 h at around 600°C and is much larger than the short-term age hardening. Fig. 3 shows the microstructure evolution with time in the specimen head portion at 650°C. The TEM observations show that the precipitation of M23C6 carbides started to occur from short times less than 1 h but that of the s phase needed long times above about 10 000 h at around 600–650°C. Therefore, the short-term and long-term age hardening in Fig. 2a results from the precipitation of M23C6 carbides and s phase, respectively. In the gauge portion of creep-ruptured specimens, the hardness decreases with time at temperatures except for 600°C where there is a tendency to increase again at long times above 10 000 h. The change in hardness with time for the gauge portion results from the change in dislocation density produced by creep deformation as well as the precipitation hardening due to the M23C6 carbides and s phase described above. It should be noted that the solid lines in Fig. 2b are connecting the
790
H. Tanaka et al. / Materials Science and Engineering A319–321 (2001) 788–791
Fig. 3. Change in TEM microstructure during thermal aging at 650°C. (a) As-received, (b) 71.9 h, (c) 13 718.5 h and (d) 100 491.4 h.
suggests that the solid lines in Fig. 2b cannot represent the test duration dependence alone but that they also involve the effect of stress level. In order to exclude a possible influence of stress level, the change in hardness was measured as a function of time during creep at a constant load condition. The creep tests were carried out at 650°C and at three different stress levels, 177, 118 and 61 MPa at which the time to rupture tr was 71.9, 2621.3 and 100 491.4 h, respectively. Fig. 4 shows the respective creep curves. The creep tests were interrupted at times shown by the arrows in Fig. 4 and then the hardness was measured. Fig. 5 shows the change in hardness during creep as functions of time (Fig. 5a) and normalized time t/tr (Fig. 5b). In this figure, the hardness in the specimen head portion under no stress is also shown by the dotted line for comparison. At a high stress of 177 MPa, the hardness increases for up to 50 h, corresponding to t/tr= 0.7, then decreases slightly and again increases just before creep-rupture. The hardness in the specimen gauge portion is much larger than that in the specimen head portion under no stress over a whole range of test duration, indicating that the hardening during creep is mainly caused by the strain hardening. The precipitation of M23C6 carbides occurs during creep but the precipitation hardening DH due to M23C6 carbides is relatively small as described above. The TEM observations showed that the dislocation density increased for up to 50 h (t/tr= 0.7) at which the accumulated creep strain is 0.16 as can be seen from
Fig. 4. Creep curves of type 304h steel at 177, 118 and 61 MPa at 650°C.
data points for the creep-ruptured specimens which were tested at different stress levels as shown in Fig. 1. In general, resultant dislocation density and resultant dislocation arrangements in the specimens are strongly influenced by stress level in creep rupture testing. This
Fig. 5. Vickers hardness during creep as functions of (a) time and (b) normalized time t/tr at 650°C.
H. Tanaka et al. / Materials Science and Engineering A319–321 (2001) 788–791
Fig. 4. At 65 h (t/tr =0.9), the matrix having low density of dislocations was surrounded by walls of high dislocation density, indicating that rearrangement of excess dislocations had occurred during creep. In the final stage of creep above 65 h, an increase in dislocation density was observed again in the matrix. At 118 MPa, there is not a large difference in hardness between the specimen gauge and head portions for up to about 1000 h but the hardness in the specimen gauge portion increases rapidly with time above 1000 h followed by softening after reaching a maximum at about t/tr= 0.8. At short times less than 1000 h, the hardening during creep is determined mainly by the precipitation hardening due to the M23C6 carbides but the strain hardening is not large, because the creep strain is relatively small 0.04 or less as can be seen in Fig. 4. The rapid increase in hardness above about 1000 h correlates with the rapid increase in creep strain, indicating strain hardening. The softening in the final stage of creep is considered to result from the recovery or rearrangement of excess dislocations and the coarsening of precipitates. At a low stress of 61 MPa, the hardening during creep is approximately the same as the age hardening for almost whole range of test duration, except for the final stage of creep where the softening occurs as at 118 MPa. This suggests that the hardening during creep is determined substantially by the precipitation hardening due to the M23C6 carbides at short times and due to the s phase at long times above 10 000 h. The strain hardening is much smaller at 61 MPa than at 177 and 118 MPa. The effects of recovery or rearrangement of excess dislocations and of coarsening of precipitates become more significant with decreasing stress, because the test duration becomes longer and the creep deformation rate becomes lower with decreasing stress. The present results indicate that the difference in hardness between the specimen gauge and head portions disappears with decreasing stress level and that the hardening during creep under stresses less than 61 MPa can be approximately given by the age hardening under no stress, except for the final stage of creep. In Fig. 2b, the hardness in the creep-ruptured specimens decreases with increasing temperature except for 600°C. This is also correlated with an increased effect of recovery of excess dislocations with increasing temperature, depending on self-diffusion rates. The increase in hardness at long times above 10 000 h at 600°C is caused by the large precipitation hardening due to the s phase.
3.3. Application to e6aluation of long-term operating components Most of high-temperature plants using type 304 austenitic steel are being operated for long times greater
791
than 100 000 h, indicating a low stress condition presumably much less than 61 MPa at 650°C. Therefore, for estimating materials degradation or remaining life for type 304 austenitic steel components, the hardness can be approximated by the age hardening under no stress and the results shown in Fig. 2a can be used as reference or standard data. By comparing the hardness of the components with the data shown in Fig. 2a and by observing microstructure, we can also estimate the operating conditions, such as operating temperature and time, for the components.
4. Conclusions During thermal aging under no stress, the precipitation of M23C6 carbides and s phase causes the age hardening at short times below 1000 h and at long times above about 10 000 h, respectively, at around 650°C. The hardening behavior during creep at 650°C depends on stress levels as well as the precipitation of M23C6 carbides and s phase. At a high stress of 177 MPa, the hardness increases for up to t/tr=0.7, then decreases slightly and again increases just before creeprupture. The hardening during creep is much larger than the age hardening, indicating that the hardening during creep is mainly caused by the strain hardening. The strain hardening disappears with decreasing stress level and increasing test duration. For estimating materials degradation or remaining life for type 304 austenitic steel components, which are usually operated under stresses presumably less than 61 MPa, the hardness can be approximated by the age hardening under no stress, except for the final stage just before creep-rupture.
References [1] H. Tanaka, M. Murata, F. Abe, K. Yagi, Mater. Sci. Eng. A234 – 236 (1997) 1049. [2] Y. Minami, H. Kimura, Y Ihara, Mater. Sci. Technol. 2 (1986) 795. [3] V.A. Biss, V.K. Sikka, Metall. Trans. 12A (1981) 1360. [4] N. Shinya, J. Kyono, H. Tanaka, M. Murata, S. Yokoi, TetsuTo-Hagane 69 (1983) 1668 in Japanese. [5] H. Tanaka, M. Murata, M. Kaise, N. Shinya, Tetsu-To-Hagane 74 (1988) 2009 in Japanese. [6] M. Murata, H. Tanaka, E. Abe, H. Irie, Proceedings of the Eighth. International Conference on Creep and Fracture of Engineering Materials and Structures, Key Eng. Mater. 171 –174 (1999) 513. [7] National Research Institute for Metals Creep Data Sheet, Metallographic Atlas of Long-Term Crept Materials, National Research Institute for Metals, Japan, No. M-1 (1999). [8] National Research Institute for Metals Creep Data Sheet, National Research Institute for Metals, Japan, No. 4B (1986).