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A new approach to improve creep resistance of high Cr martensitic steel
Journal of Nuclear Materials 417 (2011) 29–32
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
A new approach to improve creep resistance of high Cr martensitic steel Manabu Tamura a,⇑, Takuya Kumagai b, Kazuhisa Sakai b, Kei Shinozuka a, Hisao Esaka a a b
Department of Materials Science and Engineering, National Defense Academy, 1-10-20 Hashirimizu Yokosuka 239-8686, Japan National Defense Academy, 1-10-20 Hashirimizu Yokosuka 239-8686, Japan
a r t i c l e
i n f o
Article history: Available online 24 December 2010
a b s t r a c t A modified 9Cr–1Mo steel was cooled to 200 °C from the normalizing temperature and then directly heated to the tempering temperature. It was found that the time to rupture at 650–700 °C for the steel heat-treated at 200 °C increased three times over than that of the modified 9Cr–1Mo steel conventionally normalized and tempered. The microstructure of the improved steel was tempered martensite and the size of martensite blocks was larger than for the conventional treatment. The hardness of the improved steel was adequately recovered after tempering. Aging tests showed that the particle sizes of Cr23C6 and VN type carbonitride in the improved steel were finer in the conventional steel. The above-mentioned heat treatment was applied to the reduced activation martensitic steel F-82H and the improvement was confirmed. Ó 2010 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental
High chromium ferritic/martensitic heat resistant steels have been developed for designing fusion reactors, because modified ferritic/martensitic steels could reduce the induced activation as compared to austenitic steel [1]. One of the critical issues for the components of ferritic/martensitic steels is ensuring sufficient creep strength under the operation conditions. Many ferritic/martensitic heat resistant steels with very high creep strengths have been developed recently [2], however many researchers have only focused on the modification of chemical composition, and not on the heat treatment. Meanwhile, the authors recently proposed a new heat treatment, which improved the creep strength of modified 9Cr steel [3]. Although the conventional modified 9Cr steel has been cooled to room temperature after heating at the normalizing temperature, the newly heat-treated modified 9Cr steel was cooled below the Ms point without cooling to room temperature, and then directly heated to the tempering temperature. The modified 9Cr steel was strengthened by finely dispersed VN particles [4,5]. However, the reduced activation martensitic steel F-82H [6] does not contain N. Therefore, it was not known if the new heat treatment is effective for F-82H. In this study basic properties of modified 9Cr steel processed by the new heat treatment are presented and then creep properties of F-82H processed by the new heat treatment are compared with F-82H treated by the conventional method.
The starting modified 9Cr steel was a 40 mm thick plate (ASME SA387 Grade 91) and reduced activation steel was F-82H in a 25 mm thick plate (pre-IEA heat). Each reference material was renormalized at 1150 °C followed by tempering at 750 °C (designated NT hereafter). The normalizing temperature was raised compared to the usual case in order to minimize un-dissolved precipitates. After heating at the normalizing temperature, the test blocks were moved to a furnace preheated to 200 °C, below the Ms point. The blocks were heated at 200 °C for 1 h and then directly heated to the tempering temperature (this heat treatment is denoted as Z200 hereafter). The size of martensite blocks was measured for both NT and Z200 using an optical microscope. Ten micrographs with magnification 1800 were prepared for each specimen to measure the size. Two lines at right angles to each other were drawn on each photograph. The lines were marked at the cross points of the lines and boundaries where differences in orientation contrast or a line of precipitates were observed. The average block sizes of the steels were then calculated. Vickers hardness was also measured under loading of 10 kg. Creep tests for each steel were carried out at 650 and 700 °C. In order to estimate the stability of the microstructure after a long time creep exposure, test blocks of Grade 91 were tempered or aged at 750 °C for 1000 h maximum. Electrolytic dissolution of the samples was carried out and the residue was extracted using a polycarbonate filter having pores of 0.1 lm diameter. The extracted residue was analyzed by X-ray diffraction (XRD). The peak profiles of (511) M23C6, (220) NbX, and (220) VX were measured, where M denotes metallic elements and X denotes C and/or N. The observed integral breadths were
⇑ Corresponding author. Address: 6-45-2 Hinominami, konann-ku, Yokohama 2340055, Japan. Tel./fax: +81 45 891 1039. E-mail address: [email protected] (M. Tamura). 0022-3115/$ - see front matter Ó 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jnucmat.2010.12.043
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M. Tamura et al. / Journal of Nuclear Materials 417 (2011) 29–32
103
500 o
650 C, 150.1 MPa : NT : Z200
101
NT Z200
Particle size / nm
Creep rate / %/h
102
100 10-1 10-2 10-2
100 50
Grade 91 10 -1 10 100
Grade 91
10-3 10-3
M23C6 NbX VX
10-1
100 101 Time / h
102
103
Fig. 1. Creep rate vs. time for conventionally normalized and tempered Grade 91 (NT) and Grade 91 heat-treated by the new process (Z200).
Table 1 Hardness, amount of extracted residue and block size of Grade 91 with variation of heat treatment. N denotes conventional normalizing. Steel
Hardness (HV)
Amount of extracted residue (%)
Block size before aging (lm)
N NT Z200
399 234 255
0.26 1.95 1.95
– 4.0 4.9
converted to particle size of the precipitates using the Scherrer’s relation [7]. 3. Results and discussion 3.1. Grade 91 Fig. 1 shows the relation of creep rate to creep time of Grade 91/ NT and Grade 91/Z200 tested at 650 °C/150.1 MPa. This figure shows that time to rupture increases about three times as much by the new heat treatment. There was a transition creep region terminating in a minimum creep rate and converting to accelerating creep. From the start of the test to about 10 h, the creep rate of Grade 91/Z200 was slightly lower than that of Grade 91/NT. However, in Grade 91/Z200 transition creep continued for a longer time and the minimum creep rate was lower. As a result, the time to rupture of Grade 91/Z200 was longer than that of Grade 91/NT. Both rupture elongation and reduction of area of Grade 91/Z200 were in no way inferior to that of Grade 91/NT. The microstructure of as normalized Grade 91 was a mosaic pattern of martensite blocks and showed scarce precipitation. On the other hand, microstructures of Grade 91/NT and Grade 91/ Z200 were tempered martensite and many precipitates were seen on block boundaries and/or within the blocks. However, no significant difference was seen between Grade 91/NT and Grade 91/ Z200. Table 1 shows Vickers hardness and amount of extracted residue for the normalized state, NT, and Z200 of Grade 91. Though Grade 91/Z200 was slightly harder than Grade 91/NT, both were soft compared with the normalized state. The amounts of extracted residue for NT and Z200 of Grade 91 were the same. The amount of extracted residue for the normalized Grade 91 was small and corresponded to mainly un-dissolved Nb carbonitride. Table 1 shows also the average block size for NT and Z200 of Grade 91. The block size for Grade 91/Z200 was larger than for Grade 91/NT. The block boundaries defined in this study include lines of precipitates with-
101
102
103
Tempering time at 750oC / h Fig. 2. Variation in particle sizes of M23C6, NbX, and VX as a function of tempering time at 750 °C. Conventionally normalized and tempered Grade 91 is represented by open symbols (NT) and solid symbols (Z200) denote data for Grade 91 heattreated by the new process.
in blocks without large angle boundaries, i.e. prior austenite, packets, and block boundaries, and therefore, the block boundaries measured in this study include some lath boundaries. The measured block size listed in Table 1 indicates that the new heat treatment changed the morphology of the martensite blocks as compared to the usual normalized-and-tempered microstructure. Typical precipitates in the as tempered Grade 91 were M23C6, NbX, and VX [8]. The size of these particles was measured using XRD and the results are shown in Fig. 2 as a function of tempering time at 750 °C. The ranges of particle sizes measured for M23C6 and VX shown in Fig. 2 are 50–70 and 20–40 nm, respectively. These values were comparable to the reported values for modified 9Cr steel [9]. We have also confirmed that the sizes of precipitates measured by XRD agreed with the sizes in high Cr steels measured by TEM [9,10]. The size of NbX was rather large, ranging from 45 to 70 nm. This was caused by the coexistence of un-dissolved NbX. The sizes of M23C6 and VX for Grade 91/Z200 were finer than in Grade 91/NT. This situation continued after tempering at 750 °C for 24 h. Though the particle size after tempering for 1000 h was measured only for Grade 91/Z200, the growth rates of M23C6 and VX particles for both NT and Z200 of Grade 91, i.e. the slope of each line shown in Fig. 2, were very low compared to that for the Ostwald ripening. That is, as shown in Fig. 2, the sizes of the precipitates in Grade 91/Z200 were finer than those of NT after a long-term high temperature exposure at 750 °C. Therefore, it is expected that the finer size of precipitates in Grade 91/Z200 would be retained during a long-term exposure at creep temperatures. 3.2. Strengthening mechanism in Grade 91/Z200 As shown in Fig. 1, it was confirmed that the new heat treatment, Z200, clearly prolonged the time to rupture of Grade 91. Fig. 2 shows the prospect of long-term stability of microstructure of Grade 91/Z200 at creep temperatures, which is an important factor in developing heat resistant materials. Creep tests on weld joints, toughness measurements, etc. are necessary before applying the new heat treatment to practical use; however the creep mechanism is the first thing to be considered. First, some amount of austenite is considered to be left at 200 °C without forming martensite, based on the dilatation curve for Grade 91 [11]. If the retained austenite is assumed to still be stable after heating at 750 °C, it should be retained at room temperature or be transformed into martensite during cooling to room temperature after tempering. However, retained austenite was not detected by XRD measurement on a bulk specimen of Grade 91/ Z200. Moreover, the initial hardness and transition creep rate for Grade 91/NT and Grade 91/Z200 were similar. Therefore, the
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M. Tamura et al. / Journal of Nuclear Materials 417 (2011) 29–32
At 200oC
At 500oC-750oC
Retained γ M23C6
C,N
VX
VX Martensite Diffusion of C, N to γ
M23C6 Precipitation of M23C6 and VX
M23C6
1) Transformation of γ to α maintaining the same orientation with adjacent martensite (α) and larger blocks are formed. 2) M23C6 precipitates with higher density on the boundaries.
Fig. 3. Schematic illustration of the strengthening mechanism by the new heat treatment.
and adjacent tempered martensite, which originally transformed to martensite at 200 °C and is tempered at 750 °C. This mechanism allows the block size of Grade 91/Z200 to be larger than that of Grade 91/NT. Fig. 3 schematically shows the strengthening mechanism by the new heat treatment. The above discussion mainly focused on the stability of block boundaries. On the other hand, Fig. 2 clearly showed that the precipitates in Grade 91/Z200, especially the smallest compound, VX, were finer than those in Grade 91/NT. Finer precipitates of VX in Grade 91/Z200 are dispersed uniformly within martensite blocks and ferrite. This can also explain the increase in time to rupture for the new heat treatment. Further studies are necessary to explaining why the new heat treatment results in finer precipitates. 3.3. F-82H The strengthening produced by the new heat treatment for Grade 91 could be caused by the stabilization of block boundaries due to increase in the block size and the refining of precipitates, especially VX. Since F-82H does not contain N, the strengthening effect by VX cannot be expected. However, it is worthwhile to apply the new heat treatment to F-82H. The time to rupture as a function of stress is presented in Fig. 4. This shows that the time to rupture of F-82H/Z200 is about twice that of F-82H/NT. The experimental results of rupture elongation of F-82H/Z200 are also indicated near the data points and are comparable to those of F-82H/NT.
200 NT Z200
Stress / MPa
retained austenite, which was untransformed at 200 °C in Grade 91, may have decomposed into ferrite and carbonitride during heating or holding at the tempering temperature. That is, so-called fresh martensite, which might exist in Grade 91/Z200, is not responsible for the strengthening by the new heat treatment. As shown in Table 1, the block size of Grade 91/Z200 was larger than that of Grade 91/NT, but the amount of extracted residue was the same. Most of the precipitates extracted from Grade 91/NT and Grade 91/Z200 were M23C6 and they precipitate preferentially on prior austenitic, packet, block, and lath boundaries. All these boundaries were counted as block boundaries in this study. These facts show that the precipitates occupied the block boundaries of Grade 91/Z200 with higher density than in Grade 91/NT, assuming similar particle size for both. Fig. 2 showed that the particle size for Grade 91/Z200 was finer than that of Grade 91/NT, so the occupation density on the block boundaries by precipitates could be higher than the above estimation. The occupation by precipitates with higher density stabilizes the block boundaries of Grade 91/Z200 and restrains the growth of the width of block and/or of lath boundaries. This results in the delay of the onset of accelerating creep. Abe [12] studied the relation between time to rupture of a high Cr ferritic steel and percentage of occupation near prior austenitic grain boundary by precipitates and proposed a similar explanation for creep strengthening of high Cr ferritic steel. Retained austenite was usually observed as thin films along block boundaries or lath boundaries [13]. Since only interstitial elements can move at 200 °C, excess C and N in martensite diffuse into the adjacent retained austenite, resulting in the formation of austenite with higher contents of C and/or N. This could cause the high density of precipitation of carbonitride above 500 °C in the retained austenite. Although M23C6 could precipitate in the matrix of austenite with high coherency, M23C6 preferentially precipitates on the boundaries or austenite/ferrite interfaces. The precipitation of M23C6 decreases Cr content in austenite, which causes the retained austenite to be unstable. Okamoto and Oka [14] showed that the bainitic decomposition of austenite at martensite/austenite interfaces took place during isothermal aging even about 200 °C in hypereutectoid steels. Although the bainitic reaction has not been reported to occur in high Cr steels, it is not denied that bainitic-like decomposition of austenite in Grade 91/Z200 could be possible during heating at 200 °C or tempering. Fultz et al. [15] observed in 9Ni steel that the austenite/martensite interface was originally quite coherent over 10 nm. Therefore, it is assumed that a certain orientation relation exists or the same orientation is maintain between the newly decomposed ferrite
150 650oC 33
o
700 C
100
25
50
37 30
24 24
F-82H
100
101 102 103 Time to rupture / h
104
Fig. 4. Time to rupture vs. stress for conventionally normalized and tempered F82H (NT) and F-82H processed by the new method (Z200). Numbers near data symbol indicates rupture elongation in per cent.
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M. Tamura et al. / Journal of Nuclear Materials 417 (2011) 29–32
The Vickers hardness of F-82H/Z200 is 231, little harder than the usual normalized and tempered state of 220, but F-82H/Z200 is judged to be adequately tempered. The microstructure of F-82H/ Z200 was tempered martensite. The block size measurements for F-82H showed the frequency of block sizes of F-82H/Z200 in the range of 1–2 lm was low compared to that of F-82H/NT. In contrast, the frequency of block sizes in the range of 5–10 lm was high for F-82H/Z200. The average size (2.4 lm) for F-82H/Z200 was larger than for NT (2.2 lm). These features were similar to the case of Grade 91 (Table 1). Precipitates in both F-82H/Z200 and F-82H/NT were all M23C6. The amount of extracted residue for both heat treatments was equal at 1.7%. The particle sizes measured by XRD were 35 nm for F-82H/Z200 and 42 nm for F-82H/NT. As a result, the block boundaries of F-82H/Z200 are more strongly stabilized than those of F-82H/NT. Therefore, the similar strengthening mechanism as was proposed to Grade 91 may operate in F-82H/Z200, even though the strengthening by VX was absent in F-82H. F-82H/Z200 was adequately softened by tempering and the slope of the time to rupture vs. stress relation is the same as that of F-82H/NT, shown in Fig. 4. If the creep of F-82H/Z200 is improved by a fresh martensite or a high dislocation density, the slope would be steeper than that of F-82H/NT [16]. Moreover, the creep rate vs. time relation for F-82H was similar to that of Grade 91, shown in Fig. 1. The creep rate in the transition creep region for F-82H/Z200 was very close to that of F-82H/NT. Therefore, a fresh martensite or a higher dislocation density is ruled out as the major reason for the strengthening shown in Fig. 4. The block size of F-82H was smaller than that of Grade 91. However, larger block size is preferable for martensitic steels to prevent recovery near the boundaries. Therefore, further study is necessary to improve the creep resistance of F-82H. We have confirmed that the new heat treatment effectively improved the creep resistance of F-82H without reducing creep ductility. Components of fusion reactors will be large, complex structure so it will take a long time to cool to room temperature from the normalizing temperature. If the direct tempering after cooling only to a temperature below the Ms point is acceptable, the fabrication time could be shortened without degrading the creep properties. Further studies aiming at a practical application of the new heat treatment are desirable.
4. Conclusions The creep strength of a modified 9Cr steel (ASME Grade 91) was improved by applying a new heat treatment: after cooling to 200 °C, below the Ms point, in the cooling process of normalizing, the steel is directly reheated to the tempering temperature without cooling to room temperature. A strengthening mechanism for this new heat treatment was proposed. The new process increased the size of martensite blocks, resulting in a higher density of boundary precipitates. These microstructural features could stabilize the block boundaries, which increases the creep resistance. The size of precipitates, decreased by the new process, could also contribute to the creep improvement. This process was applied to the reduced activation steel F-82H and the effectiveness of the new process was confirmed. References [1] R.L. Klueh, D.R. Harries, Development of high (7–12%) chromium martensitic steels, in: R.L. Klueh, D.R. Harries (Eds.), High-Chromium Ferritic and Martensitic Steels for Nuclear Applications, ASTM, West Conshohocken, PA, 2001, pp. 5–27. [2] F. Masuyama, ISIJ Int. 41 (2001) 612–625. [3] M. Tamura, T. Kumagai, K. Shinozuka, H. Esaka, CAMP ISIJ 21 (2008) 1418. [4] V.K. Sikka, C.T. Ward, K.C. Thomas, Modified 9 Cr–1 Mo Steel – An Improved Alloy for Steam Generator Application, in: Ferritic Steel for High Temperature Applications, AMS, Metals Park, OH, 1983, pp. 65–84. [5] L. Lundin, S. Fallman, H.-O. Andren, Mater. Sci. Technol. 13 (1997) 233–242. [6] M. Tamura, H. Hayakawa, M. Tanimura, A. Hishinuma, T. Kondo, J. Nucl. Mater. 141–143 (1986) 1067–1073. [7] F.W. Jones, Proc. Roy. Soc. 116A (1938) 16–43. [8] W.B. Jones, C.R. Hills, D.H. Polonis, Metall. Mater. Trans. A 22 (1991) 1049–1058. [9] M. Tamura, H. Kusuyama, K. Shinozuka, H. Esaka, ISIJ Int. 47 (2007) 317–326. [10] M. Tamura, M. Nakamura, K. Shinozuka, H. Esaka, Metall. Mater. Trans. A 39 (2008) 1060–1076. [11] F. Masuyama, N. Nishimura, Phase transformation and properties of Gr. 91 at around critical temperature, in: Y.-Y. Wang (Ed.), Experience with CreepStrength Enhanced Ferritic Steels and New Emerging Computational Methods, ASME, New York, NY, 2004, pp. 85–91. [12] F. Abe, Sci. Technol. Adv. Mater. 9 (2008) 013002-1–013002-15. [13] H.K.D.H. Bhadeshia, Tempered Martensite, pp. 1–16.
Contents lists available at ScienceDirect
Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
A new approach to improve creep resistance of high Cr martensitic steel Manabu Tamura a,⇑, Takuya Kumagai b, Kazuhisa Sakai b, Kei Shinozuka a, Hisao Esaka a a b
Department of Materials Science and Engineering, National Defense Academy, 1-10-20 Hashirimizu Yokosuka 239-8686, Japan National Defense Academy, 1-10-20 Hashirimizu Yokosuka 239-8686, Japan
a r t i c l e
i n f o
Article history: Available online 24 December 2010
a b s t r a c t A modified 9Cr–1Mo steel was cooled to 200 °C from the normalizing temperature and then directly heated to the tempering temperature. It was found that the time to rupture at 650–700 °C for the steel heat-treated at 200 °C increased three times over than that of the modified 9Cr–1Mo steel conventionally normalized and tempered. The microstructure of the improved steel was tempered martensite and the size of martensite blocks was larger than for the conventional treatment. The hardness of the improved steel was adequately recovered after tempering. Aging tests showed that the particle sizes of Cr23C6 and VN type carbonitride in the improved steel were finer in the conventional steel. The above-mentioned heat treatment was applied to the reduced activation martensitic steel F-82H and the improvement was confirmed. Ó 2010 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental
High chromium ferritic/martensitic heat resistant steels have been developed for designing fusion reactors, because modified ferritic/martensitic steels could reduce the induced activation as compared to austenitic steel [1]. One of the critical issues for the components of ferritic/martensitic steels is ensuring sufficient creep strength under the operation conditions. Many ferritic/martensitic heat resistant steels with very high creep strengths have been developed recently [2], however many researchers have only focused on the modification of chemical composition, and not on the heat treatment. Meanwhile, the authors recently proposed a new heat treatment, which improved the creep strength of modified 9Cr steel [3]. Although the conventional modified 9Cr steel has been cooled to room temperature after heating at the normalizing temperature, the newly heat-treated modified 9Cr steel was cooled below the Ms point without cooling to room temperature, and then directly heated to the tempering temperature. The modified 9Cr steel was strengthened by finely dispersed VN particles [4,5]. However, the reduced activation martensitic steel F-82H [6] does not contain N. Therefore, it was not known if the new heat treatment is effective for F-82H. In this study basic properties of modified 9Cr steel processed by the new heat treatment are presented and then creep properties of F-82H processed by the new heat treatment are compared with F-82H treated by the conventional method.
The starting modified 9Cr steel was a 40 mm thick plate (ASME SA387 Grade 91) and reduced activation steel was F-82H in a 25 mm thick plate (pre-IEA heat). Each reference material was renormalized at 1150 °C followed by tempering at 750 °C (designated NT hereafter). The normalizing temperature was raised compared to the usual case in order to minimize un-dissolved precipitates. After heating at the normalizing temperature, the test blocks were moved to a furnace preheated to 200 °C, below the Ms point. The blocks were heated at 200 °C for 1 h and then directly heated to the tempering temperature (this heat treatment is denoted as Z200 hereafter). The size of martensite blocks was measured for both NT and Z200 using an optical microscope. Ten micrographs with magnification 1800 were prepared for each specimen to measure the size. Two lines at right angles to each other were drawn on each photograph. The lines were marked at the cross points of the lines and boundaries where differences in orientation contrast or a line of precipitates were observed. The average block sizes of the steels were then calculated. Vickers hardness was also measured under loading of 10 kg. Creep tests for each steel were carried out at 650 and 700 °C. In order to estimate the stability of the microstructure after a long time creep exposure, test blocks of Grade 91 were tempered or aged at 750 °C for 1000 h maximum. Electrolytic dissolution of the samples was carried out and the residue was extracted using a polycarbonate filter having pores of 0.1 lm diameter. The extracted residue was analyzed by X-ray diffraction (XRD). The peak profiles of (511) M23C6, (220) NbX, and (220) VX were measured, where M denotes metallic elements and X denotes C and/or N. The observed integral breadths were
⇑ Corresponding author. Address: 6-45-2 Hinominami, konann-ku, Yokohama 2340055, Japan. Tel./fax: +81 45 891 1039. E-mail address: [email protected] (M. Tamura). 0022-3115/$ - see front matter Ó 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jnucmat.2010.12.043
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M. Tamura et al. / Journal of Nuclear Materials 417 (2011) 29–32
103
500 o
650 C, 150.1 MPa : NT : Z200
101
NT Z200
Particle size / nm
Creep rate / %/h
102
100 10-1 10-2 10-2
100 50
Grade 91 10 -1 10 100
Grade 91
10-3 10-3
M23C6 NbX VX
10-1
100 101 Time / h
102
103
Fig. 1. Creep rate vs. time for conventionally normalized and tempered Grade 91 (NT) and Grade 91 heat-treated by the new process (Z200).
Table 1 Hardness, amount of extracted residue and block size of Grade 91 with variation of heat treatment. N denotes conventional normalizing. Steel
Hardness (HV)
Amount of extracted residue (%)
Block size before aging (lm)
N NT Z200
399 234 255
0.26 1.95 1.95
– 4.0 4.9
converted to particle size of the precipitates using the Scherrer’s relation [7]. 3. Results and discussion 3.1. Grade 91 Fig. 1 shows the relation of creep rate to creep time of Grade 91/ NT and Grade 91/Z200 tested at 650 °C/150.1 MPa. This figure shows that time to rupture increases about three times as much by the new heat treatment. There was a transition creep region terminating in a minimum creep rate and converting to accelerating creep. From the start of the test to about 10 h, the creep rate of Grade 91/Z200 was slightly lower than that of Grade 91/NT. However, in Grade 91/Z200 transition creep continued for a longer time and the minimum creep rate was lower. As a result, the time to rupture of Grade 91/Z200 was longer than that of Grade 91/NT. Both rupture elongation and reduction of area of Grade 91/Z200 were in no way inferior to that of Grade 91/NT. The microstructure of as normalized Grade 91 was a mosaic pattern of martensite blocks and showed scarce precipitation. On the other hand, microstructures of Grade 91/NT and Grade 91/ Z200 were tempered martensite and many precipitates were seen on block boundaries and/or within the blocks. However, no significant difference was seen between Grade 91/NT and Grade 91/ Z200. Table 1 shows Vickers hardness and amount of extracted residue for the normalized state, NT, and Z200 of Grade 91. Though Grade 91/Z200 was slightly harder than Grade 91/NT, both were soft compared with the normalized state. The amounts of extracted residue for NT and Z200 of Grade 91 were the same. The amount of extracted residue for the normalized Grade 91 was small and corresponded to mainly un-dissolved Nb carbonitride. Table 1 shows also the average block size for NT and Z200 of Grade 91. The block size for Grade 91/Z200 was larger than for Grade 91/NT. The block boundaries defined in this study include lines of precipitates with-
101
102
103
Tempering time at 750oC / h Fig. 2. Variation in particle sizes of M23C6, NbX, and VX as a function of tempering time at 750 °C. Conventionally normalized and tempered Grade 91 is represented by open symbols (NT) and solid symbols (Z200) denote data for Grade 91 heattreated by the new process.
in blocks without large angle boundaries, i.e. prior austenite, packets, and block boundaries, and therefore, the block boundaries measured in this study include some lath boundaries. The measured block size listed in Table 1 indicates that the new heat treatment changed the morphology of the martensite blocks as compared to the usual normalized-and-tempered microstructure. Typical precipitates in the as tempered Grade 91 were M23C6, NbX, and VX [8]. The size of these particles was measured using XRD and the results are shown in Fig. 2 as a function of tempering time at 750 °C. The ranges of particle sizes measured for M23C6 and VX shown in Fig. 2 are 50–70 and 20–40 nm, respectively. These values were comparable to the reported values for modified 9Cr steel [9]. We have also confirmed that the sizes of precipitates measured by XRD agreed with the sizes in high Cr steels measured by TEM [9,10]. The size of NbX was rather large, ranging from 45 to 70 nm. This was caused by the coexistence of un-dissolved NbX. The sizes of M23C6 and VX for Grade 91/Z200 were finer than in Grade 91/NT. This situation continued after tempering at 750 °C for 24 h. Though the particle size after tempering for 1000 h was measured only for Grade 91/Z200, the growth rates of M23C6 and VX particles for both NT and Z200 of Grade 91, i.e. the slope of each line shown in Fig. 2, were very low compared to that for the Ostwald ripening. That is, as shown in Fig. 2, the sizes of the precipitates in Grade 91/Z200 were finer than those of NT after a long-term high temperature exposure at 750 °C. Therefore, it is expected that the finer size of precipitates in Grade 91/Z200 would be retained during a long-term exposure at creep temperatures. 3.2. Strengthening mechanism in Grade 91/Z200 As shown in Fig. 1, it was confirmed that the new heat treatment, Z200, clearly prolonged the time to rupture of Grade 91. Fig. 2 shows the prospect of long-term stability of microstructure of Grade 91/Z200 at creep temperatures, which is an important factor in developing heat resistant materials. Creep tests on weld joints, toughness measurements, etc. are necessary before applying the new heat treatment to practical use; however the creep mechanism is the first thing to be considered. First, some amount of austenite is considered to be left at 200 °C without forming martensite, based on the dilatation curve for Grade 91 [11]. If the retained austenite is assumed to still be stable after heating at 750 °C, it should be retained at room temperature or be transformed into martensite during cooling to room temperature after tempering. However, retained austenite was not detected by XRD measurement on a bulk specimen of Grade 91/ Z200. Moreover, the initial hardness and transition creep rate for Grade 91/NT and Grade 91/Z200 were similar. Therefore, the
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M. Tamura et al. / Journal of Nuclear Materials 417 (2011) 29–32
At 200oC
At 500oC-750oC
Retained γ M23C6
C,N
VX
VX Martensite Diffusion of C, N to γ
M23C6 Precipitation of M23C6 and VX
M23C6
1) Transformation of γ to α maintaining the same orientation with adjacent martensite (α) and larger blocks are formed. 2) M23C6 precipitates with higher density on the boundaries.
Fig. 3. Schematic illustration of the strengthening mechanism by the new heat treatment.
and adjacent tempered martensite, which originally transformed to martensite at 200 °C and is tempered at 750 °C. This mechanism allows the block size of Grade 91/Z200 to be larger than that of Grade 91/NT. Fig. 3 schematically shows the strengthening mechanism by the new heat treatment. The above discussion mainly focused on the stability of block boundaries. On the other hand, Fig. 2 clearly showed that the precipitates in Grade 91/Z200, especially the smallest compound, VX, were finer than those in Grade 91/NT. Finer precipitates of VX in Grade 91/Z200 are dispersed uniformly within martensite blocks and ferrite. This can also explain the increase in time to rupture for the new heat treatment. Further studies are necessary to explaining why the new heat treatment results in finer precipitates. 3.3. F-82H The strengthening produced by the new heat treatment for Grade 91 could be caused by the stabilization of block boundaries due to increase in the block size and the refining of precipitates, especially VX. Since F-82H does not contain N, the strengthening effect by VX cannot be expected. However, it is worthwhile to apply the new heat treatment to F-82H. The time to rupture as a function of stress is presented in Fig. 4. This shows that the time to rupture of F-82H/Z200 is about twice that of F-82H/NT. The experimental results of rupture elongation of F-82H/Z200 are also indicated near the data points and are comparable to those of F-82H/NT.
200 NT Z200
Stress / MPa
retained austenite, which was untransformed at 200 °C in Grade 91, may have decomposed into ferrite and carbonitride during heating or holding at the tempering temperature. That is, so-called fresh martensite, which might exist in Grade 91/Z200, is not responsible for the strengthening by the new heat treatment. As shown in Table 1, the block size of Grade 91/Z200 was larger than that of Grade 91/NT, but the amount of extracted residue was the same. Most of the precipitates extracted from Grade 91/NT and Grade 91/Z200 were M23C6 and they precipitate preferentially on prior austenitic, packet, block, and lath boundaries. All these boundaries were counted as block boundaries in this study. These facts show that the precipitates occupied the block boundaries of Grade 91/Z200 with higher density than in Grade 91/NT, assuming similar particle size for both. Fig. 2 showed that the particle size for Grade 91/Z200 was finer than that of Grade 91/NT, so the occupation density on the block boundaries by precipitates could be higher than the above estimation. The occupation by precipitates with higher density stabilizes the block boundaries of Grade 91/Z200 and restrains the growth of the width of block and/or of lath boundaries. This results in the delay of the onset of accelerating creep. Abe [12] studied the relation between time to rupture of a high Cr ferritic steel and percentage of occupation near prior austenitic grain boundary by precipitates and proposed a similar explanation for creep strengthening of high Cr ferritic steel. Retained austenite was usually observed as thin films along block boundaries or lath boundaries [13]. Since only interstitial elements can move at 200 °C, excess C and N in martensite diffuse into the adjacent retained austenite, resulting in the formation of austenite with higher contents of C and/or N. This could cause the high density of precipitation of carbonitride above 500 °C in the retained austenite. Although M23C6 could precipitate in the matrix of austenite with high coherency, M23C6 preferentially precipitates on the boundaries or austenite/ferrite interfaces. The precipitation of M23C6 decreases Cr content in austenite, which causes the retained austenite to be unstable. Okamoto and Oka [14] showed that the bainitic decomposition of austenite at martensite/austenite interfaces took place during isothermal aging even about 200 °C in hypereutectoid steels. Although the bainitic reaction has not been reported to occur in high Cr steels, it is not denied that bainitic-like decomposition of austenite in Grade 91/Z200 could be possible during heating at 200 °C or tempering. Fultz et al. [15] observed in 9Ni steel that the austenite/martensite interface was originally quite coherent over 10 nm. Therefore, it is assumed that a certain orientation relation exists or the same orientation is maintain between the newly decomposed ferrite
150 650oC 33
o
700 C
100
25
50
37 30
24 24
F-82H
100
101 102 103 Time to rupture / h
104
Fig. 4. Time to rupture vs. stress for conventionally normalized and tempered F82H (NT) and F-82H processed by the new method (Z200). Numbers near data symbol indicates rupture elongation in per cent.
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The Vickers hardness of F-82H/Z200 is 231, little harder than the usual normalized and tempered state of 220, but F-82H/Z200 is judged to be adequately tempered. The microstructure of F-82H/ Z200 was tempered martensite. The block size measurements for F-82H showed the frequency of block sizes of F-82H/Z200 in the range of 1–2 lm was low compared to that of F-82H/NT. In contrast, the frequency of block sizes in the range of 5–10 lm was high for F-82H/Z200. The average size (2.4 lm) for F-82H/Z200 was larger than for NT (2.2 lm). These features were similar to the case of Grade 91 (Table 1). Precipitates in both F-82H/Z200 and F-82H/NT were all M23C6. The amount of extracted residue for both heat treatments was equal at 1.7%. The particle sizes measured by XRD were 35 nm for F-82H/Z200 and 42 nm for F-82H/NT. As a result, the block boundaries of F-82H/Z200 are more strongly stabilized than those of F-82H/NT. Therefore, the similar strengthening mechanism as was proposed to Grade 91 may operate in F-82H/Z200, even though the strengthening by VX was absent in F-82H. F-82H/Z200 was adequately softened by tempering and the slope of the time to rupture vs. stress relation is the same as that of F-82H/NT, shown in Fig. 4. If the creep of F-82H/Z200 is improved by a fresh martensite or a high dislocation density, the slope would be steeper than that of F-82H/NT [16]. Moreover, the creep rate vs. time relation for F-82H was similar to that of Grade 91, shown in Fig. 1. The creep rate in the transition creep region for F-82H/Z200 was very close to that of F-82H/NT. Therefore, a fresh martensite or a higher dislocation density is ruled out as the major reason for the strengthening shown in Fig. 4. The block size of F-82H was smaller than that of Grade 91. However, larger block size is preferable for martensitic steels to prevent recovery near the boundaries. Therefore, further study is necessary to improve the creep resistance of F-82H. We have confirmed that the new heat treatment effectively improved the creep resistance of F-82H without reducing creep ductility. Components of fusion reactors will be large, complex structure so it will take a long time to cool to room temperature from the normalizing temperature. If the direct tempering after cooling only to a temperature below the Ms point is acceptable, the fabrication time could be shortened without degrading the creep properties. Further studies aiming at a practical application of the new heat treatment are desirable.
4. Conclusions The creep strength of a modified 9Cr steel (ASME Grade 91) was improved by applying a new heat treatment: after cooling to 200 °C, below the Ms point, in the cooling process of normalizing, the steel is directly reheated to the tempering temperature without cooling to room temperature. A strengthening mechanism for this new heat treatment was proposed. The new process increased the size of martensite blocks, resulting in a higher density of boundary precipitates. These microstructural features could stabilize the block boundaries, which increases the creep resistance. The size of precipitates, decreased by the new process, could also contribute to the creep improvement. This process was applied to the reduced activation steel F-82H and the effectiveness of the new process was confirmed. References [1] R.L. Klueh, D.R. Harries, Development of high (7–12%) chromium martensitic steels, in: R.L. Klueh, D.R. Harries (Eds.), High-Chromium Ferritic and Martensitic Steels for Nuclear Applications, ASTM, West Conshohocken, PA, 2001, pp. 5–27. [2] F. Masuyama, ISIJ Int. 41 (2001) 612–625. [3] M. Tamura, T. Kumagai, K. Shinozuka, H. Esaka, CAMP ISIJ 21 (2008) 1418. [4] V.K. Sikka, C.T. Ward, K.C. Thomas, Modified 9 Cr–1 Mo Steel – An Improved Alloy for Steam Generator Application, in: Ferritic Steel for High Temperature Applications, AMS, Metals Park, OH, 1983, pp. 65–84. [5] L. Lundin, S. Fallman, H.-O. Andren, Mater. Sci. Technol. 13 (1997) 233–242. [6] M. Tamura, H. Hayakawa, M. Tanimura, A. Hishinuma, T. Kondo, J. Nucl. Mater. 141–143 (1986) 1067–1073. [7] F.W. Jones, Proc. Roy. Soc. 116A (1938) 16–43. [8] W.B. Jones, C.R. Hills, D.H. Polonis, Metall. Mater. Trans. A 22 (1991) 1049–1058. [9] M. Tamura, H. Kusuyama, K. Shinozuka, H. Esaka, ISIJ Int. 47 (2007) 317–326. [10] M. Tamura, M. Nakamura, K. Shinozuka, H. Esaka, Metall. Mater. Trans. A 39 (2008) 1060–1076. [11] F. Masuyama, N. Nishimura, Phase transformation and properties of Gr. 91 at around critical temperature, in: Y.-Y. Wang (Ed.), Experience with CreepStrength Enhanced Ferritic Steels and New Emerging Computational Methods, ASME, New York, NY, 2004, pp. 85–91. [12] F. Abe, Sci. Technol. Adv. Mater. 9 (2008) 013002-1–013002-15. [13] H.K.D.H. Bhadeshia, Tempered Martensite, pp. 1–16.