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Microstructural stability of 11Cr ODS steel
Journal of Nuclear Materials 472 (2016) 247e251
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Microstructural stability of 11Cr ODS steel* Tetsuya Yamashiro a, *, Shigeharu Ukai b, Naoko Oono b, Satoshi Ohtsuka c, Takeji Kaito c a
Materials Science and Engineering, Graduate School of Engineering, Hokkaido University, N13, W-8, Kita-ku, Sapporo 060-8628, Japan Materials Science and Engineering, Faculty of Engineering, Hokkaido University, N13, W-8, Kita-ku, Sapporo 060-8628, Japan c Advanced Nuclear System R&D Directorate, Japan Atomic Energy Agency (JAEA), 4002, Narita, Oarai, Ibaraki-pref. 311-1393, Japan b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 2 May 2015 Received in revised form 30 December 2015 Accepted 4 January 2016 Available online 6 January 2016
Aiming at further improvement of high-temperature oxidation and corrosion resistance, 11CrODS steel with martensitic base structure has been previously developed, as a candidate fuel cladding material for 4th generation advanced nuclear reactors. In this study, the microstructure of 11CrODS steel was characterized by means of EBSD and nanoindentation hardness measurement. The continuous cooling transformation (CCT) diagram was constructed. Upper critical cooling rate, which is minimum cooling rate necessary to form martensitic structure, was derived to be 60 C/min (3600 C/h). In contrast, lower critical cooling rate preventing from martensite formation, was derived to be 10 C/min (600 C/h). An area fraction of so called residual ferrite was estimated by image processing of EBSD-IQ map to be 21% of the total area. This fraction of the residual ferrite in 11CrODS steel was evaluated by considering the driving force for a to g reverse transformation. © 2016 Published by Elsevier B.V.
Keywords: Continuous cooling transformation diagram ODS Residual ferrite a-g phase transformation Pinning force
1. Introduction The oxide dispersion strengthened (ODS) ferritic steels have excellent radiation resistance and high-temperature strength. Therefore, they are promising candidates for cladding materials of Generation IV sodium-cooled fast reactors [1,2] and for blanket structural materials of the advanced fusion reactors [3,4]. The ODS ferritic steels have two types of structures known as alpha ferrite and martensite. We have focused on developing the 9CrODS steel with martensitic structure, since on one hand its structure and processing can be effectively controlled in terms of a-g phase transformation [5e11]. On the other hand, the structure and morphology of the 12CrODS steel as the typical ferritic steel can be only controlled by recrystallization [12]. However in this case there is a lack of repeatability and consequently this kind of a control is unsuitable to mass production [13,14]. In order to further improve a high-temperature oxidation and corrosion resistance in 9CrODS steel, we have developed a 11CrODS steel by increasing the Crcontent from 9 to 11 wt. % [15]. The 11CrODS steel has martensitic structure even at 11 wt. % Cr; thus it has the advantage for controlling the processing-fabrication route by a-g phase transformation. * Presented at the NuMat 2014 Conference, 27e30 October 2014, Clearwater, Florida, USA. * Corresponding author. E-mail address: [email protected] (T. Yamashiro).
http://dx.doi.org/10.1016/j.jnucmat.2016.01.002 0022-3115/© 2016 Published by Elsevier B.V.
A nominal chemical composition of the 11CrODS steel is Fee11Cr-0.13C-1.3W-0.4Ni- 0.3Ti-0.35Y2O3 (wt. %). We constructed phase stability map of the 11CrODS steel in a tungsten-nickel contents diagram, which is compared with those of 9CrODS and 12CrODS steels at 1050 C, keeping 0.3 wt. % Ti, 0.35 wt. % Y2O3 and Fe-balance. This is shown in Fig. 1, where 1050 C corresponds to the normalizing temperature of the ODS steels. These calculations were carried out by means of thermodynamic analysis software Pandat [16,17]. Each line for 9Cr, 11Cr and 12Cr indicates phase boundary between single g and g þ d two phase field. The single gphase is stabilized on the left side from the boundary, while the g þ d field is stable on its right side. In the case of 12CrODS steel, higher Ni-content is required at even lower content of tungsten, to produce the single g-austenite phase at 1050 C. With decreasing Cr-content to 11Cr and 9Cr, the Ni-content required for stabilization of the single g-austenite phase can be decreased. The chemical compositions of 11CrODS steel (1.3W-0.4Ni) and 9CrODS steel (2W without Ni addition) are represented by solid circle in Fig. 1, from which the single g-austenite is stable at 1050 C for 9CrODS steel; thus leading to the martensite formation after cooling. However, the structure of 9CrODS steels is characteristic for a dual phase composed of both martensite and ferrite, and this ferrite was designated as the metastable residual phase [8,18]. It is well known that the residual phase significantly improves the hightemperature strength [19e21]. For this reason, an estimation of
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T. Yamashiro et al. / Journal of Nuclear Materials 472 (2016) 247e251
Fig. 2. Length change of 11CrODS steel determined from dilatometric measurements during cooling with various cooling rate.
Fig. 1. Variation of phase diagram with Ni and W content.
1100a using a Berkovich type of indent, which separately gives a nanohardness of each phase.
the residual ferrite fraction is also important for 11CrODS steel. In this study, 11CrODS steel bars were manufactured, and their structural stability including continuous cooling transform (CCT) diagram were studied. A fraction of the metastable residual ferrite is also assessed from thermodynamic calculations.
3. Result & discussion 3.1. CCT diagram
2. Experimental procedure
The results of dilatometric measurement are shown in Fig. 2. The specimen length changes were recorded during cooling from 1050 C. Nine cooling rates were used, namely: 3 C/min, 6 C/min, 8 C/min, 10 C/min, 12 C/min, 30 C/min, 60 C/min, 300 C/min and 600 C/min. The symbols Ar3 and Ar1 correspond to temperatures of the beginning and of the end of the austenite to ferrite transformation, respectively. Ms and Mf denote temperatures of the beginning and of the end of the martensitic transformation. Temperatures Ar1, Ar3, Ms and Mf measured at each cooling rate are shown in Table 2. The lower critical cooling rate was assessed to be 10 C/min (600 C/h). Below this rate only ferrite transformation takes place. The upper critical cooling rate is found to be around 60 C/min (3600 C/h). For higher rates only martensite transformation occurs. The CCT diagram was constructed from the dilatometric data, and it is shown in Fig. 3. For comparison, diagrams for F82H and 9CrODS steels are also plotted. The F82H is the martensitic type of steel not containing oxide particles, and it is the promising candidate for the advanced fusion blanket structural materials [22]. In the case of F82H, the lower and upper critical cooling rates are shifted to the right side. It means that the martensite transformation easily proceeds even at lower cooling rate, because the ferrite transformation undergoing along prior austenite grain
The 11CrODS steel bar was manufactured by means of mechanical alloying and hot-extrusion. The pre-alloyed powders produced by gas atomizing of the nominal composition Fee11Cr0.13C-1.3W-0.4Ni-0.3Ti (wt. %) were supplied by Sanyo Special Steel Co. Ltd. The size of atomized powders is less than 100 mm, and they were mechanically alloyed together with Y2O3 powder by attrition type ball mill. The following conditions of ball milling were chosen: 220 rpm, 48 h and balls: powder ratio of 15: 1 (mass ratio). The mechanically alloyed powder was consolidated into bar by hotextrusion at 1150 C. The extrusion ratio is 9.2, and outer diameter of the bar is 24 mm. The result of chemical analysis of the extruded bar is shown in Table 1, where an excess oxygen (Ex.O) means the amount of total oxygen minus oxygen bounded with yttrium as Y2O3. The CCT diagram was constructed using a Rigaku model TMA8140C thermo mechanical analyzer. Rectangular shape of specimens in the dimension of 3 mm 3 mm 10 mm were used for this dilatometric studies. The specimen length change was measured along the extrusion direction during cooling from 1050 C with the cooling rate ranging from 3 C/min to 600 C/min. The specimens after dilatometric measurements were polished with colloidal silica. Electron back scattering diffraction (EBSD) measurement was carried out under the condition of accelerating voltage of 20 kV by JEOL JSM-6500F. The image quality (IQ) map was obtained to distinguish martensite from ferrite transformed from austenite, and from residual ferrite. An area of residual ferrite can be dark in IQ map, because the residual ferrite contains a large amount of strain, as compared with the transformed ferrite from austenite. Based on the above information, the average area fraction of the residual ferrite was estimated by means of the image processing of the low magnification EBSD-IQ map. Nanoindentation measurement was also conducted by ELIONIX ENT-
Table 2 Ar3, Ar1, Ms and Mf temperatures determined from dilatometric measurement at various cooling rate. Cooling rate ( C/min)
Ar3 ( C) Ar1 ( C) Ms ( C) Mf ( C)
3
6
8
10
12
30
60
300
600
756 692 e e
756 688 e e
748 672 e e
760 684 e e
728 640 390 275
720 620 405 220
e e 365 220
e e 355 214
e e 360 190
Table 1 Chemical composition of the manufactured 11CrODS steel specimen. Element
C
Si
Mn
P
S
Ni
Cr
W
Ti
Al
Y
O
N
Ar
Y2O3
Ex.O
wt%
0.13
0.04
0.04
0.01
0
0.4
10.9
1.3
0.28
0.02
0.27
0.14
0
0.01
0.34
0.07
Fig. 3. Continuous cooling transformation diagram for 11CrODS steel compared with those for 9CrODS and F82H steels.
boundaries is suppressed due to coarser size of the prior austenite grains [23]. If compared with 9CrODS steel [19], characteristic temperatures for the ferrite and martensite transformations are slightly lower in 11CrODS steel. This could be attributed to an austenite stabilization induced by nickel addition and smaller tungsten content in the 11CrODS steel.
3.2. Residual ferrite characterization
Fig. 4. Computed pseudo-binary Fee11Cr-1.3W-0.4Ni-0.3Ti phase diagram with variable carbon content.
Fig. 4 shows a computed phase diagram of 11CrODS steel in a pseudo-binary system of Fee11Cr-0.3Ti-0.4Ni-1.3W vs. carbon content without Y2O3, which was calculated by Pandat software [16,17]. At the standard composition of 0.13 wt% C, only g-phase is stable at normalizing temperature of 1050 C, which is consistent with Fig. 1. Following the CCT diagram (Fig. 3) and pseudo-binary phase diagram shown in Fig. 4, it is concluded that single ferrite
Fig. 5. Results of EBSD-image quarity (IQ) map, nanoindentation hardness and orientation distribution function (ODF) for fine grains and coarse grains in 11CrODS steel.
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T. Yamashiro et al. / Journal of Nuclear Materials 472 (2016) 247e251
Fig. 6. IQ map measured in several regions of the sample.
Fig. 7. Measured residual ferrite fractions in terms of free energy difference computed for a-g phases at 1050 C (J/mol).
phase is formed by the ferrite-transformation from single austenite phase, when 11CrODS steel is cooled down with the adequately slow cooling rate of 3 C/min. The EBSD-IQ map of the specimen cooled at 3 C/min is shown in Fig. 5(a). This structure does not correspond to single full ferrite, but it obviously consists of dual phases: coarser grains in white region and fine grains in dark region, where the hot-extruded direction is represented by the vertical arrow. In order to distinguish both phases, nano-indentation hardness was measured at 15 points, which are marked by small triangular shape of Berkovich indent in Fig. 5(a). The indents marked with circles correspond to the nanoindentation hardness marked with solid circles in Fig. 5(b). Obviously fine grain region by solid circles has higher hardness than coarse grain region. EBSD analyses were also conducted in the fine grain region and coarse grain region. Their orientation distribution function (ODF) maps are shown in Fig. 5 (c). The coarse grain region exhibits two series of peaks around {311}<011> and peaks toward {332}<113>. It is
T. Yamashiro et al. / Journal of Nuclear Materials 472 (2016) 247e251
known that both textures belong to the ferrite structure transformed from the austenite during slow cooling at 3 C/min. In contrast, the fine grain region shows the typical a-fiber texture; <110> direction is parallel to the hot-extruded direction and planes rotate 360 around <110> axis. The {111}<110> is prominent among the a-fiber texture. The fine grain regions were formed at the hot-extrusion processing, and they kept the same texture even followed by the normalizing heat treatment. Consequently, these results imply that the fine grain regions correspond to the residual ferrite. The same microstructure is reported by Cayron et al. in EUROFER ODS steels of 9Cr-0.1Ce1W-0.2Y2O3 (wt. %) [18]. The average fraction area of the residual ferrite was estimated by means of image analysis of IQ map at several regions (Fig. 6) and it was found to be 21%. This result is plotted as the solid circle in Fig. 7, where the horizontal axis represents the free energy difference between a and g phases at 1050 C. It was shown in the previous study [7] that the residual ferrite can be retained, when the pinning force for a-g interfacial boundaries by the oxide particles overcomes the driving force for the a to g reverse transformation. This driving force is closely related to the a/g free energy difference at 1050 C, Ga(1050 C)－Gg(1050 C); thus, the residual ferrite is stabilized with decreasing Ga－Gg. The detail chemical compositions of specimens indicated in Fig. 7 are listed in Ref. [15]. The residual ferrite fraction in each specimen was controlled by modifying contents of elements such as C, Cr, Ni, W and excess oxygen (Ex.O), and was estimated by means of the following three different methods as given in the ref. [15]. The X-ray intensity coming from ferrite was measured at 1050 C by X-ray diffraction (XRD) method. The W distribution, enriched in ferrite, was measured by electron probe microanalyzer (EPMA), because W element preferentially partitions into ferrite rather than austenite. Metallographic observation was also carried out after chemical etching. The ferrite fraction estimated by these three methods varies widely, however there is a tendency of increasing ferrite fraction with decreasing the a/g free energy difference, Ga(1050 C)－Gg(1050 C). The ferrite fraction derived from EBSD-IQ map in this study locates within the trend given in Ref. [15]. This implies that the estimation of fraction of residual ferrite from IQ map is reasonable. It is concluded that 11CrODS steel contains relatively higher fraction of the residual ferrite than 9CrODS steel. 4. Conclusions With the aim of further improvement of corrosion resistance, 11 wt% Cr ODS steels have been previously developed, which maintain martensitic base structure by adjusting content of alloying elements. The nominal chemical composition of 11CrODS steel is Fee11Cr-0.13C-1.3W-0.4Ni-0.3Ti- 0.35Y2O3 (wt%). In this study, a
251
continuous cooling transformation (CCT) diagram of 11CrODS steel was constructed, and its microstructure was evaluated with the aim to establish the basis of structural control and material processing. The CCT diagram constructed in terms of cooling rate showed that upper critical cooling rate inducing only martensite formation is 60 C/min (3600 C/h) while the lower cooling rate without martensite formation corresponds to 10 C/min (600 C/h). Characteristic temperatures for the ferrite and martensite transformations in 11CrODS steel are slightly lower than those of 9CrODS steel. From image analysis of IQ map, the area fraction of the residual ferrite, which is retained by pinning of a-g interfacial boundaries by oxide particles, was revealed to be 21% of the total area. This fraction is slightly higher than that obtained for 9CrODS steel, and it plays a key role in improving the high temperature strength of martensitic ODS steel.
References [1] Y. de Carlan, J.-L. Bechade, P. Dubuisson, et al., J. Nucl. Mater. 386e388 (2009) 430e432. [2] P. Dubuisson, Y. de Carlan, V. Garat, M. Blat, J. Nucl. Mater 428 (2012) 6e12. [3] G.R. Odette, M.J. Alinger, B.D. Wirth, Annu. Rev. Mater. Res. 38 (2008) 471e503. [4] R. Lindau, A. Moslang, M. Rieth, et al., Fusion Eng. Des. 75e79 (2005) 989e996. [5] S. Ukai, Y. Kudo, X. Wu, N. Oono, S. Hayashi, S. Ohtsuka, T. Kaito, J. Nucl. Mater 455 (2014) 700e703. [6] S. Ukai, M. Fujiwara, J. Nucl. Mater. 307e311 (2002) 749e757. [7] M. Yamamoto, S. Ukai, S. Hayashi, T. Kaito, Satoshi Ohtsuka, J. Nucl. Mater 417 (2011) 237e240. [8] S. Ukai, T. Kaito, S. Otsuka, T. Narita, M. Fujiwara, T. Kobayashi, ISIJ Int. 43 (12) (2003) 2038e2045. [9] S. Ohtsuka, S. Ukai, M. Fujiwara, H. Sakasegawa, T. Kaito, T. Narita, J. Nucl. Mater. 367e370 (2007) 160e165. [10] S. Ohtsuka, S. Ukai, M. Fujiwara, T. Kaito, T. Narita, J. Nucl. Mater. 329e333 (2004) 372e376. [11] S. Ohtsuka, S. Ukai, M. Fujiwara, J. Nucl. Mater 351 (2006) 241e246. [12] A. Steckmeyer, M. Praud, B. Fournier, et al., J. Nucl. Mater 405 (2010) 95e100. [13] T. Narita, S. Ukai, B. Leng, S. Ohtsuka, T. Kaito, J. Nucl. Scie. Technol. 50 (2013) 314e320. [14] B. Leng, S. Ukai, T. Narita, Y. Sugino, Q. Tang, N. Oono, S. Hayashi, F. Wan, S. Ohtsuka, T. Kaito, Mater. Trans. 53 (2012) 652e657. [15] S. Ohtsuka, T. Kaito, T. Tanno, Y. Yano, S. Koyama, K. Tanaka, J. Nucl. Mater. 442 (2013) S89eS94. [16] J.C. Zhao, Michael F. Henry, Adv. Eng. Mater 4 (2002) 501e508. [17] W. Cao, S.-L. Chen, F. Zhang, K. Wu, Y. Yang, Y.A. Chang, R. Schmid-Fetzer, W.A. Oates, Calphad 33 (2009) 328e342. [18] C. Cayron, E. Rath, I. Chu, S. Launois, J. Nucl. Mater 335 (2004) 83e102. [19] S. Ukai, S. Ohtsuka, Energy Mater 2 (1) (2007) 26e35. [20] M. Yamamoto, S. Ukai, S. Hayashi, T. Kaito, S. Ohtsuka, Mater. Sci. Eng. A 527 (2010) 4418e4423. [21] S. Ukai, S. Ohtsuka, T. Kaito, H. Sakasegawa, N. Chikata, S. Hayashi, Mater. Sci. Eng. A 510e511 (2009) 115e120. [22] R.L. Klueh, D.S. Gelles, S. Jitsukawa, et al., J. Nucl. Mater. 307e311 (2002) 455e465. [23] A. Alamo, A. Alamo, J.C. Brachet, et al., J. Nucl. Mater. 258e263 (1998) 1228e1235.
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Microstructural stability of 11Cr ODS steel* Tetsuya Yamashiro a, *, Shigeharu Ukai b, Naoko Oono b, Satoshi Ohtsuka c, Takeji Kaito c a
Materials Science and Engineering, Graduate School of Engineering, Hokkaido University, N13, W-8, Kita-ku, Sapporo 060-8628, Japan Materials Science and Engineering, Faculty of Engineering, Hokkaido University, N13, W-8, Kita-ku, Sapporo 060-8628, Japan c Advanced Nuclear System R&D Directorate, Japan Atomic Energy Agency (JAEA), 4002, Narita, Oarai, Ibaraki-pref. 311-1393, Japan b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 2 May 2015 Received in revised form 30 December 2015 Accepted 4 January 2016 Available online 6 January 2016
Aiming at further improvement of high-temperature oxidation and corrosion resistance, 11CrODS steel with martensitic base structure has been previously developed, as a candidate fuel cladding material for 4th generation advanced nuclear reactors. In this study, the microstructure of 11CrODS steel was characterized by means of EBSD and nanoindentation hardness measurement. The continuous cooling transformation (CCT) diagram was constructed. Upper critical cooling rate, which is minimum cooling rate necessary to form martensitic structure, was derived to be 60 C/min (3600 C/h). In contrast, lower critical cooling rate preventing from martensite formation, was derived to be 10 C/min (600 C/h). An area fraction of so called residual ferrite was estimated by image processing of EBSD-IQ map to be 21% of the total area. This fraction of the residual ferrite in 11CrODS steel was evaluated by considering the driving force for a to g reverse transformation. © 2016 Published by Elsevier B.V.
Keywords: Continuous cooling transformation diagram ODS Residual ferrite a-g phase transformation Pinning force
1. Introduction The oxide dispersion strengthened (ODS) ferritic steels have excellent radiation resistance and high-temperature strength. Therefore, they are promising candidates for cladding materials of Generation IV sodium-cooled fast reactors [1,2] and for blanket structural materials of the advanced fusion reactors [3,4]. The ODS ferritic steels have two types of structures known as alpha ferrite and martensite. We have focused on developing the 9CrODS steel with martensitic structure, since on one hand its structure and processing can be effectively controlled in terms of a-g phase transformation [5e11]. On the other hand, the structure and morphology of the 12CrODS steel as the typical ferritic steel can be only controlled by recrystallization [12]. However in this case there is a lack of repeatability and consequently this kind of a control is unsuitable to mass production [13,14]. In order to further improve a high-temperature oxidation and corrosion resistance in 9CrODS steel, we have developed a 11CrODS steel by increasing the Crcontent from 9 to 11 wt. % [15]. The 11CrODS steel has martensitic structure even at 11 wt. % Cr; thus it has the advantage for controlling the processing-fabrication route by a-g phase transformation. * Presented at the NuMat 2014 Conference, 27e30 October 2014, Clearwater, Florida, USA. * Corresponding author. E-mail address: [email protected] (T. Yamashiro).
http://dx.doi.org/10.1016/j.jnucmat.2016.01.002 0022-3115/© 2016 Published by Elsevier B.V.
A nominal chemical composition of the 11CrODS steel is Fee11Cr-0.13C-1.3W-0.4Ni- 0.3Ti-0.35Y2O3 (wt. %). We constructed phase stability map of the 11CrODS steel in a tungsten-nickel contents diagram, which is compared with those of 9CrODS and 12CrODS steels at 1050 C, keeping 0.3 wt. % Ti, 0.35 wt. % Y2O3 and Fe-balance. This is shown in Fig. 1, where 1050 C corresponds to the normalizing temperature of the ODS steels. These calculations were carried out by means of thermodynamic analysis software Pandat [16,17]. Each line for 9Cr, 11Cr and 12Cr indicates phase boundary between single g and g þ d two phase field. The single gphase is stabilized on the left side from the boundary, while the g þ d field is stable on its right side. In the case of 12CrODS steel, higher Ni-content is required at even lower content of tungsten, to produce the single g-austenite phase at 1050 C. With decreasing Cr-content to 11Cr and 9Cr, the Ni-content required for stabilization of the single g-austenite phase can be decreased. The chemical compositions of 11CrODS steel (1.3W-0.4Ni) and 9CrODS steel (2W without Ni addition) are represented by solid circle in Fig. 1, from which the single g-austenite is stable at 1050 C for 9CrODS steel; thus leading to the martensite formation after cooling. However, the structure of 9CrODS steels is characteristic for a dual phase composed of both martensite and ferrite, and this ferrite was designated as the metastable residual phase [8,18]. It is well known that the residual phase significantly improves the hightemperature strength [19e21]. For this reason, an estimation of
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T. Yamashiro et al. / Journal of Nuclear Materials 472 (2016) 247e251
Fig. 2. Length change of 11CrODS steel determined from dilatometric measurements during cooling with various cooling rate.
Fig. 1. Variation of phase diagram with Ni and W content.
1100a using a Berkovich type of indent, which separately gives a nanohardness of each phase.
the residual ferrite fraction is also important for 11CrODS steel. In this study, 11CrODS steel bars were manufactured, and their structural stability including continuous cooling transform (CCT) diagram were studied. A fraction of the metastable residual ferrite is also assessed from thermodynamic calculations.
3. Result & discussion 3.1. CCT diagram
2. Experimental procedure
The results of dilatometric measurement are shown in Fig. 2. The specimen length changes were recorded during cooling from 1050 C. Nine cooling rates were used, namely: 3 C/min, 6 C/min, 8 C/min, 10 C/min, 12 C/min, 30 C/min, 60 C/min, 300 C/min and 600 C/min. The symbols Ar3 and Ar1 correspond to temperatures of the beginning and of the end of the austenite to ferrite transformation, respectively. Ms and Mf denote temperatures of the beginning and of the end of the martensitic transformation. Temperatures Ar1, Ar3, Ms and Mf measured at each cooling rate are shown in Table 2. The lower critical cooling rate was assessed to be 10 C/min (600 C/h). Below this rate only ferrite transformation takes place. The upper critical cooling rate is found to be around 60 C/min (3600 C/h). For higher rates only martensite transformation occurs. The CCT diagram was constructed from the dilatometric data, and it is shown in Fig. 3. For comparison, diagrams for F82H and 9CrODS steels are also plotted. The F82H is the martensitic type of steel not containing oxide particles, and it is the promising candidate for the advanced fusion blanket structural materials [22]. In the case of F82H, the lower and upper critical cooling rates are shifted to the right side. It means that the martensite transformation easily proceeds even at lower cooling rate, because the ferrite transformation undergoing along prior austenite grain
The 11CrODS steel bar was manufactured by means of mechanical alloying and hot-extrusion. The pre-alloyed powders produced by gas atomizing of the nominal composition Fee11Cr0.13C-1.3W-0.4Ni-0.3Ti (wt. %) were supplied by Sanyo Special Steel Co. Ltd. The size of atomized powders is less than 100 mm, and they were mechanically alloyed together with Y2O3 powder by attrition type ball mill. The following conditions of ball milling were chosen: 220 rpm, 48 h and balls: powder ratio of 15: 1 (mass ratio). The mechanically alloyed powder was consolidated into bar by hotextrusion at 1150 C. The extrusion ratio is 9.2, and outer diameter of the bar is 24 mm. The result of chemical analysis of the extruded bar is shown in Table 1, where an excess oxygen (Ex.O) means the amount of total oxygen minus oxygen bounded with yttrium as Y2O3. The CCT diagram was constructed using a Rigaku model TMA8140C thermo mechanical analyzer. Rectangular shape of specimens in the dimension of 3 mm 3 mm 10 mm were used for this dilatometric studies. The specimen length change was measured along the extrusion direction during cooling from 1050 C with the cooling rate ranging from 3 C/min to 600 C/min. The specimens after dilatometric measurements were polished with colloidal silica. Electron back scattering diffraction (EBSD) measurement was carried out under the condition of accelerating voltage of 20 kV by JEOL JSM-6500F. The image quality (IQ) map was obtained to distinguish martensite from ferrite transformed from austenite, and from residual ferrite. An area of residual ferrite can be dark in IQ map, because the residual ferrite contains a large amount of strain, as compared with the transformed ferrite from austenite. Based on the above information, the average area fraction of the residual ferrite was estimated by means of the image processing of the low magnification EBSD-IQ map. Nanoindentation measurement was also conducted by ELIONIX ENT-
Table 2 Ar3, Ar1, Ms and Mf temperatures determined from dilatometric measurement at various cooling rate. Cooling rate ( C/min)
Ar3 ( C) Ar1 ( C) Ms ( C) Mf ( C)
3
6
8
10
12
30
60
300
600
756 692 e e
756 688 e e
748 672 e e
760 684 e e
728 640 390 275
720 620 405 220
e e 365 220
e e 355 214
e e 360 190
Table 1 Chemical composition of the manufactured 11CrODS steel specimen. Element
C
Si
Mn
P
S
Ni
Cr
W
Ti
Al
Y
O
N
Ar
Y2O3
Ex.O
wt%
0.13
0.04
0.04
0.01
0
0.4
10.9
1.3
0.28
0.02
0.27
0.14
0
0.01
0.34
0.07
Fig. 3. Continuous cooling transformation diagram for 11CrODS steel compared with those for 9CrODS and F82H steels.
boundaries is suppressed due to coarser size of the prior austenite grains [23]. If compared with 9CrODS steel [19], characteristic temperatures for the ferrite and martensite transformations are slightly lower in 11CrODS steel. This could be attributed to an austenite stabilization induced by nickel addition and smaller tungsten content in the 11CrODS steel.
3.2. Residual ferrite characterization
Fig. 4. Computed pseudo-binary Fee11Cr-1.3W-0.4Ni-0.3Ti phase diagram with variable carbon content.
Fig. 4 shows a computed phase diagram of 11CrODS steel in a pseudo-binary system of Fee11Cr-0.3Ti-0.4Ni-1.3W vs. carbon content without Y2O3, which was calculated by Pandat software [16,17]. At the standard composition of 0.13 wt% C, only g-phase is stable at normalizing temperature of 1050 C, which is consistent with Fig. 1. Following the CCT diagram (Fig. 3) and pseudo-binary phase diagram shown in Fig. 4, it is concluded that single ferrite
Fig. 5. Results of EBSD-image quarity (IQ) map, nanoindentation hardness and orientation distribution function (ODF) for fine grains and coarse grains in 11CrODS steel.
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Fig. 6. IQ map measured in several regions of the sample.
Fig. 7. Measured residual ferrite fractions in terms of free energy difference computed for a-g phases at 1050 C (J/mol).
phase is formed by the ferrite-transformation from single austenite phase, when 11CrODS steel is cooled down with the adequately slow cooling rate of 3 C/min. The EBSD-IQ map of the specimen cooled at 3 C/min is shown in Fig. 5(a). This structure does not correspond to single full ferrite, but it obviously consists of dual phases: coarser grains in white region and fine grains in dark region, where the hot-extruded direction is represented by the vertical arrow. In order to distinguish both phases, nano-indentation hardness was measured at 15 points, which are marked by small triangular shape of Berkovich indent in Fig. 5(a). The indents marked with circles correspond to the nanoindentation hardness marked with solid circles in Fig. 5(b). Obviously fine grain region by solid circles has higher hardness than coarse grain region. EBSD analyses were also conducted in the fine grain region and coarse grain region. Their orientation distribution function (ODF) maps are shown in Fig. 5 (c). The coarse grain region exhibits two series of peaks around {311}<011> and peaks toward {332}<113>. It is
T. Yamashiro et al. / Journal of Nuclear Materials 472 (2016) 247e251
known that both textures belong to the ferrite structure transformed from the austenite during slow cooling at 3 C/min. In contrast, the fine grain region shows the typical a-fiber texture; <110> direction is parallel to the hot-extruded direction and planes rotate 360 around <110> axis. The {111}<110> is prominent among the a-fiber texture. The fine grain regions were formed at the hot-extrusion processing, and they kept the same texture even followed by the normalizing heat treatment. Consequently, these results imply that the fine grain regions correspond to the residual ferrite. The same microstructure is reported by Cayron et al. in EUROFER ODS steels of 9Cr-0.1Ce1W-0.2Y2O3 (wt. %) [18]. The average fraction area of the residual ferrite was estimated by means of image analysis of IQ map at several regions (Fig. 6) and it was found to be 21%. This result is plotted as the solid circle in Fig. 7, where the horizontal axis represents the free energy difference between a and g phases at 1050 C. It was shown in the previous study [7] that the residual ferrite can be retained, when the pinning force for a-g interfacial boundaries by the oxide particles overcomes the driving force for the a to g reverse transformation. This driving force is closely related to the a/g free energy difference at 1050 C, Ga(1050 C)－Gg(1050 C); thus, the residual ferrite is stabilized with decreasing Ga－Gg. The detail chemical compositions of specimens indicated in Fig. 7 are listed in Ref. [15]. The residual ferrite fraction in each specimen was controlled by modifying contents of elements such as C, Cr, Ni, W and excess oxygen (Ex.O), and was estimated by means of the following three different methods as given in the ref. [15]. The X-ray intensity coming from ferrite was measured at 1050 C by X-ray diffraction (XRD) method. The W distribution, enriched in ferrite, was measured by electron probe microanalyzer (EPMA), because W element preferentially partitions into ferrite rather than austenite. Metallographic observation was also carried out after chemical etching. The ferrite fraction estimated by these three methods varies widely, however there is a tendency of increasing ferrite fraction with decreasing the a/g free energy difference, Ga(1050 C)－Gg(1050 C). The ferrite fraction derived from EBSD-IQ map in this study locates within the trend given in Ref. [15]. This implies that the estimation of fraction of residual ferrite from IQ map is reasonable. It is concluded that 11CrODS steel contains relatively higher fraction of the residual ferrite than 9CrODS steel. 4. Conclusions With the aim of further improvement of corrosion resistance, 11 wt% Cr ODS steels have been previously developed, which maintain martensitic base structure by adjusting content of alloying elements. The nominal chemical composition of 11CrODS steel is Fee11Cr-0.13C-1.3W-0.4Ni-0.3Ti- 0.35Y2O3 (wt%). In this study, a
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continuous cooling transformation (CCT) diagram of 11CrODS steel was constructed, and its microstructure was evaluated with the aim to establish the basis of structural control and material processing. The CCT diagram constructed in terms of cooling rate showed that upper critical cooling rate inducing only martensite formation is 60 C/min (3600 C/h) while the lower cooling rate without martensite formation corresponds to 10 C/min (600 C/h). Characteristic temperatures for the ferrite and martensite transformations in 11CrODS steel are slightly lower than those of 9CrODS steel. From image analysis of IQ map, the area fraction of the residual ferrite, which is retained by pinning of a-g interfacial boundaries by oxide particles, was revealed to be 21% of the total area. This fraction is slightly higher than that obtained for 9CrODS steel, and it plays a key role in improving the high temperature strength of martensitic ODS steel.
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