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Improvement of high temperature strength and low temperature toughness of high manganese-chromium austenitic steels
EISEVIER
Journai of Nuclear Materials 212-215 (1994) 766-771
Improvement of high temperature strength and low temperature toughness of high manganese-chromium austenitic steels Kazuya Miyahara a, Dong-Su Bae b, Yukio Shimoide ’ a Department of Materials Science and Engineering, Nagoya University, Furoh-cho, Chikusa-ku, Nagoya 464, Japan b Graduate School, Nagoya University, ’ ~~ariment of Mechanical Engineering, Daido ~~titute of Tec~no~~~, Daido-cho, ~inami-~ Nagoya 457, Japan
Abstract High Mn-Cr austenitic steels are still considered to be an important high temperature structural material from the point of view of fast-induced radioactivity decay (FIRD) and non-magneticity. The objective of the present study is to investigate the mechanical properties of 12% Cr-15% Mn austenitic stainless steels and to compare these properties with those of the reference materials of JPCAs and JFMS, which are being investigated for the development of fusion reactor structural materials in Japan. The effects of the alloying elements V, Ti, Ta, etc. were investigated to determine the improvement of mechanical properties. Tiny precipitates of VN and Ti(C, N) raised the high-temperature strength considerably. Content of 0.1 to 0.2% C, however, formed very coarse precipitates of M,C, type carbide on the grain boundaries, which deteriorated low temperature toughness inducing intergranular fracture. Microst~ctural evolution during long-term aging was also investigated.
1. introduction
austenitic stainless steels) and JFMS (a 10% Cr ferritic-martensitic steel).
In the 1950s to 197Os, high Mn-Cr heat-resisting steels were developed [l-3] from the point of view of the saving of nickel resources and material cost. Recently, however, much attention has been paid on the development of a high Mn-Cr steel as a structural material for a fusion reactor for the attainment of fast-induced radioactive decay (FIRD [4,SJ). Accordingly, fundamental studies are being made on the various aspects of the high Mn-Cr alloys and steels, such as effects of neutron irradiation on microstructural evolution [6,7], phase stability [8,9], elevated temperature strength [lo-121 and low temperature toughness [131.
In the present study, the effects of C, N, V, Ti, Ta and P on the mechanical properties (creep-rupture strength and low temperature toughness) and microstructural characteristics of 12% Cr-15% Mn austenitic steels have been investigated. The results of strength measurements are compared with those of the reference materials of JPCAs (modified type 316 00%3115/94/%07.00
2. Experimental procedures The chemical compositions of materials are shown in Table 1, The addition of 0.9% Ni is based on the ~nsideration of the effect of nickel on the ductility of Mn-Cr stainless steels, where nickel has a concentration limit of 1.0% [S] under shallow burial guidelines for radioactive waste (US NRC Proposed Rule lOCFR61, Class Cl. These alloys were melted in a vacuum induction furnace. Each alloy was hot-rolled and solution-treated at 1373 K for 1 h. After the solution treatment, some specimens were used for microstructure observation, tensile test, creep-rupture test and Charpy impact test, and the other specimens were aged at 873 K for 1 to 8000 h before measurements and observation. The chemical compositions of the reference materials of JPCA, JPCA-2 and JFMS are shown in Table 2.
0 1994 Elsevier Science S.V. AI1 rights resewed
SSDI 0022-3115(93)E0379-N
K Miyaharaet al. /Journal of Nuclear Materials212-215 (1994) 764-771
767
JPCA and JPCA-2 are solution treated at 1363 K for 1 h and JFMS is normalized at 1323 K for 0.5 h and tempered at 1048 K for 1 h.
3. Experimental results and discussions 3.1. Microstructural evolution during long-term aging Fig. 1 shows the Vickers hardness change during the aging treatment at 873 K. Large age-hardening peaks for V-containing materials at 1000 h are caused by the fine precipitates of VN in the matrix, as shown in Fig. 2a. Small age-hardening peaks for V-free materials at aging times of 1 to 10 h are due to the M&-type carbide formed on grain boundaries. The M,C, carbide also formed on grain boundaries in the V-containing materials. Ti-containing materials also show large age-hardening peaks at about 1000 h aging, which are attributed to tiny precipitates of Ti(C, N) shown in Fig. 2b. Ta-containing material did not indicate any hardening at about 500 to 1000 h aging. This is explained from the TEM observation (Fig. 2c) that show coarse and spherical precipitates of TaN that formed, but precipitation of tiny VN or Ti(C, N) are restrained after a 1000 h aging treatment. Fig. 1 also indicates that all materials show a remarkable increase of hardness during long-term aging for more than 3000 h. Since the formation of many rectangular and coarse M,C, carbides are observed in the matrix of materials aged for 8000 h, as shown in Fig. 3, such increase in hardness is considered to be
0
073 K PClNW
0
lC2NWi
ged at
Fig. 2. TEM photographs of (a) lC2NWV, (b) lC2NWVTi and (c) lC2NWVTiTa aged at 873 K for 1000 h.
Aging time / ks
Fig. 1. Vickers hardness of the 12Cr-15Mn steels as functions of aging time at 873 K.
due to these precipitates. A few u-phase particles were also observed in the materials aged for 8000 h [ll]. It is interesting to note that JPCA-2 shows a similar increase in hardness after long-term aging, as shown in Fig. 1. Weiss et al. reported that a type 316 steel also showed the pronounced increase in hardness after long-term aging, and they attributed the increase to intragranular precipitates of M,C, carbide [14]. Since the microstructural evolution during long-term aging is important for high-temperature materials, detailed
768
I: Mi~ahara et al. /Journal
~~~~~eaF l~ater~als 212-215
11994f 1616-771
Table 1 Chemical composition of materials (mass%) N
C 2ClN 2ClNw 2ClNwV 2ClNWVTi lC2N 1c2Nw lC2NwV 1CZNWVTi 1CZNWVTiTa
0.20 0.19 0.20 0.19 0.10 0.089 0.095 0.10 0.10
0.10 0.13 0.13 0.13 0.24 0.15 0.216 0.18 0.22
Si 0.10 0.10 0.003 0.10 0.10 0.003 0.073 < 0.10 < 0.01 < < < < <
Mn 15.3 16.2 15.36 15.1 15.53 15.23 14.97 15.1 14.78
P 0.001 < 0.003 < 0.003 < 0.003 0.002 < 0.003 0.002 < 0.003 0.003
comparison of such aging behavior between the 12% Cr-15% Mn steels and type 316 steels, including modified steels, is necessary. 3.2. Creep-rupture
strength
Fig. 4 shows the creep curves at 873 K under a stress of 216 MPa. The addition of V is effective in
S 0.004 0.006 0.006 0.004 0.005 0.004 0.004 0.005 0.006
Cr 12.20 9.40 9.60 11.97 12.57 9.50 10.72 11.85 10.51
Ni 0.91 0.84 0.88 0.86 0.96 0.86 0.88 0.84 0.87
W < 0.005 1.84
V
Ti
Ta
-
_
-
0.51 0.48
0.16
-
1.83 2.18 _ 1.80
-
_
2.01 2.15 1.90
0.62 0.46 0.23
0.10 0.19
0.55
raising the creep strength in comparison with V-free materials of 2ClNW and lC2NW. The combined addition of V and Ti is very effective for improving creep strength. Such strengthening effects of V and Ti are due to the formation of fine precipitates of VN or Ti(C, N) as shown in Figs. 2a and 2b. The addition of 0.5% Ta, however, was not effective in raising the
d
Fig. 3. TEM photographs of 2ClNWV aged at 873 K for 8000 h. (a) Bright-field image, (b) dark-field image of Mz3C6, (c) electron diffraction pattern of M,,C, and y-phase, and (d) indices of the diffraction pattern of (c).
K Miyaharaet al./Journal of Nuclear Materials212-215 (1994) 766-771 Time I h 2000
4000
I ’ I ’ I I,,
I
’
I
6000 ’
Time to
,
6000 I
’
IV,
10000
a,
10’ ‘I’ 1’
5cK
8
169 ruptureI h
ld
103
104
\ ‘,‘1”1 ’ “,‘,‘I
1 “‘,,‘I
102
104
105
’ ‘*‘,“I
’
873K
1 C2N 4oc
g
3oc
z $ * P "n 2oc 4"
1C2NWVTiTa
1 C2NWV
1OC
I ,,,,
I
10’
10000
103
105
106
Time to rupture / ks
,1,,,,,,,,,,,,,,,,,,,,,,,,,,,
20000
30000
40000
Fig. 5. Relationship between rupture time and applied stress for the 12Cr-15Mn steels and the reference materials. (The rupture strength of JPCA-2 was obtained from the LarsonMiller plot of Asakura’s data [15] with rupture time of 5 to 2 x lo4 h at 923 K.)
Time I ks
Fig. 4. Creep curves of the 12Cr-15Mn steels tested at 873 K under an applied stress of 216 MPa.
3.3. Low-temperature toughness and fracture mode creep strength because very coarse precipitates of TaN, which are shown in Fig. 2c, formed, and the precipitation of tiny VN and Ti(C, N) were restrained, as described in section 3.1. Fig. 5 shows the creep-rupture strength of various 12Cr-15Mn materials at 873 K. The V-containing and Ti-containing materials (2ClNWV, lC2NWV and 2ClNWVTi) show the largest creep-rupture strengths among the 12Cr-15Mn materials. These rupture strengths, however, are a little lower than that of JPCA-2 1151(including JPCA [161X However, they are higher than those of ordinary type 316 stainless steel [17] and JFMS [18].
Fig. 6 shows the temperature dependence of toughness of the materials, including the results of JPCA-2 and JFMS. As for the 12% Cr-15% Mn austenitic steels, even the solution-treated materials indicate a considerable temperature dependence of Charpy impact value. Such large temperature dependence is in contrast to that of JPCA-2. Observation with an optical microscope and transmission electron microscope indicated the existence of a large volume fraction of Emartensite and twins in the 12% Cr-15% Mn steels, as shown in Figs. 7 and 8. Accordingly, it is concluded that such l-martensite and twins caused the large tem-
Table 2 Chemical composition of reference materials of austenitic and ferritic steels (mass%) Austenitic steel a
JPCA JPCA-2
C
Si
Mn
P
S
Cr
Ni
MO
Ca
Ti
B
0.06 0.055
0.50 0.53
1.50 1.88
0.025 0.024
0.009
14.00 15.27
16.00 15.80
2.50 2.66
0.048
0.25 0.24
0.0040 0.0032
C
Si
Mn
P
S
Cr
Ni
MO
V
Nb
N
0.05
0.58
0.58
0.009
0.006
9.85
0.94
2.31
0.12
0.06
0.01
Ferritic steel b
JFMS
a 1363 KST. b1323Kx1.8ksAC+1048Kx3.6ksAC.
770
K. ~iyahar~ et uf. ~~~ur~~l
of Nuclear ~ateriu~s 212-215
(‘19941
to the e-martensite, VN and T#C, N).
164-771
twins and/or
tiny precipitates
of
The Ta-containing materiat showed the iargest impact value among the 12% Cr-15% Mn steels. Kn the Ta-containing material, as described above, coarse and spherical precipitates of TaN are formed, restraining the precipitation of tiny VN or Ti(C, N). Such precipitation behavior is considered to cause the ductile and partially cleavage fracture at all test temperature. The addition of P degraded the toughness of the aged material. From the above results, it is noted that the combination of a moderate fraction of VN, Ti(C, N) and TaN due to adjusting the contents of N, V, Ti and Ta with reducing C content is important for the improvement of toughness and high temperature strength. 4oc Test temperature
/
K
Fig. 6. Effect of test temperature bn Charpy impact value of the 12Cr-15Mn steels. (0): Normalized and tempered [HI.
perature dependence of toughness of the 12% Cr-15% Mn steels. The fracture mode of the 316 steel and JPCA-2 was transgranular ductile fracture. On the other hand, soiution-treated 12% Cr-1.5% Mn steels showed a combined fracture mode of cleavage (hereafter, ‘cleavage’ means cleavage or quasi cleavage) and ductile fracture at all test temperatures. As for the aged 12% Cr-15% Mn steels, all the materials showed a combined fracture mode of cleavage and intergranular below 273 K, except for the Ta-containing material, which showed ductile and partially cleavage fracture at all temperature. The intergranular fracture of the aged 12% Cr-15% Mn steels is due to the coarse precipitates of M,C, carbides on the grain boundaries. Cleavage fracture is considered to be due
4. Summary
The following summarizes the present study: (1) Addition of V and Ti raised creep-rupture strength considerably due to the tiny precipitates of VN and Ti(C, N). Combined addition of V and Ti lead the largest creep-rupture strength of the 12Cr-15Mn steels. The highest strength is a little lower than that of JPCAs but much higher than that of type 316 steel and JFMS. (2) A high content of 0.1 to 0.2% C resulted in very coarse precipitates of M,,C, carbide on grain boundaries and caused intergranular fracture even at near room temperature. Therefore, C content should be reduced. (3) Addition of V lowered the toughness because of the fine precipitates of VN that caused cleavage fracture at low temperature. P also reduced lowtemperature toughness. Ta improved the toughness by the precipitation of coarse and spherical TaN.
Fig. 7. Optical micrographs of (a) 2ClNW and (b) 2ClNWV solution-treated
and cooled to 77 K.
K Miyahara et al. /Journal
of Nuclear Materials 212-215
(1994) 766-771
Fig. 8. r-martensite in 2ClNW solution-treated and cooled to 77 K. (a) Bright-field image, (b) dark-field image from the spot, (cl diffraction pattern of matrix and e-martensite, and (d) indices of diffraction spots of e-martensite.
Combined additions of V and Ti improved the toughness. Accordingly, it is necessary to obtain the combination of a moderate fraction of VN, Ti(C, N) and TaN to increase toughness while keeping the high temperature strength.
Acknowledgement
The present authors are grateful to Dr. T. Nakazawa, Steel Research Laboratories, Nippon Steel Co., Ltd. for supplying the materials for the present study.
References
Ill R.R. Brady and D.T. Boughner, The Iron Age 184 (1959) 90.
Dl F.C. Hull, Mater. Des. Eng. 59 (1964) 92. 131 R. Choubey, R.B. Gibbons and K. Prasad, Metals Technol. (1977) 524. [41 R.L. Klueh and P.J. Maziasz, Mater. Sci. Eng. Al27 (1990) 17.
771
l(lOi2)
I51D.G. Doran, H.L. Heinisch and F.M. Mann, J. Nucl. Mater. 133&134 (198.5) 582. &I F.A. Gamer, H.R. Brager, D.S. Gelles and J.M. McCarthy, J. Nucl. Mater. 148 (1987) 294. [II K. Miyahara, F.A. Gamer and Y. Hosoi, J. Nucl. Mater. 191-194 (1992) 1198. @I F.A. Gamer, F. Abe and T. Noda, J. Nucl. Mater. 155-157 (1988) 870. I91 Y. Okazaki, K. Miyahara, Y. Hosoi, M. Tanino and H. Komatsu, J. Jpn. inst. Metals 53 (1989) 964. DO1 K. Miyahara, D.S. Bae, H. Sakai, J. Inami and Y. Hosoi, Tetsu to Hagane 78 (1992) 1737. Dll K. Miyahara, D.S. Bae, H. Sakai and Y. Hosoi, Mater. Trans. Jpn. Inst. Metals 34 (1993) 1036. WI F. Abe, T. Noda, H. Araki and S. Nakazawa, J. Nucl. Mater. 191-194 (1992) 668. [131 Y. Shimoide, K. Miyahara, D.S. Bae, H. Katoh and Y. Hosoi, Mater. Trans. Jpn. Inst. Metals 34 (1993) 1006. 1141B. Weiss and R. SticMer, Metall. Trans. 3 (1972) 851. 1151 K. Asakura and T. Fujita, Report of the 123rd Committee on Heat-Resisting Metals and Alloys, Japanese Society for the Promotion of Science, vol. 28 (1987) p. 41. 1161 M. Fujiwara, private communication. u71 NRIM Creep Data Sheet 14B (1988). Ml K. Asakura and T. Fujita, Report of the 123rd Committee on Heat-Resisting Metals and Alloys, Japanese Society for the Promotion of Science, vol. 27 (1986) pp. 93-104.
Journai of Nuclear Materials 212-215 (1994) 766-771
Improvement of high temperature strength and low temperature toughness of high manganese-chromium austenitic steels Kazuya Miyahara a, Dong-Su Bae b, Yukio Shimoide ’ a Department of Materials Science and Engineering, Nagoya University, Furoh-cho, Chikusa-ku, Nagoya 464, Japan b Graduate School, Nagoya University, ’ ~~ariment of Mechanical Engineering, Daido ~~titute of Tec~no~~~, Daido-cho, ~inami-~ Nagoya 457, Japan
Abstract High Mn-Cr austenitic steels are still considered to be an important high temperature structural material from the point of view of fast-induced radioactivity decay (FIRD) and non-magneticity. The objective of the present study is to investigate the mechanical properties of 12% Cr-15% Mn austenitic stainless steels and to compare these properties with those of the reference materials of JPCAs and JFMS, which are being investigated for the development of fusion reactor structural materials in Japan. The effects of the alloying elements V, Ti, Ta, etc. were investigated to determine the improvement of mechanical properties. Tiny precipitates of VN and Ti(C, N) raised the high-temperature strength considerably. Content of 0.1 to 0.2% C, however, formed very coarse precipitates of M,C, type carbide on the grain boundaries, which deteriorated low temperature toughness inducing intergranular fracture. Microst~ctural evolution during long-term aging was also investigated.
1. introduction
austenitic stainless steels) and JFMS (a 10% Cr ferritic-martensitic steel).
In the 1950s to 197Os, high Mn-Cr heat-resisting steels were developed [l-3] from the point of view of the saving of nickel resources and material cost. Recently, however, much attention has been paid on the development of a high Mn-Cr steel as a structural material for a fusion reactor for the attainment of fast-induced radioactive decay (FIRD [4,SJ). Accordingly, fundamental studies are being made on the various aspects of the high Mn-Cr alloys and steels, such as effects of neutron irradiation on microstructural evolution [6,7], phase stability [8,9], elevated temperature strength [lo-121 and low temperature toughness [131.
In the present study, the effects of C, N, V, Ti, Ta and P on the mechanical properties (creep-rupture strength and low temperature toughness) and microstructural characteristics of 12% Cr-15% Mn austenitic steels have been investigated. The results of strength measurements are compared with those of the reference materials of JPCAs (modified type 316 00%3115/94/%07.00
2. Experimental procedures The chemical compositions of materials are shown in Table 1, The addition of 0.9% Ni is based on the ~nsideration of the effect of nickel on the ductility of Mn-Cr stainless steels, where nickel has a concentration limit of 1.0% [S] under shallow burial guidelines for radioactive waste (US NRC Proposed Rule lOCFR61, Class Cl. These alloys were melted in a vacuum induction furnace. Each alloy was hot-rolled and solution-treated at 1373 K for 1 h. After the solution treatment, some specimens were used for microstructure observation, tensile test, creep-rupture test and Charpy impact test, and the other specimens were aged at 873 K for 1 to 8000 h before measurements and observation. The chemical compositions of the reference materials of JPCA, JPCA-2 and JFMS are shown in Table 2.
0 1994 Elsevier Science S.V. AI1 rights resewed
SSDI 0022-3115(93)E0379-N
K Miyaharaet al. /Journal of Nuclear Materials212-215 (1994) 764-771
767
JPCA and JPCA-2 are solution treated at 1363 K for 1 h and JFMS is normalized at 1323 K for 0.5 h and tempered at 1048 K for 1 h.
3. Experimental results and discussions 3.1. Microstructural evolution during long-term aging Fig. 1 shows the Vickers hardness change during the aging treatment at 873 K. Large age-hardening peaks for V-containing materials at 1000 h are caused by the fine precipitates of VN in the matrix, as shown in Fig. 2a. Small age-hardening peaks for V-free materials at aging times of 1 to 10 h are due to the M&-type carbide formed on grain boundaries. The M,C, carbide also formed on grain boundaries in the V-containing materials. Ti-containing materials also show large age-hardening peaks at about 1000 h aging, which are attributed to tiny precipitates of Ti(C, N) shown in Fig. 2b. Ta-containing material did not indicate any hardening at about 500 to 1000 h aging. This is explained from the TEM observation (Fig. 2c) that show coarse and spherical precipitates of TaN that formed, but precipitation of tiny VN or Ti(C, N) are restrained after a 1000 h aging treatment. Fig. 1 also indicates that all materials show a remarkable increase of hardness during long-term aging for more than 3000 h. Since the formation of many rectangular and coarse M,C, carbides are observed in the matrix of materials aged for 8000 h, as shown in Fig. 3, such increase in hardness is considered to be
0
073 K PClNW
0
lC2NWi
ged at
Fig. 2. TEM photographs of (a) lC2NWV, (b) lC2NWVTi and (c) lC2NWVTiTa aged at 873 K for 1000 h.
Aging time / ks
Fig. 1. Vickers hardness of the 12Cr-15Mn steels as functions of aging time at 873 K.
due to these precipitates. A few u-phase particles were also observed in the materials aged for 8000 h [ll]. It is interesting to note that JPCA-2 shows a similar increase in hardness after long-term aging, as shown in Fig. 1. Weiss et al. reported that a type 316 steel also showed the pronounced increase in hardness after long-term aging, and they attributed the increase to intragranular precipitates of M,C, carbide [14]. Since the microstructural evolution during long-term aging is important for high-temperature materials, detailed
768
I: Mi~ahara et al. /Journal
~~~~~eaF l~ater~als 212-215
11994f 1616-771
Table 1 Chemical composition of materials (mass%) N
C 2ClN 2ClNw 2ClNwV 2ClNWVTi lC2N 1c2Nw lC2NwV 1CZNWVTi 1CZNWVTiTa
0.20 0.19 0.20 0.19 0.10 0.089 0.095 0.10 0.10
0.10 0.13 0.13 0.13 0.24 0.15 0.216 0.18 0.22
Si 0.10 0.10 0.003 0.10 0.10 0.003 0.073 < 0.10 < 0.01 < < < < <
Mn 15.3 16.2 15.36 15.1 15.53 15.23 14.97 15.1 14.78
P 0.001 < 0.003 < 0.003 < 0.003 0.002 < 0.003 0.002 < 0.003 0.003
comparison of such aging behavior between the 12% Cr-15% Mn steels and type 316 steels, including modified steels, is necessary. 3.2. Creep-rupture
strength
Fig. 4 shows the creep curves at 873 K under a stress of 216 MPa. The addition of V is effective in
S 0.004 0.006 0.006 0.004 0.005 0.004 0.004 0.005 0.006
Cr 12.20 9.40 9.60 11.97 12.57 9.50 10.72 11.85 10.51
Ni 0.91 0.84 0.88 0.86 0.96 0.86 0.88 0.84 0.87
W < 0.005 1.84
V
Ti
Ta
-
_
-
0.51 0.48
0.16
-
1.83 2.18 _ 1.80
-
_
2.01 2.15 1.90
0.62 0.46 0.23
0.10 0.19
0.55
raising the creep strength in comparison with V-free materials of 2ClNW and lC2NW. The combined addition of V and Ti is very effective for improving creep strength. Such strengthening effects of V and Ti are due to the formation of fine precipitates of VN or Ti(C, N) as shown in Figs. 2a and 2b. The addition of 0.5% Ta, however, was not effective in raising the
d
Fig. 3. TEM photographs of 2ClNWV aged at 873 K for 8000 h. (a) Bright-field image, (b) dark-field image of Mz3C6, (c) electron diffraction pattern of M,,C, and y-phase, and (d) indices of the diffraction pattern of (c).
K Miyaharaet al./Journal of Nuclear Materials212-215 (1994) 766-771 Time I h 2000
4000
I ’ I ’ I I,,
I
’
I
6000 ’
Time to
,
6000 I
’
IV,
10000
a,
10’ ‘I’ 1’
5cK
8
169 ruptureI h
ld
103
104
\ ‘,‘1”1 ’ “,‘,‘I
1 “‘,,‘I
102
104
105
’ ‘*‘,“I
’
873K
1 C2N 4oc
g
3oc
z $ * P "n 2oc 4"
1C2NWVTiTa
1 C2NWV
1OC
I ,,,,
I
10’
10000
103
105
106
Time to rupture / ks
,1,,,,,,,,,,,,,,,,,,,,,,,,,,,
20000
30000
40000
Fig. 5. Relationship between rupture time and applied stress for the 12Cr-15Mn steels and the reference materials. (The rupture strength of JPCA-2 was obtained from the LarsonMiller plot of Asakura’s data [15] with rupture time of 5 to 2 x lo4 h at 923 K.)
Time I ks
Fig. 4. Creep curves of the 12Cr-15Mn steels tested at 873 K under an applied stress of 216 MPa.
3.3. Low-temperature toughness and fracture mode creep strength because very coarse precipitates of TaN, which are shown in Fig. 2c, formed, and the precipitation of tiny VN and Ti(C, N) were restrained, as described in section 3.1. Fig. 5 shows the creep-rupture strength of various 12Cr-15Mn materials at 873 K. The V-containing and Ti-containing materials (2ClNWV, lC2NWV and 2ClNWVTi) show the largest creep-rupture strengths among the 12Cr-15Mn materials. These rupture strengths, however, are a little lower than that of JPCA-2 1151(including JPCA [161X However, they are higher than those of ordinary type 316 stainless steel [17] and JFMS [18].
Fig. 6 shows the temperature dependence of toughness of the materials, including the results of JPCA-2 and JFMS. As for the 12% Cr-15% Mn austenitic steels, even the solution-treated materials indicate a considerable temperature dependence of Charpy impact value. Such large temperature dependence is in contrast to that of JPCA-2. Observation with an optical microscope and transmission electron microscope indicated the existence of a large volume fraction of Emartensite and twins in the 12% Cr-15% Mn steels, as shown in Figs. 7 and 8. Accordingly, it is concluded that such l-martensite and twins caused the large tem-
Table 2 Chemical composition of reference materials of austenitic and ferritic steels (mass%) Austenitic steel a
JPCA JPCA-2
C
Si
Mn
P
S
Cr
Ni
MO
Ca
Ti
B
0.06 0.055
0.50 0.53
1.50 1.88
0.025 0.024
0.009
14.00 15.27
16.00 15.80
2.50 2.66
0.048
0.25 0.24
0.0040 0.0032
C
Si
Mn
P
S
Cr
Ni
MO
V
Nb
N
0.05
0.58
0.58
0.009
0.006
9.85
0.94
2.31
0.12
0.06
0.01
Ferritic steel b
JFMS
a 1363 KST. b1323Kx1.8ksAC+1048Kx3.6ksAC.
770
K. ~iyahar~ et uf. ~~~ur~~l
of Nuclear ~ateriu~s 212-215
(‘19941
to the e-martensite, VN and T#C, N).
164-771
twins and/or
tiny precipitates
of
The Ta-containing materiat showed the iargest impact value among the 12% Cr-15% Mn steels. Kn the Ta-containing material, as described above, coarse and spherical precipitates of TaN are formed, restraining the precipitation of tiny VN or Ti(C, N). Such precipitation behavior is considered to cause the ductile and partially cleavage fracture at all test temperature. The addition of P degraded the toughness of the aged material. From the above results, it is noted that the combination of a moderate fraction of VN, Ti(C, N) and TaN due to adjusting the contents of N, V, Ti and Ta with reducing C content is important for the improvement of toughness and high temperature strength. 4oc Test temperature
/
K
Fig. 6. Effect of test temperature bn Charpy impact value of the 12Cr-15Mn steels. (0): Normalized and tempered [HI.
perature dependence of toughness of the 12% Cr-15% Mn steels. The fracture mode of the 316 steel and JPCA-2 was transgranular ductile fracture. On the other hand, soiution-treated 12% Cr-1.5% Mn steels showed a combined fracture mode of cleavage (hereafter, ‘cleavage’ means cleavage or quasi cleavage) and ductile fracture at all test temperatures. As for the aged 12% Cr-15% Mn steels, all the materials showed a combined fracture mode of cleavage and intergranular below 273 K, except for the Ta-containing material, which showed ductile and partially cleavage fracture at all temperature. The intergranular fracture of the aged 12% Cr-15% Mn steels is due to the coarse precipitates of M,C, carbides on the grain boundaries. Cleavage fracture is considered to be due
4. Summary
The following summarizes the present study: (1) Addition of V and Ti raised creep-rupture strength considerably due to the tiny precipitates of VN and Ti(C, N). Combined addition of V and Ti lead the largest creep-rupture strength of the 12Cr-15Mn steels. The highest strength is a little lower than that of JPCAs but much higher than that of type 316 steel and JFMS. (2) A high content of 0.1 to 0.2% C resulted in very coarse precipitates of M,,C, carbide on grain boundaries and caused intergranular fracture even at near room temperature. Therefore, C content should be reduced. (3) Addition of V lowered the toughness because of the fine precipitates of VN that caused cleavage fracture at low temperature. P also reduced lowtemperature toughness. Ta improved the toughness by the precipitation of coarse and spherical TaN.
Fig. 7. Optical micrographs of (a) 2ClNW and (b) 2ClNWV solution-treated
and cooled to 77 K.
K Miyahara et al. /Journal
of Nuclear Materials 212-215
(1994) 766-771
Fig. 8. r-martensite in 2ClNW solution-treated and cooled to 77 K. (a) Bright-field image, (b) dark-field image from the spot, (cl diffraction pattern of matrix and e-martensite, and (d) indices of diffraction spots of e-martensite.
Combined additions of V and Ti improved the toughness. Accordingly, it is necessary to obtain the combination of a moderate fraction of VN, Ti(C, N) and TaN to increase toughness while keeping the high temperature strength.
Acknowledgement
The present authors are grateful to Dr. T. Nakazawa, Steel Research Laboratories, Nippon Steel Co., Ltd. for supplying the materials for the present study.
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