- Email: [email protected]
Evolution of microstructure and acceleration of creep rate in tempered martensitic 9Cr-W steels
ELSEVIER
Materials
Evolution
Science
and Engineering
A234-236
(1997)
1045-1048
of microstructure and acceleration of creep rate in tempered martensitic 9Cr-W steels Fujio Abe * National
Research
Received
Institute
27 February
for
Metals,
1997; received
1-2-1
Sengen,
in revised
Tsukuba
form
305, Japan
28 March
1997
Abstract The effect of microstructural evolution on creep rates in the acceleration creep region has been investigated for tempered martensitic 9Cr-W steels. Creep tests were carried out at 823, 873 and 923 K for up to 15000 h. The creep rupture strength increased with increasing W concentration or by the addition of V and Ta which formed fine MC carbides. In the accelerated creep region, the logarithm of creep rate increased linearly with strain but not linearly with time. The acceleration of creep rate by strain correlates with the microstructural stability and increases with increasing W concentration or by the addition of V and Ta. 0 1997 Elsevier Science S.A. Keywords:
9Cr
steel; Tempered
martensite;
Creep
rate;
Creep
rupture
1. Introduction
2. Experimental procedure
Tempered martensitic 9-12Cr steels strengthened by W have become of much interest as advanced ferritic/ martensitic steelsin application to boiler and turbine of ultra-supercritical electric power plants [11, fuel cladding of fast breeder reactor [2] and first wall and blanket of fusion reactor [3]. These steels are alloyed with many elements, C, Cr, W, MO, V, Nb, N, B, etc., for improving
strength
solid solution
and precipitation
strength-
ening but the main alloying constituents are C-Cr-W. In tempered martensitic steels, it is well known that the secondary or steady-state creep region is substantially not present and that a greater part of the creep life is dominated by the tertiary or acceleration creep region. The purpose of the present research is to investigate the microstructural evolution and its effect on creep rates in the acceleration creep region for simple 9Cr-W and solute modified 9Cr-WVTa steels.
Four simple 9Cr-(0,1,2,4)W-O.lC lute modified
9Cr-1WVTa
steels and two so-
and 9Cr-3WVTa
steels were
used. The chemical compositions of the steelsare given in Table 1. The steel rods were austenitized, quenched and then tempered [4-71. The steels, except for the 9Cr-4W steel, were 100% tempered martensite, while the 9Cr-4W steel contained 10 vol.% d-ferrite. The tempered
martensite
consisted
of lath
subgrains
that
contained a high density of dislocations and fine carbides of about 0.1 pm or less. The major component of carbides in as tempered condition was M,& in each steel. Creep tests were carried out at 823, 873 and 923 K up to about 15000 h, using specimens of 6 mm in gage diameter and 30 mm in gage length.
3. Results and discussion 3.1. Creep behavior and microstructural evolution during creep
+ 81 * Tel.: [email protected] 0921-5093/97/$17.00 PZI SO921
298
592215;
fax:
0 1997 Elsevier
-5093(97)00404-8
+81
Science
298
592201;
S.A. All rights
e-mail:
reserved.
Fig. 1 shows the creep rupture strength of the steels as a function of Larson-Miller Parameter (LMP) with a constant C= 20. The decrease in creep rupture strength with LMP becomes less pronounced with in-
1046
F. Abe/Materials
Table 1 Chemical
compositions
9Cr 9Cr-1W 9Cr-2W 9Cr-4W 9Cr-1 WVTa 9Cr-3WVTa
of steels examined
Science
and Engineering
Cr
W
0.104 0.101 0.100 0.101 0.125 0.167
8.96 9.01 8.92 9.09 9.24 9.16
0.99 1.92 3.93 1.06 3.08
V
Ta
0.18 0.16
3oo-
AA
0
0
-
-
20-
4tO. A q
k,,++
l
0
A 0 v A l
l
A 0 av 0 M ‘b4 03 A co A A.o 00
60-
i% 40-
I
90
A
0
Fig. 1. Creep rupture the 9Cr-W steels.
0
N
0.49 0.48 0.48 0.50
0.30 0.29 0.28 0.29
0.009 0.011 0.012 0.006 0.003 0.003
0.001 0.002 0.002 0.002 0.003 0.003
3.2. Creep
rates in accelerated
creep region
Fig. 2 shows the creep rate of the 9Cr-1W steel as a function of true strain. In the acceleration creep region
1 E-2
l+ g
qn
lE-3
3
VA v
t
1 E-4
q
A
-
0
-
n
1 E-5
I
1
,
[ 20
stress vs. Larson
+ log Miller
fr(h)
873K, 137MPa
873K, 118MPa 0 873K, 98MPa l 873K, 78MPa A 873K, 69MPa A 923K. 29MPa
6
1 E-6
20000 22000 18000 Larson-Miller parameter T(K)
Si
1 E-l
0
SCr-1W 9Cr-2W 9Cr-4W 9Cr-1WVTa SCr-3WVTa I 1
Mn
tion of Fe,W also occurred during creep in the 9Cr-4W and 9Cr-3WVTa steels [7]. The microstructural evolution became less pronounced with increasing W concentration or by the addition of V and Ta and was much larger during creep than during aging under unstressedconditions. This suggeststhat the stressand strain effects are important. The increased creep rupture strength of the 9Cr-W steels correlated with the increased stability of the tempered martensite microstructure. V and Ta formed fine MC carbides having a size of several 10 nm in the matrix in as tempered condition. The fine MC carbides were effective as obstacles for the recovery of dislocations and the migrations of lath boundaries.
823 - 923 K 2
1045-1048
0.005 0.007
873 K 104h 10’h 4 4
Q
E E ‘g)L
B
0.10 0.10
creasing W concentration or by the addition of V and Ta. The addition of small amounts of V and Ta significantly increases the long-term creep rupture strength. The creep rupture strength at LMP = 20 952, corresponding to 10’ h at 873 K, is evaluated to be 81 and 145 MPa for the 9Cr-1WVTa and 9Cr-3WVTa steels, respectively. The creep curves of the present steels consisted of the primary or transient creep region, where the creep rate decreaseswith time and the tertiary or accelerated creep region, where the creep rate increaseswith time. The stressdependence of the minimum creep rate was described by a power law or Norton’s law as &,, = A#‘, where A is a constant and 12the stress exponent. The stress exponent y1increased with increasing W concentration or by the addition of V and Ta. During creep, the recovery of excessdislocations, the agglomeration of carbides and the growth of martensite lath subgrains occurred [446]. The extensive precipita-
1
(1997)
(wt%)
C
I
A234-236
I
0
10
I
20
I
30
40
True strain (%)
1
Parameter
(LMP)
for
Fig. 2. Creep rate-true strain curves for the 9Cr-1W steel. The time to rupture was 42, 130, 482, 1751, 4331 h at 137, 118, 98, 78 and 69 MPa, respectively, at 873 K and was 12 560 h at 29 MPa and 923 K.
F. Abe /Materials
1oo-----l s .
_ 0 90 A SCr-1W 80 q 9Cr-2W _ v 9Cr-4W A 9Cr-1WVTa 60- + 9Cr-3Ma
Science
+
and Engineering
.J
+
+ 0
*cd
v
v
V 0
A A
I
I
I
’
18000
Larson-Miller T(K) Fig. 3. Relationship
between
1
I 20000 [ 2.
1 22000
I
parameter + log
d In B/d& and LMP
fr(h)
1
for the 9Cr-W
steels.
after reaching a minimum creep rate, the logarithm of creep rate increases linearly with true strain over a wide range of strain. The linear acceleration of creep rate with strain has been reported for several Cr-Mo steels [&lo] and the present 9Cr-2W steel [5]. It should be noted that the acceleration of creep rate occurs in two stages; stage I just after reaching a minimum creep rate of about 10% true strain and stage II above about 10% true strain. The d In i/d& is much larger in stage I than in stage II. Stage I becomes more pronounced with decreasing stress level which suggests that stage I is important for the analysis of long-term creep behavior. At low stresses, the creep rate rapidly decreased just before reaching a minimum creep rate and the rapidly decreased creep rate recovered in stage I. Although the main microstructural evolution in the transient creep region was the recovery of excess dislocations to lath
boundaries,the rapid decreasein creepratejust before reaching a minimum creep rate results from the formation of creep-resistant microstructure where the distributions of excess dislocations, fine precipitates and lath boundaries are highly resistant to further recovery of excess dislocations. Thus the appearance of stage I correlates with the recovery or evolution of previously formed creep-resistant microstructure. Fig. 3 shows the d In i/d& in stage I for the steels as a function of LMP. The 9Cr steel containing no W exhibits the smallest d In i/de, 12-15 for LMP < 21000, of the steels examined. The d In S1d.zincreases with increasing W concentration or by the addition of V and Ta, although the variation in d In c/d& among the steels is not large at LMP < 21000. At LMP > 21000, the d In c/d&
A234-236
(1997)
1045
IO48
1047
abruptly increases with LMP for the steels containing high W concentrations or containing V and Ta. Under constant load conditions, as in the present creep test, an increase in stress by a decrease in cross section with strain causes the acceleration of creep rate and contributes to the d In i/d&. For uniaxial extension at constant volume, we have the well-known relation given by 0 = (1 + E)g,, where g is the stress at the engineering strain E and crO the initial stress. Because E = ln(1 + E), where E is the true strain, we obtain d In i/d& = (d In i/d In a)(d In a/d&) = n, where n is the stress exponent of the Norton’s law. The iz for the minimum creep rate was evaluated to be 5, 7, 13, 15, 12 and 25 for the 9Cr, 9Cr-lW, 9Cr-2W, 9Cr-4W, 9Cr1WVTa and 9Cr-3WVTa steels, respectively, at 873 K. Therefore, the d In S/d& is much larger than the 12in each steel. This suggeststhat the contribution of stress increase by a decreasein cross section to the d In iId& is rather small and that a greater part of the d In k/d& results from the effect of microstructural degradation by creep strain. The contribution of microstructural degradation to the d In i/dE, corresponding to (d In i/ dc - n), increases with increasing W concentration or by the addition of V and Ta. This is the same trend as the creep rupture strength and the microstrnctural stability during creep described in the previous section.
4. Conclusion The creep rupture strength of the 9Cr-W steels increases with increasing W concentration or by the addition of V and Ta; the lo5 h-rupture strength at 873 K was evaluated to be 81 and 145 MPa for the 9Cr1WVTa and 9Cr-3WVTa steels, respectively. In the accelerated creep region, the logarithm of creep rate increaseslinearly with creep strain. The acceleration of creep rate by strain occurred in two stages: stage I at low strains just after reaching a minimum creep rate and stage II at high strains. Stage I became more pronounced with decreasing stresslevel and the d In ,6/ d& was much larger in stage I than in stage II. Stage I is important for the analysis of long-term creep behavior. The d In i/d& increased with increasing W concentration or by the addition of V and Ta.
References [l]
E. Metcalfe, W.T. Bakker, Proceedings of the EPRI/National Power Conference, National Power, 1995, p.1. [2] S. Nomura, S. Shikakura, S. Ukai, I. Seshimo, M. Harada, I. Shibahara, M. Katsuragawa, Proceedings of the International Conference on Fast Reactors and Related Fuel Cycles, Kyoto, Japan, 1991, p. 7.4-l. [3] R.L. Klueh, K. Ehrlich, F. Abe, J. NucI. Mater. 191-194 (1992) 116.
1048 [4] F. Abe, S. (1992) 469. [5] F. Abe, S. [6] F. Abe, S. [7] F. Abe, T.
F. Abe/Materials Nakazawa,
H.
Araki,
T. Noda,
Science Met.
Trans.
Nakazawa, Met. Trans. 23A (1992) 3025. Nakazawa, Mater. Sci. Tech. 8 (1992) 1063. Noda, M. Okada, J. Nucl. Mater. 195 (1992)
and Engineering 23A
51.
A234-236
(1997)
10451048
[8] M. Prager, Pressure Vessel and Piping 288 (1994) 401. [9] R. Woo, R. Sandstrom, J. Storesund, Mater. Temp. 12 (1994) 271. [lo] S. Straub, M. Meier, J. Ostermann, W. Blum, VGB Kraftwerkstechnik 73 (1993) 646.
Materials
Evolution
Science
and Engineering
A234-236
(1997)
1045-1048
of microstructure and acceleration of creep rate in tempered martensitic 9Cr-W steels Fujio Abe * National
Research
Received
Institute
27 February
for
Metals,
1997; received
1-2-1
Sengen,
in revised
Tsukuba
form
305, Japan
28 March
1997
Abstract The effect of microstructural evolution on creep rates in the acceleration creep region has been investigated for tempered martensitic 9Cr-W steels. Creep tests were carried out at 823, 873 and 923 K for up to 15000 h. The creep rupture strength increased with increasing W concentration or by the addition of V and Ta which formed fine MC carbides. In the accelerated creep region, the logarithm of creep rate increased linearly with strain but not linearly with time. The acceleration of creep rate by strain correlates with the microstructural stability and increases with increasing W concentration or by the addition of V and Ta. 0 1997 Elsevier Science S.A. Keywords:
9Cr
steel; Tempered
martensite;
Creep
rate;
Creep
rupture
1. Introduction
2. Experimental procedure
Tempered martensitic 9-12Cr steels strengthened by W have become of much interest as advanced ferritic/ martensitic steelsin application to boiler and turbine of ultra-supercritical electric power plants [11, fuel cladding of fast breeder reactor [2] and first wall and blanket of fusion reactor [3]. These steels are alloyed with many elements, C, Cr, W, MO, V, Nb, N, B, etc., for improving
strength
solid solution
and precipitation
strength-
ening but the main alloying constituents are C-Cr-W. In tempered martensitic steels, it is well known that the secondary or steady-state creep region is substantially not present and that a greater part of the creep life is dominated by the tertiary or acceleration creep region. The purpose of the present research is to investigate the microstructural evolution and its effect on creep rates in the acceleration creep region for simple 9Cr-W and solute modified 9Cr-WVTa steels.
Four simple 9Cr-(0,1,2,4)W-O.lC lute modified
9Cr-1WVTa
steels and two so-
and 9Cr-3WVTa
steels were
used. The chemical compositions of the steelsare given in Table 1. The steel rods were austenitized, quenched and then tempered [4-71. The steels, except for the 9Cr-4W steel, were 100% tempered martensite, while the 9Cr-4W steel contained 10 vol.% d-ferrite. The tempered
martensite
consisted
of lath
subgrains
that
contained a high density of dislocations and fine carbides of about 0.1 pm or less. The major component of carbides in as tempered condition was M,& in each steel. Creep tests were carried out at 823, 873 and 923 K up to about 15000 h, using specimens of 6 mm in gage diameter and 30 mm in gage length.
3. Results and discussion 3.1. Creep behavior and microstructural evolution during creep
+ 81 * Tel.: [email protected] 0921-5093/97/$17.00 PZI SO921
298
592215;
fax:
0 1997 Elsevier
-5093(97)00404-8
+81
Science
298
592201;
S.A. All rights
e-mail:
reserved.
Fig. 1 shows the creep rupture strength of the steels as a function of Larson-Miller Parameter (LMP) with a constant C= 20. The decrease in creep rupture strength with LMP becomes less pronounced with in-
1046
F. Abe/Materials
Table 1 Chemical
compositions
9Cr 9Cr-1W 9Cr-2W 9Cr-4W 9Cr-1 WVTa 9Cr-3WVTa
of steels examined
Science
and Engineering
Cr
W
0.104 0.101 0.100 0.101 0.125 0.167
8.96 9.01 8.92 9.09 9.24 9.16
0.99 1.92 3.93 1.06 3.08
V
Ta
0.18 0.16
3oo-
AA
0
0
-
-
20-
4tO. A q
k,,++
l
0
A 0 v A l
l
A 0 av 0 M ‘b4 03 A co A A.o 00
60-
i% 40-
I
90
A
0
Fig. 1. Creep rupture the 9Cr-W steels.
0
N
0.49 0.48 0.48 0.50
0.30 0.29 0.28 0.29
0.009 0.011 0.012 0.006 0.003 0.003
0.001 0.002 0.002 0.002 0.003 0.003
3.2. Creep
rates in accelerated
creep region
Fig. 2 shows the creep rate of the 9Cr-1W steel as a function of true strain. In the acceleration creep region
1 E-2
l+ g
qn
lE-3
3
VA v
t
1 E-4
q
A
-
0
-
n
1 E-5
I
1
,
[ 20
stress vs. Larson
+ log Miller
fr(h)
873K, 137MPa
873K, 118MPa 0 873K, 98MPa l 873K, 78MPa A 873K, 69MPa A 923K. 29MPa
6
1 E-6
20000 22000 18000 Larson-Miller parameter T(K)
Si
1 E-l
0
SCr-1W 9Cr-2W 9Cr-4W 9Cr-1WVTa SCr-3WVTa I 1
Mn
tion of Fe,W also occurred during creep in the 9Cr-4W and 9Cr-3WVTa steels [7]. The microstructural evolution became less pronounced with increasing W concentration or by the addition of V and Ta and was much larger during creep than during aging under unstressedconditions. This suggeststhat the stressand strain effects are important. The increased creep rupture strength of the 9Cr-W steels correlated with the increased stability of the tempered martensite microstructure. V and Ta formed fine MC carbides having a size of several 10 nm in the matrix in as tempered condition. The fine MC carbides were effective as obstacles for the recovery of dislocations and the migrations of lath boundaries.
823 - 923 K 2
1045-1048
0.005 0.007
873 K 104h 10’h 4 4
Q
E E ‘g)L
B
0.10 0.10
creasing W concentration or by the addition of V and Ta. The addition of small amounts of V and Ta significantly increases the long-term creep rupture strength. The creep rupture strength at LMP = 20 952, corresponding to 10’ h at 873 K, is evaluated to be 81 and 145 MPa for the 9Cr-1WVTa and 9Cr-3WVTa steels, respectively. The creep curves of the present steels consisted of the primary or transient creep region, where the creep rate decreaseswith time and the tertiary or accelerated creep region, where the creep rate increaseswith time. The stressdependence of the minimum creep rate was described by a power law or Norton’s law as &,, = A#‘, where A is a constant and 12the stress exponent. The stress exponent y1increased with increasing W concentration or by the addition of V and Ta. During creep, the recovery of excessdislocations, the agglomeration of carbides and the growth of martensite lath subgrains occurred [446]. The extensive precipita-
1
(1997)
(wt%)
C
I
A234-236
I
0
10
I
20
I
30
40
True strain (%)
1
Parameter
(LMP)
for
Fig. 2. Creep rate-true strain curves for the 9Cr-1W steel. The time to rupture was 42, 130, 482, 1751, 4331 h at 137, 118, 98, 78 and 69 MPa, respectively, at 873 K and was 12 560 h at 29 MPa and 923 K.
F. Abe /Materials
1oo-----l s .
_ 0 90 A SCr-1W 80 q 9Cr-2W _ v 9Cr-4W A 9Cr-1WVTa 60- + 9Cr-3Ma
Science
+
and Engineering
.J
+
+ 0
*cd
v
v
V 0
A A
I
I
I
’
18000
Larson-Miller T(K) Fig. 3. Relationship
between
1
I 20000 [ 2.
1 22000
I
parameter + log
d In B/d& and LMP
fr(h)
1
for the 9Cr-W
steels.
after reaching a minimum creep rate, the logarithm of creep rate increases linearly with true strain over a wide range of strain. The linear acceleration of creep rate with strain has been reported for several Cr-Mo steels [&lo] and the present 9Cr-2W steel [5]. It should be noted that the acceleration of creep rate occurs in two stages; stage I just after reaching a minimum creep rate of about 10% true strain and stage II above about 10% true strain. The d In i/d& is much larger in stage I than in stage II. Stage I becomes more pronounced with decreasing stress level which suggests that stage I is important for the analysis of long-term creep behavior. At low stresses, the creep rate rapidly decreased just before reaching a minimum creep rate and the rapidly decreased creep rate recovered in stage I. Although the main microstructural evolution in the transient creep region was the recovery of excess dislocations to lath
boundaries,the rapid decreasein creepratejust before reaching a minimum creep rate results from the formation of creep-resistant microstructure where the distributions of excess dislocations, fine precipitates and lath boundaries are highly resistant to further recovery of excess dislocations. Thus the appearance of stage I correlates with the recovery or evolution of previously formed creep-resistant microstructure. Fig. 3 shows the d In i/d& in stage I for the steels as a function of LMP. The 9Cr steel containing no W exhibits the smallest d In i/de, 12-15 for LMP < 21000, of the steels examined. The d In S1d.zincreases with increasing W concentration or by the addition of V and Ta, although the variation in d In c/d& among the steels is not large at LMP < 21000. At LMP > 21000, the d In c/d&
A234-236
(1997)
1045
IO48
1047
abruptly increases with LMP for the steels containing high W concentrations or containing V and Ta. Under constant load conditions, as in the present creep test, an increase in stress by a decrease in cross section with strain causes the acceleration of creep rate and contributes to the d In i/d&. For uniaxial extension at constant volume, we have the well-known relation given by 0 = (1 + E)g,, where g is the stress at the engineering strain E and crO the initial stress. Because E = ln(1 + E), where E is the true strain, we obtain d In i/d& = (d In i/d In a)(d In a/d&) = n, where n is the stress exponent of the Norton’s law. The iz for the minimum creep rate was evaluated to be 5, 7, 13, 15, 12 and 25 for the 9Cr, 9Cr-lW, 9Cr-2W, 9Cr-4W, 9Cr1WVTa and 9Cr-3WVTa steels, respectively, at 873 K. Therefore, the d In S/d& is much larger than the 12in each steel. This suggeststhat the contribution of stress increase by a decreasein cross section to the d In iId& is rather small and that a greater part of the d In k/d& results from the effect of microstructural degradation by creep strain. The contribution of microstructural degradation to the d In i/dE, corresponding to (d In i/ dc - n), increases with increasing W concentration or by the addition of V and Ta. This is the same trend as the creep rupture strength and the microstrnctural stability during creep described in the previous section.
4. Conclusion The creep rupture strength of the 9Cr-W steels increases with increasing W concentration or by the addition of V and Ta; the lo5 h-rupture strength at 873 K was evaluated to be 81 and 145 MPa for the 9Cr1WVTa and 9Cr-3WVTa steels, respectively. In the accelerated creep region, the logarithm of creep rate increaseslinearly with creep strain. The acceleration of creep rate by strain occurred in two stages: stage I at low strains just after reaching a minimum creep rate and stage II at high strains. Stage I became more pronounced with decreasing stresslevel and the d In ,6/ d& was much larger in stage I than in stage II. Stage I is important for the analysis of long-term creep behavior. The d In i/d& increased with increasing W concentration or by the addition of V and Ta.
References [l]
E. Metcalfe, W.T. Bakker, Proceedings of the EPRI/National Power Conference, National Power, 1995, p.1. [2] S. Nomura, S. Shikakura, S. Ukai, I. Seshimo, M. Harada, I. Shibahara, M. Katsuragawa, Proceedings of the International Conference on Fast Reactors and Related Fuel Cycles, Kyoto, Japan, 1991, p. 7.4-l. [3] R.L. Klueh, K. Ehrlich, F. Abe, J. NucI. Mater. 191-194 (1992) 116.
1048 [4] F. Abe, S. (1992) 469. [5] F. Abe, S. [6] F. Abe, S. [7] F. Abe, T.
F. Abe/Materials Nakazawa,
H.
Araki,
T. Noda,
Science Met.
Trans.
Nakazawa, Met. Trans. 23A (1992) 3025. Nakazawa, Mater. Sci. Tech. 8 (1992) 1063. Noda, M. Okada, J. Nucl. Mater. 195 (1992)
and Engineering 23A
51.
A234-236
(1997)
10451048
[8] M. Prager, Pressure Vessel and Piping 288 (1994) 401. [9] R. Woo, R. Sandstrom, J. Storesund, Mater. Temp. 12 (1994) 271. [lo] S. Straub, M. Meier, J. Ostermann, W. Blum, VGB Kraftwerkstechnik 73 (1993) 646.