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Microstructure and mechanical properties of unirradiated low activation Ferritic steel
Journal of Nuclear Materials 141-143 (1986) 1107-1112 North-Holland, Amsterdam
MICROSTRUCTURE
AND
1107
MECHANICAL
PROPERTIES
OF UNIRRADIATED
LOW
ACTIVATION FERRITIC STEEL Chen-Yih
HSU and Thomas
A. LECHTENBERG
GA Technologies Inc., P.O. Box 85608, San Diego, CA 92138, USA
Transmission electron micrographs of normalized and tempered 9Cr-2.5W~.3V~.lYZ low activation ferritic steel showed tempered lath-type martensite with precipitation of rod and plate-like carbides at lath and grain boundaries. X-ray diffraction analysis of the extracted replicas revealed nearly 100% M& carbides (a = 1.064 nm), with no indication of Fe,W-type Laves phase even after thermal aging at 600”C/1000 h. Thermal aging increased the number density of rod-like Mz3Chalong prior austenite grain boundaries and martensite lath boundaries. The elevated-temperature tensile strengths of this steel are about 10% higher than the average strengths of commercial heats of 9Cr-1Mo and modified 9Cr-1Mo steels up to 650°C with equivalent uniform elongation and -50% decrease in total elongation. The DBTT was determined to be -25°C which is similar to other 9Cr-1Mo steels. Fractographic examination of tensile tested specimens shows a mixed mode of equiaxed and elongaied dimples at test temperatures above 400°C. Modification of the GA3X alloy composition for optimization of materials properties is discussed. However, the proposed low activation ferritic steel shows the promise of improved mechanical properties over 9Cr-1Mo steels. 1. Introduction available alloys to be used for first structural components will become during service and will pose a waste disposal
Commercially
wall and blanket activated
problem after service. Guidelines (10CFR61) for the classification of nuclear wastes have been issued by the US Nuclear Regulatory Commission. In 1982, DOE convened a panel to examine the issues and incentives associated with the development of low activation materials. The available data show that the 9-12%Cr ferriticl martensitic steels are attractive materials for fusion reactor structural component applications. The primary objective of this work is to develop ferritic steels that exhibit equivalent mechanical properties and performance to the commercially developed alloys, HT-9 (12Cr-1MoVW) and/or 9Cr-1MoVNb steels, but which may allow near-surface disposal after fusion reactor service. These steels are hereafter referred to as low activation. To achieve this, the elements MO, Ni, Nb, and N, which are added for elevated temperature strength, must be reduced to very low levels. W and V can be used to substitute for MO. In equal atomic fractions, W and MO show similar solid solution hardening characteristics and Fe-W-C alloys develop analogous precipitates with a similar precipitation sequence to Fe-MO-C alloys. Therefore, the Cr-W steels offer promise for the development of a replacement for Cr-Mo steels. An experimental heat (GA3X) of low activation ferritic steel was prepared with a nominal composition of 9Cr-2.5W4.3V4.15C steel. The addition of 0.3% V to Cr-W steels will result in a pronounced effect on the precipitate formation and enhance the elevated temperature properties, but may also make the steel less weldable. For this reason, the suggested carbon content is kept below 0.15% in Cr-W steels. Titanium is a strong carbide former with an evi0022-3115/86/$03.50 0 Elsevier Science Publishers (North-Holland Physics Publishing Division)
dence of inhibiting grain coarsening while retaining good notch toughness, so the addition of 0.1% Ti is also suggested as a modification to GA3X although work on this not reported here. This paper summarizes the metallographic examination and mechanical property evaluation of the unirradiated GA3X low activation ferritic steel compared to commercial heats of 9Cr1Mo steels, and also provides baseline data for a comparison with the irradiated GA3X steel. 2. Experimental
procedures
The melting and fabrication practice of the GA3X low activation ferritic steel are described in ref. [l]. The chemical composition of this GA3X heat is shown in table 1. The specimens discussed here were water quenched from lOOO”C/lh and then tempered at 7OO”C/l h. A two-stage replication technique was used to prepare the carbon film extraction replicas of the precipitates [2]. For X-ray diffractometry analysis, 5 mg of the precipitate were smeared onto a single crystal quartz slide in a graphite crystal monochromator-scintillation detector system. In the preparation of thin-foil specimens, 3 mm disks were electropolished in a solution of 50% ethyl alcohol, 45% butyl cellosolve, and 5% HC104. Carbon film extraction replicas and thin-foil samples were examined on-a Philips 300M transmission electron microscope operating at 100 kV.
Table 1 Chemical composition of GA3X ferritic steel (wt%) C
Cr
W
V
Mn
S
0.16-0.17
7.4-7.6
1.9-2.0
0.014-0.016
0.002
0.005
B.V.
1108
C.-Y. Hsu, T. A. Lechtenherg
I Unirradiated
The tensile tests were performed on flat tensile specimens with a reduced gage section 25.4 mm long by 3.18 mm wide by 1.0 mm thick. The flat tensile specimens were dipped in toluene pretreat coating solution in order to avoid any significant oxidation during the tensile test at temperature in the range of 2@-65O”C. Fractographic examination of tensile tested specimens was performed on a scanning electron microscope. The ductile brittle transition temperature (DBTT) and impact properties were determined from f size miniature Charpy V-notch specimens (23.6 mm x 3.33 mm x 3.33 mm with a 30” notch angle and 0.5 mm notch depth). The miniature Charpy V-notch specimens were fatigue precracked to approximately a/w = 0.5, and were then tested in the instrumented impact testing machine at the Hanford Engineering Development Laboratory. 3. Results and discussion 3.1. Metallographic examination The weight percentage of the extracted precipitates was increased from 2.0% of the quenched and tempered specimens to 3.1% after thermal aging at 600°C for 100 h and then remained unchanged up to 1000 h. Further, there was no indication of Fe,W-type intermetallic compound in any of the samples analyzed [3]. X-ray diffraction analysis of the extracted precipitates revealed nearly 100% M,,C, carbides with a lattice parameter of 1.064 nm. The X-ray fluorence analysis indicated that the extracted M&carbides were in the form of (Feo,21Wo,osCro,,o)23C6. There was no change of metal contents in M& carbides after thermal aging treatment. Fig. 1 shows TEM micrographs of extraction replicas prior to and after thermal aging at 600°C for 100 h. The extracted precipitates show two types of carbides, with rod-like and plate-like shapes, similar to results on HT-9 steel reported by Lechtenberg et al. [4]. The selected area diffraction patterns from the isolated particles confirmed the carbides to be M&, and a few particles seem likely to be MC-types carbide. This is consistent with the X-ray diffraction analyses. The scarcity of MC-type carbide is due to very low V content (0.015%) and no Ti added in this steel. An appreciable amount of MC-type carbide should be detected if the target value of 0.3% V or Ti were added to this steel. The size of precipitates was typically in the range of 20-750 nm. The mean particle size increased from 100 nm to 140 nm after thermal aging at 600°C for 100 h and many more rod-like carbides were found in the thermally aged specimens than in the astempered specimens. Thin-foil TEM micrographs show tempered lath martensite within the prior-austenite grains, as shown in fig. 2. No significant amount of ferrite was seen. In
low activation ferritic strrl
both the tempered and the thermally aged specimens, heavy carbide precipitation was noted at the prioraustenite grain boundaries and martensite lath boundaries. The areal density of carbide precipitation is appreciably higher in the thermally aged specimen than in the tempered specimen as judged directly from the thin-foil microstructural observation. Particle coalcscence and an increase in the areal density of rod-like carbides were found on grain and lath boundaries after the thermal aging treatment. This observation can be attributed to enhanced diffusion at grain and lath boundaries. The evidence shows that the number density of carbides and carbide size increased, and particle morphology tended to change during thermal aging treatment. The precipitation process and morphology change of particles and the increase of elongated rodlike carbides along grain boundaries and lath boundaries may influence the impact and fracture toughness of this steel, as well as other mechanical behavior such as tensile and creep strengths. In the quenched and tempered condition, dislocations are ordered in such a way as to produce irregular subgrains within the martensite laths. After thermal aging at 600°C for 100 h there were no apparent
Olwn
Fig. 1. TEM micrographs of carbon film extraction replica of GA3X steel, (a) quenched and tempered condition, and (b) followed by thermal aging at 6OO”C/lOOh.
C.-Y. Hsu, T. A. Lechtenberg
1109
I Unirradiated low activation ferritic steel
higher than the average value for commercial heats of 9Cr-1Mo and modified 9Cr-1Mo steels at temperatures up to 400°C and are slightly higher at temperatures above 500°C. This may be partly due to the low tempering temperature (700°C) given to GA3X whereas 9Cr-1Mo and 9Cr-1MoVNb are tempered in the range of 740-760°C. A significant contributing factor may be the alloy content. The actual Cr content of this steel is only 7.5% (target 9%) and the concentration of W is only 2.0 wt% (target 2.5 wt%). The mod-
600
a
8
AVERAGE CURVE OF MOOlFlEO SCr-1Mo
. NOMINAL
100 0' 0
Fig. 2. TEM thin-foil micrographs of GA3X steel, (a) quenched and tempered condition, and (b) followed by thermal aging at 6OO”C/lOOh. changes in the dislocation sub-structures. Changes in dislocation structure may become more pronounced after prolonged thermal aging at service temperatures or during irradiation. Resistance to temper embrittlement is controlled by maintaining low levels of impurity elements such as P, S, Sb, Sn, Mn, and Si. Furthermore, the formation of intermetallic compounds such as Laves phases may promote the segregation of impurity elements and thus degrade the resistance to temper embrittlement. In the GA3X steel, the Mn and S contents are less than 0.005% while the other impurity elements are negligible. Based on the analysis proposed by Watanabe et al.[5], the effect of these impurities combined with a fine grain size lead us to conclude little temper embrittlement will occur.
COMPOSITION:
A SCr-1MoV Nb (REF.51
I
I
1
I
1
I
100
200
300
400
500
600
700
TEST TEMPERATURE,°C 900
700
g
600
AVERAGECURVEOF MODIFIED SCr-1Mo
z $ 500 I k 9 400 Ly > k t
300
:: 2 200
NOMINAL COMPOSITION: 0 9Cr-2.5W-0.3V-0.l5C
109
A SCr-1MoV Nb IfiEFs 51 0
3.2. Mechanical
properties
0
100
200
300
400
500
600
700
TEST TEMPERATURE,°C
The ultimate tensile strength and 0.2% offset yield strength of this ferritic steel are compared with the average values of 9Cr-1Mo and modified 9Cr-1Mo steels [6,7] in fig. 3. Tensile strengths are about 10%
Fig. 3. The tensile strengths of 9Cr-2.5W4I.3V41.5C (GA3X) steel compared to average curves of 9Cr-1Mo and modified 9Cr-1Mo steels (band), (a) ultimate tensile strength, and (b) 0.2% offset yield strength.
1110
C.-Y. Hsu. T.A. Lechtenberg I Unirradiated low activation ferritic steel
ified 9Cr-1Mo steels contains 0.21 wt% V and 0.01 wt% Ti and GA3X steel contains only 0.015 wt% V. Hence, the enhanced tensile strength observed in GA3X steel is likely to be due to the effect of a 0.17% carbon content, which is higher than carbon content (0.1%) of both 9Cr-1Mo steels. In comparison with the experimental heat of unirradiated 9Cr-IMoVNb steel [8], the GA3X steel shows higher strengths and ductility from 20°C to 4OO”C, and becomes slightly weaker at temperatures above 500°C but still maintains a higher ductility. In contrast to 0.015% V in GA3X steel, the 9Cr-1MoVNb steel contains 0.24% V and 0.08% Nb which contributed significantly to the elevated temperature strengths. So, the strength of GA3X steel may be significantly improved if the alloying elements added for elevated-temperature strength had been closer to target contents. It is indicated that the low activation ferritic steel with a nominal composition of 9Cr-2.5W-O.3Va.15C may show a promising potential of improved strength at elevated temperatures, compared to the 9Cr-1Mo steels. The uniform and total elongations of GA3X steel and the average data of commercial heats of 9Cr-1Mo and modified 9Cr-1Mo steels are shown in fig. 4. The uniform elongation of this steel is about equivalent to that of modified 9Cr-1Mo steel while the total elongation is -50% less than both 9Cr-1Mo steels. The workhardening exponent is approximately equivalent to the uniform elongation. This indicates that the workhardening response of this steel is similar to that of 9Cr-1Mo steels which indicate similar precipitation effects on dislocation movement and multiplication. Changes in fracture toughness (K,,) can be estimated from tensile behavior and the associated microstructural evolution of the steel [9-141. An excellent agreement between tensile-base prediction and KIC fracture toughness measurements was obtained on irradiated stainless steel by using the modified Krafft correlation [13]. In the case of low work-hardening coefficients, the modified Krafft correlation can be simplified as: Kit-
(oYer)lh,
In the fractographic examination of the tensile tested specimens [15], a ductile dimple rupture caused by a void coalescence mechanism was observed for specimens tested from 20°C up to 400°C. At temperatures above 500°C the fracture mode was still ductile as shown by the dimples, but more than 30% of the dimples were elongated, which was usually indicated to be the result of ductile tearing [16]. This mixed fracture mode suggests that there might have been a strong shear component and/or cross slip involved in the frac-
AVERAGECURVEOF MODIFIED SCr-1Yo
0
100
200
300
400
500
700
600
6OC
TEST TEMPERATURE,PC
t AVERAGECURVE OF MODIFIED SCr-1Mo
AVERAGE CURVE OF SCr-1Mo STEEL
(I)
where a,, is the yield strength and .sris the true fracture strain. The opposing changes in yield strength and ductility tend to counterbalance each other. The workhardening coefficients of GA3X ferritic steel are 0.1 at lower temperatures and become 0.06 at higher temperatures. In the same class of 9Cr ferritic steels, the variation of K,, fracture toughness can be estimated by comparing the differences of tensile properties. Based on eq. (l), the plane-strain KIC of GA3X steel is predicted to be about lO-20% lower than the commercial heats of modified 9Cr-1Mo steel [6], and is expected to be slightly higher than results on an experimental heat of 9Cr-1MoVNb steel at the same temperatures [71.
0
100
200
300
400
500
600
700
600
TEST TEMPERATURE,%
Fig. 4. Elongation (a) uniform
data vs test temperature on GA3X elongation, and (b) total elongation.
steel,
C.-Y. Hsu, T.A. Lechtenberg
I Unirradiated low activation ferritic steel
ture at the high temperatures (?0.47’,,,). Braski and Maziasz [17] studied the elevated-temperature tensile properties of PCA austenitic stainless steel and found shear-elongated dimples at test temperatures from 300°C to 600°C in the 25% cold-worked condition. Once the cold-worked material was heat treated at 75O”C/2 h, MC-type carbides precipitated and the mixed-mode vanished. It is suggested that the shear elongated dimples may be eliminated with more MC carbide precipitation as would be the case by increasing V and/or Ti contents of GA3X heat. However, the reason for the formation of shear dimples is not yet fully understood and deserves further investigation. The impact test results on the precrackedi size Charpy specimens of this ferritic steel are tabulated in table 2. The ductile-brittle transition temperature was determined to be -25°C with an upper shelf energy of 2.2 J (37.2 J/cm2 normalized fracture energy). The DBTT and impact energy of this low activation ferritic steel is very similar to the modified 9Cr-1Mo steels ]I81. The experimental results of GA3X steel provide some clues concerning optimum alloy composition. The mechanical properties of steel can be optimized by controlling microstructure through processing and by modifying alloy composition. For example, significant strengthening was contributed by the 0.17% carbon content. High carbon content can make the steel less weldable, and increase particle coarsening may degrade ductility and fracture toughness. The carbon content should be reduced to 0.15% or lower. The decrease in strength due to the lowering of carbon content can be compensated by increasing Cr, V, and Ti alloying addition. The Cr content is to be maintained at a level of 9% to provide oxidation resistance at elevated temperatures. The phase stability is a key concern for the justification of alloy composition. V and Ti are strong carbide formers. The addition of 0.3% V and
Fracture
Maximum
Normalized
temperature
energy (J)
load (kN)
energy
100°C 50°C 21°C 0°C -10°C -25°C -50°C -80°C
2.225 2.499 2.264 2.180 1.658 1.113 0.498 0.050
0.645 0.733 0.693 0.726 0.648 0.574 0.571 0.536
37.56 38.29 36.11 34.10 27.16 17.73 8.01 0.85
(J/cm2)
1111
intermetallic phase precipitation, such as Fe2W type Laves phase, after a prolonged thermal treatment or during irradiation service. The formation of intermetallit phases might degrade tensile ductility and increase the DBTT. The extent of that degradation has not yet be determined. So, it is appropriate to maintain a W content at a level of 2.0% in order to reduce the precipitation potential of Fe,W-type Laves phase during irradiation. However, the preliminary results point out that the proposed low activation ferritic steel shows the promise of improved mechanical properties over both 9Cr-1Mo and modified 9Cr-1Mo steels. 4. Conclusions The preliminary results demonstrated that a fully tempered martensitic alloy can be successfully produced with a nominal composition 9Cr-2.5W-O.3V0.15C steel, acceptable for near-surface waste disposal after fusion reactor service. Quenched and tempered specimens were thermally aged at 600°C up to 1000 h to examine the phase stability of the alloy. X-ray diffraction analysis of the precipitates extracted from the tempered and thermally aged specimens revealed nearly 100% M,,C, carbides (a = 1.064 nm), with no indication of Fe,W-type Laves phase which may cause embrittlement via segregation of detrimental impurity elements and strengthening. The impurity elements (P, S, Sn, Sb, M, Si) are in small concentrations and this, combined with a fine grain size and no precipitation of Laves phase, lead us to conclude little temper embrittlement will occur. The elevated-temperature tensile strength and ductility of GA3X steel are comparable to the average data of commercial heats of 9Cr-1Mo and modified 9Cr1Mo steels up to 650°C. The impact properties and DBTT characteristic of this steel is similar to both 9Cr1Mo steels. The GA3X alloy composition for the optimization of materials properties was evaluated, and the proposed low activation ferritic steels show the promise of improved mechanical properties over 9Cr1Mo steels. The authors wish to express their thanks to the DOE Office of Fusion Energy for financial support of this work under Contract DE-AC03/84ER53158 and to Dr W.L. Hu at HEDL for preparation and testing of the impact specimens. References T.A.
Lechtenberg, Hsu and (1985) 137. [31 ChewYih Hsu and (1985) 120. [41 T.A. Lechtenberg
;:; ChewYih
DOE/ER-0045/14 T.A. Lechtenberg,
(1985) 117. DOEIER-0045/15
T.A. Lechtenberg,
DOE/ER-0045114
et al., in: Proc. of Conf. on Ferritic
Al-
1112
[S]
[6] [7]
[S] [9] [lo] [l I]
C.-Y. Hsu. T.A. Lechrenberg
I Unirradiated low activation ferriric .srcc~l
lays in Nuclear Energy Tech., eds., J. Davis and D. Michel (AIME, New York, 1984) p. 365. J. Watanabe et al., Presented at the ASME 29th Petroleum Mech. Eng. Conf., Dallas, Texas, September 15, 1974. M.K.Booker, Presented to the ASME Working Group on Materials Behavior, New York, February 27. 1985. V.K. Sikka, in: Proc. of Conf. on Ferritic Alloys in Nuclear Energy Tech., eds., J. Davis and D. Michel (AIME, New York, 1984) p. 317. R.L. Klueh and J.M. Vitek, DOE/ER-0045/12 (1984) 100. W.G. Wolfer and R.H. Jones, J. Nucl. Mater. 103-104 (1981) 1305. J.M. Krafft. Appl. Mater. Res. 3 (1964) 88. G.T. Hahn and A.R. Rosenfield, in: Applications Related Phenomena in Titanium Alloys, ASTM STP 432, ASTM (1968) p. 5.
(121 F.H. Huang, Int. J. Fracture 25 (1984) 181. [13] F.H. Huang and R.L. Fish, in: Effects of Radiation on Materials: Eleventh Conf., ASTM STP 7X2. ASTM (1982) p. 701. (141 M.L. Hamilton. F.A. Garner and J.S. Yang, DOE/ER0046/20 ( 1984). 1151 C.Y. Hsu. these proceedings, 2nd Int. Conf. on Fusion Reactor Materials, Chicago, IL, April 13-17. 1986. J. Nucl. Mater. 141-143 (1986). (161 S. Bhattacharyya et al.. IIT Research Institute, Chicago, Illinois (January 1979). [17] D.N. Braski and P.J. Maziasz, J. Nucl. Mater. 1222123 (1984) 338. [18] W.L. Hu, unpublished data. Hanford Engineering Development Laboratory.
MICROSTRUCTURE
AND
1107
MECHANICAL
PROPERTIES
OF UNIRRADIATED
LOW
ACTIVATION FERRITIC STEEL Chen-Yih
HSU and Thomas
A. LECHTENBERG
GA Technologies Inc., P.O. Box 85608, San Diego, CA 92138, USA
Transmission electron micrographs of normalized and tempered 9Cr-2.5W~.3V~.lYZ low activation ferritic steel showed tempered lath-type martensite with precipitation of rod and plate-like carbides at lath and grain boundaries. X-ray diffraction analysis of the extracted replicas revealed nearly 100% M& carbides (a = 1.064 nm), with no indication of Fe,W-type Laves phase even after thermal aging at 600”C/1000 h. Thermal aging increased the number density of rod-like Mz3Chalong prior austenite grain boundaries and martensite lath boundaries. The elevated-temperature tensile strengths of this steel are about 10% higher than the average strengths of commercial heats of 9Cr-1Mo and modified 9Cr-1Mo steels up to 650°C with equivalent uniform elongation and -50% decrease in total elongation. The DBTT was determined to be -25°C which is similar to other 9Cr-1Mo steels. Fractographic examination of tensile tested specimens shows a mixed mode of equiaxed and elongaied dimples at test temperatures above 400°C. Modification of the GA3X alloy composition for optimization of materials properties is discussed. However, the proposed low activation ferritic steel shows the promise of improved mechanical properties over 9Cr-1Mo steels. 1. Introduction available alloys to be used for first structural components will become during service and will pose a waste disposal
Commercially
wall and blanket activated
problem after service. Guidelines (10CFR61) for the classification of nuclear wastes have been issued by the US Nuclear Regulatory Commission. In 1982, DOE convened a panel to examine the issues and incentives associated with the development of low activation materials. The available data show that the 9-12%Cr ferriticl martensitic steels are attractive materials for fusion reactor structural component applications. The primary objective of this work is to develop ferritic steels that exhibit equivalent mechanical properties and performance to the commercially developed alloys, HT-9 (12Cr-1MoVW) and/or 9Cr-1MoVNb steels, but which may allow near-surface disposal after fusion reactor service. These steels are hereafter referred to as low activation. To achieve this, the elements MO, Ni, Nb, and N, which are added for elevated temperature strength, must be reduced to very low levels. W and V can be used to substitute for MO. In equal atomic fractions, W and MO show similar solid solution hardening characteristics and Fe-W-C alloys develop analogous precipitates with a similar precipitation sequence to Fe-MO-C alloys. Therefore, the Cr-W steels offer promise for the development of a replacement for Cr-Mo steels. An experimental heat (GA3X) of low activation ferritic steel was prepared with a nominal composition of 9Cr-2.5W4.3V4.15C steel. The addition of 0.3% V to Cr-W steels will result in a pronounced effect on the precipitate formation and enhance the elevated temperature properties, but may also make the steel less weldable. For this reason, the suggested carbon content is kept below 0.15% in Cr-W steels. Titanium is a strong carbide former with an evi0022-3115/86/$03.50 0 Elsevier Science Publishers (North-Holland Physics Publishing Division)
dence of inhibiting grain coarsening while retaining good notch toughness, so the addition of 0.1% Ti is also suggested as a modification to GA3X although work on this not reported here. This paper summarizes the metallographic examination and mechanical property evaluation of the unirradiated GA3X low activation ferritic steel compared to commercial heats of 9Cr1Mo steels, and also provides baseline data for a comparison with the irradiated GA3X steel. 2. Experimental
procedures
The melting and fabrication practice of the GA3X low activation ferritic steel are described in ref. [l]. The chemical composition of this GA3X heat is shown in table 1. The specimens discussed here were water quenched from lOOO”C/lh and then tempered at 7OO”C/l h. A two-stage replication technique was used to prepare the carbon film extraction replicas of the precipitates [2]. For X-ray diffractometry analysis, 5 mg of the precipitate were smeared onto a single crystal quartz slide in a graphite crystal monochromator-scintillation detector system. In the preparation of thin-foil specimens, 3 mm disks were electropolished in a solution of 50% ethyl alcohol, 45% butyl cellosolve, and 5% HC104. Carbon film extraction replicas and thin-foil samples were examined on-a Philips 300M transmission electron microscope operating at 100 kV.
Table 1 Chemical composition of GA3X ferritic steel (wt%) C
Cr
W
V
Mn
S
0.16-0.17
7.4-7.6
1.9-2.0
0.014-0.016
0.002
0.005
B.V.
1108
C.-Y. Hsu, T. A. Lechtenherg
I Unirradiated
The tensile tests were performed on flat tensile specimens with a reduced gage section 25.4 mm long by 3.18 mm wide by 1.0 mm thick. The flat tensile specimens were dipped in toluene pretreat coating solution in order to avoid any significant oxidation during the tensile test at temperature in the range of 2@-65O”C. Fractographic examination of tensile tested specimens was performed on a scanning electron microscope. The ductile brittle transition temperature (DBTT) and impact properties were determined from f size miniature Charpy V-notch specimens (23.6 mm x 3.33 mm x 3.33 mm with a 30” notch angle and 0.5 mm notch depth). The miniature Charpy V-notch specimens were fatigue precracked to approximately a/w = 0.5, and were then tested in the instrumented impact testing machine at the Hanford Engineering Development Laboratory. 3. Results and discussion 3.1. Metallographic examination The weight percentage of the extracted precipitates was increased from 2.0% of the quenched and tempered specimens to 3.1% after thermal aging at 600°C for 100 h and then remained unchanged up to 1000 h. Further, there was no indication of Fe,W-type intermetallic compound in any of the samples analyzed [3]. X-ray diffraction analysis of the extracted precipitates revealed nearly 100% M,,C, carbides with a lattice parameter of 1.064 nm. The X-ray fluorence analysis indicated that the extracted M&carbides were in the form of (Feo,21Wo,osCro,,o)23C6. There was no change of metal contents in M& carbides after thermal aging treatment. Fig. 1 shows TEM micrographs of extraction replicas prior to and after thermal aging at 600°C for 100 h. The extracted precipitates show two types of carbides, with rod-like and plate-like shapes, similar to results on HT-9 steel reported by Lechtenberg et al. [4]. The selected area diffraction patterns from the isolated particles confirmed the carbides to be M&, and a few particles seem likely to be MC-types carbide. This is consistent with the X-ray diffraction analyses. The scarcity of MC-type carbide is due to very low V content (0.015%) and no Ti added in this steel. An appreciable amount of MC-type carbide should be detected if the target value of 0.3% V or Ti were added to this steel. The size of precipitates was typically in the range of 20-750 nm. The mean particle size increased from 100 nm to 140 nm after thermal aging at 600°C for 100 h and many more rod-like carbides were found in the thermally aged specimens than in the astempered specimens. Thin-foil TEM micrographs show tempered lath martensite within the prior-austenite grains, as shown in fig. 2. No significant amount of ferrite was seen. In
low activation ferritic strrl
both the tempered and the thermally aged specimens, heavy carbide precipitation was noted at the prioraustenite grain boundaries and martensite lath boundaries. The areal density of carbide precipitation is appreciably higher in the thermally aged specimen than in the tempered specimen as judged directly from the thin-foil microstructural observation. Particle coalcscence and an increase in the areal density of rod-like carbides were found on grain and lath boundaries after the thermal aging treatment. This observation can be attributed to enhanced diffusion at grain and lath boundaries. The evidence shows that the number density of carbides and carbide size increased, and particle morphology tended to change during thermal aging treatment. The precipitation process and morphology change of particles and the increase of elongated rodlike carbides along grain boundaries and lath boundaries may influence the impact and fracture toughness of this steel, as well as other mechanical behavior such as tensile and creep strengths. In the quenched and tempered condition, dislocations are ordered in such a way as to produce irregular subgrains within the martensite laths. After thermal aging at 600°C for 100 h there were no apparent
Olwn
Fig. 1. TEM micrographs of carbon film extraction replica of GA3X steel, (a) quenched and tempered condition, and (b) followed by thermal aging at 6OO”C/lOOh.
C.-Y. Hsu, T. A. Lechtenberg
1109
I Unirradiated low activation ferritic steel
higher than the average value for commercial heats of 9Cr-1Mo and modified 9Cr-1Mo steels at temperatures up to 400°C and are slightly higher at temperatures above 500°C. This may be partly due to the low tempering temperature (700°C) given to GA3X whereas 9Cr-1Mo and 9Cr-1MoVNb are tempered in the range of 740-760°C. A significant contributing factor may be the alloy content. The actual Cr content of this steel is only 7.5% (target 9%) and the concentration of W is only 2.0 wt% (target 2.5 wt%). The mod-
600
a
8
AVERAGE CURVE OF MOOlFlEO SCr-1Mo
. NOMINAL
100 0' 0
Fig. 2. TEM thin-foil micrographs of GA3X steel, (a) quenched and tempered condition, and (b) followed by thermal aging at 6OO”C/lOOh. changes in the dislocation sub-structures. Changes in dislocation structure may become more pronounced after prolonged thermal aging at service temperatures or during irradiation. Resistance to temper embrittlement is controlled by maintaining low levels of impurity elements such as P, S, Sb, Sn, Mn, and Si. Furthermore, the formation of intermetallic compounds such as Laves phases may promote the segregation of impurity elements and thus degrade the resistance to temper embrittlement. In the GA3X steel, the Mn and S contents are less than 0.005% while the other impurity elements are negligible. Based on the analysis proposed by Watanabe et al.[5], the effect of these impurities combined with a fine grain size lead us to conclude little temper embrittlement will occur.
COMPOSITION:
A SCr-1MoV Nb (REF.51
I
I
1
I
1
I
100
200
300
400
500
600
700
TEST TEMPERATURE,°C 900
700
g
600
AVERAGECURVEOF MODIFIED SCr-1Mo
z $ 500 I k 9 400 Ly > k t
300
:: 2 200
NOMINAL COMPOSITION: 0 9Cr-2.5W-0.3V-0.l5C
109
A SCr-1MoV Nb IfiEFs 51 0
3.2. Mechanical
properties
0
100
200
300
400
500
600
700
TEST TEMPERATURE,°C
The ultimate tensile strength and 0.2% offset yield strength of this ferritic steel are compared with the average values of 9Cr-1Mo and modified 9Cr-1Mo steels [6,7] in fig. 3. Tensile strengths are about 10%
Fig. 3. The tensile strengths of 9Cr-2.5W4I.3V41.5C (GA3X) steel compared to average curves of 9Cr-1Mo and modified 9Cr-1Mo steels (band), (a) ultimate tensile strength, and (b) 0.2% offset yield strength.
1110
C.-Y. Hsu. T.A. Lechtenberg I Unirradiated low activation ferritic steel
ified 9Cr-1Mo steels contains 0.21 wt% V and 0.01 wt% Ti and GA3X steel contains only 0.015 wt% V. Hence, the enhanced tensile strength observed in GA3X steel is likely to be due to the effect of a 0.17% carbon content, which is higher than carbon content (0.1%) of both 9Cr-1Mo steels. In comparison with the experimental heat of unirradiated 9Cr-IMoVNb steel [8], the GA3X steel shows higher strengths and ductility from 20°C to 4OO”C, and becomes slightly weaker at temperatures above 500°C but still maintains a higher ductility. In contrast to 0.015% V in GA3X steel, the 9Cr-1MoVNb steel contains 0.24% V and 0.08% Nb which contributed significantly to the elevated temperature strengths. So, the strength of GA3X steel may be significantly improved if the alloying elements added for elevated-temperature strength had been closer to target contents. It is indicated that the low activation ferritic steel with a nominal composition of 9Cr-2.5W-O.3Va.15C may show a promising potential of improved strength at elevated temperatures, compared to the 9Cr-1Mo steels. The uniform and total elongations of GA3X steel and the average data of commercial heats of 9Cr-1Mo and modified 9Cr-1Mo steels are shown in fig. 4. The uniform elongation of this steel is about equivalent to that of modified 9Cr-1Mo steel while the total elongation is -50% less than both 9Cr-1Mo steels. The workhardening exponent is approximately equivalent to the uniform elongation. This indicates that the workhardening response of this steel is similar to that of 9Cr-1Mo steels which indicate similar precipitation effects on dislocation movement and multiplication. Changes in fracture toughness (K,,) can be estimated from tensile behavior and the associated microstructural evolution of the steel [9-141. An excellent agreement between tensile-base prediction and KIC fracture toughness measurements was obtained on irradiated stainless steel by using the modified Krafft correlation [13]. In the case of low work-hardening coefficients, the modified Krafft correlation can be simplified as: Kit-
(oYer)lh,
In the fractographic examination of the tensile tested specimens [15], a ductile dimple rupture caused by a void coalescence mechanism was observed for specimens tested from 20°C up to 400°C. At temperatures above 500°C the fracture mode was still ductile as shown by the dimples, but more than 30% of the dimples were elongated, which was usually indicated to be the result of ductile tearing [16]. This mixed fracture mode suggests that there might have been a strong shear component and/or cross slip involved in the frac-
AVERAGECURVEOF MODIFIED SCr-1Yo
0
100
200
300
400
500
700
600
6OC
TEST TEMPERATURE,PC
t AVERAGECURVE OF MODIFIED SCr-1Mo
AVERAGE CURVE OF SCr-1Mo STEEL
(I)
where a,, is the yield strength and .sris the true fracture strain. The opposing changes in yield strength and ductility tend to counterbalance each other. The workhardening coefficients of GA3X ferritic steel are 0.1 at lower temperatures and become 0.06 at higher temperatures. In the same class of 9Cr ferritic steels, the variation of K,, fracture toughness can be estimated by comparing the differences of tensile properties. Based on eq. (l), the plane-strain KIC of GA3X steel is predicted to be about lO-20% lower than the commercial heats of modified 9Cr-1Mo steel [6], and is expected to be slightly higher than results on an experimental heat of 9Cr-1MoVNb steel at the same temperatures [71.
0
100
200
300
400
500
600
700
600
TEST TEMPERATURE,%
Fig. 4. Elongation (a) uniform
data vs test temperature on GA3X elongation, and (b) total elongation.
steel,
C.-Y. Hsu, T.A. Lechtenberg
I Unirradiated low activation ferritic steel
ture at the high temperatures (?0.47’,,,). Braski and Maziasz [17] studied the elevated-temperature tensile properties of PCA austenitic stainless steel and found shear-elongated dimples at test temperatures from 300°C to 600°C in the 25% cold-worked condition. Once the cold-worked material was heat treated at 75O”C/2 h, MC-type carbides precipitated and the mixed-mode vanished. It is suggested that the shear elongated dimples may be eliminated with more MC carbide precipitation as would be the case by increasing V and/or Ti contents of GA3X heat. However, the reason for the formation of shear dimples is not yet fully understood and deserves further investigation. The impact test results on the precrackedi size Charpy specimens of this ferritic steel are tabulated in table 2. The ductile-brittle transition temperature was determined to be -25°C with an upper shelf energy of 2.2 J (37.2 J/cm2 normalized fracture energy). The DBTT and impact energy of this low activation ferritic steel is very similar to the modified 9Cr-1Mo steels ]I81. The experimental results of GA3X steel provide some clues concerning optimum alloy composition. The mechanical properties of steel can be optimized by controlling microstructure through processing and by modifying alloy composition. For example, significant strengthening was contributed by the 0.17% carbon content. High carbon content can make the steel less weldable, and increase particle coarsening may degrade ductility and fracture toughness. The carbon content should be reduced to 0.15% or lower. The decrease in strength due to the lowering of carbon content can be compensated by increasing Cr, V, and Ti alloying addition. The Cr content is to be maintained at a level of 9% to provide oxidation resistance at elevated temperatures. The phase stability is a key concern for the justification of alloy composition. V and Ti are strong carbide formers. The addition of 0.3% V and
Fracture
Maximum
Normalized
temperature
energy (J)
load (kN)
energy
100°C 50°C 21°C 0°C -10°C -25°C -50°C -80°C
2.225 2.499 2.264 2.180 1.658 1.113 0.498 0.050
0.645 0.733 0.693 0.726 0.648 0.574 0.571 0.536
37.56 38.29 36.11 34.10 27.16 17.73 8.01 0.85
(J/cm2)
1111
intermetallic phase precipitation, such as Fe2W type Laves phase, after a prolonged thermal treatment or during irradiation service. The formation of intermetallit phases might degrade tensile ductility and increase the DBTT. The extent of that degradation has not yet be determined. So, it is appropriate to maintain a W content at a level of 2.0% in order to reduce the precipitation potential of Fe,W-type Laves phase during irradiation. However, the preliminary results point out that the proposed low activation ferritic steel shows the promise of improved mechanical properties over both 9Cr-1Mo and modified 9Cr-1Mo steels. 4. Conclusions The preliminary results demonstrated that a fully tempered martensitic alloy can be successfully produced with a nominal composition 9Cr-2.5W-O.3V0.15C steel, acceptable for near-surface waste disposal after fusion reactor service. Quenched and tempered specimens were thermally aged at 600°C up to 1000 h to examine the phase stability of the alloy. X-ray diffraction analysis of the precipitates extracted from the tempered and thermally aged specimens revealed nearly 100% M,,C, carbides (a = 1.064 nm), with no indication of Fe,W-type Laves phase which may cause embrittlement via segregation of detrimental impurity elements and strengthening. The impurity elements (P, S, Sn, Sb, M, Si) are in small concentrations and this, combined with a fine grain size and no precipitation of Laves phase, lead us to conclude little temper embrittlement will occur. The elevated-temperature tensile strength and ductility of GA3X steel are comparable to the average data of commercial heats of 9Cr-1Mo and modified 9Cr1Mo steels up to 650°C. The impact properties and DBTT characteristic of this steel is similar to both 9Cr1Mo steels. The GA3X alloy composition for the optimization of materials properties was evaluated, and the proposed low activation ferritic steels show the promise of improved mechanical properties over 9Cr1Mo steels. The authors wish to express their thanks to the DOE Office of Fusion Energy for financial support of this work under Contract DE-AC03/84ER53158 and to Dr W.L. Hu at HEDL for preparation and testing of the impact specimens. References T.A.
Lechtenberg, Hsu and (1985) 137. [31 ChewYih Hsu and (1985) 120. [41 T.A. Lechtenberg
;:; ChewYih
DOE/ER-0045/14 T.A. Lechtenberg,
(1985) 117. DOEIER-0045/15
T.A. Lechtenberg,
DOE/ER-0045114
et al., in: Proc. of Conf. on Ferritic
Al-
1112
[S]
[6] [7]
[S] [9] [lo] [l I]
C.-Y. Hsu. T.A. Lechrenberg
I Unirradiated low activation ferriric .srcc~l
lays in Nuclear Energy Tech., eds., J. Davis and D. Michel (AIME, New York, 1984) p. 365. J. Watanabe et al., Presented at the ASME 29th Petroleum Mech. Eng. Conf., Dallas, Texas, September 15, 1974. M.K.Booker, Presented to the ASME Working Group on Materials Behavior, New York, February 27. 1985. V.K. Sikka, in: Proc. of Conf. on Ferritic Alloys in Nuclear Energy Tech., eds., J. Davis and D. Michel (AIME, New York, 1984) p. 317. R.L. Klueh and J.M. Vitek, DOE/ER-0045/12 (1984) 100. W.G. Wolfer and R.H. Jones, J. Nucl. Mater. 103-104 (1981) 1305. J.M. Krafft. Appl. Mater. Res. 3 (1964) 88. G.T. Hahn and A.R. Rosenfield, in: Applications Related Phenomena in Titanium Alloys, ASTM STP 432, ASTM (1968) p. 5.
(121 F.H. Huang, Int. J. Fracture 25 (1984) 181. [13] F.H. Huang and R.L. Fish, in: Effects of Radiation on Materials: Eleventh Conf., ASTM STP 7X2. ASTM (1982) p. 701. (141 M.L. Hamilton. F.A. Garner and J.S. Yang, DOE/ER0046/20 ( 1984). 1151 C.Y. Hsu. these proceedings, 2nd Int. Conf. on Fusion Reactor Materials, Chicago, IL, April 13-17. 1986. J. Nucl. Mater. 141-143 (1986). (161 S. Bhattacharyya et al.. IIT Research Institute, Chicago, Illinois (January 1979). [17] D.N. Braski and P.J. Maziasz, J. Nucl. Mater. 1222123 (1984) 338. [18] W.L. Hu, unpublished data. Hanford Engineering Development Laboratory.