where the temperature (T) is in °C. The correlation coefficient for the plotted data using this equation is R=0.8955.

Fig. 1. Ultimate tensile strength of unirradiated 8-9%Cr steels [1-6].

Fig. 2. Yield strength of unirradiated 8-9%Cr steels [1-4,6].

Neutron irradiation causes a pronounced increase in the yield and ultimate tensile strength of 8-9%Cr conventional and reduced activation steels at temperatures below ~400°C, but has little effect on the strength at higher temperatures. Figure 3 shows a comparison of the unirradiated and irradiated yield strength for irradiation temperatures between 50 and 600°C. The effects of fusion-relevant helium generation on the yield and ultimate strength have not been adequately studied, although only minor changes have been observed in studies performed to date [7-9].

Fig. 3. Comparison of the yield strength of unirradiated and irradiated 8-9%Cr steels [6].

The uniform elongations of unirradiated 8-9%Cr conventional and reduced activation steels exhibit relatively low values (<5%) at all temperatures from 20 to 700°C [6]. As shown in Fig. 4, the unirradiated uniform elongation decreases slowly from ~5% to ~1% over this temperature range. The corresponding total elongations range from ~10 to 30%. As shown in Fig. 5, irradiation causes a decrease in the uniform and total elongations, particularly for irradiation temperatures below 400°C [1-3,6,9,10]. Data for both reduced-activation and conventional Fe-(8-9)%Cr steels are plotted in this figure. The uniform elongation is very low (<3%) at all investigated irradiation temperatures. The total elongation remains above ~7% for all irradiation conditions investigated to date, and it increases with increasing irradiation temperature (Fig. 5). The effects of fusion-relevant helium generation on tensile elongation have not been adequately studied.

Fig. 4. Uniform and total elongation of unirradiated 8-9%Cr steels [1-3,6].

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Fig. 5. Uniform and total elongation of irradiated 8-9%Cr steels [1-3,6]. The test temperature equals the irradiation temperature.

The reduction in area (RA) as measured on unirradiated and irradiated 8-9%Cr steel tensile specimens is shown in Fig. 6 [1,3,6]. The unirradiated reduction in area is high (>80%) at all test temperatures between 20 and 700°C. Irradiation causes a decrease in the RA (particularly at low irradiation and test temperatures), but the reduction in area remains acceptably high in the limited number of tensile specimens examined to date. Helium effects have not been adequately studied.

Fig. 6. Reduction in area of unirradiated and irradiated 8-9%Cr steels [1,3,6].

Figure 7 shows representative stress-strain curves obtained on miniature "type SS-3" sheet tensile specimens (0.76 x 1.52 x 7.6 mm gage dimensions) for F82H steel tensile tested at a strain rate of 1.1x10^-3 s^-1 following neutron irradiation at 200-600°C [6,9,10]. Pronounced flow localization is observed for irradiation temperatures below ~400°C, whereas adequate strain hardening capacity occurs at temperatures >400°C.

Fig. 7. Load vs. normalized crosshead displacement tensile curves for F82H irradiated to 3-34 dpa at 200-600°C [6,9,10].

The elastic constants for F82H (8Cr-2WVTa) have been measured between 20 and 700°C [3,5], and Young's modulus has been measured from -150 to 350°C for several other reduced-activation steels including Fe-9Cr alloys [11]. The temperature-dependent elastic constants for F82H exhibit approximately bilinear behavior, which a slope change occurring near 450-500°C. The change in the temperature dependence of the elastic constants at 450-500°C was attributed to annealing effects on the martensitic structure at the higher temperatures [3,5]. The following equations for Young's modulus (E_Y) and the shear modulus (G) are obtained from the F82H experimental data [3,5] in the temperature interval between 20 and 450°C:

At temperatures above 450°C, Young's modulus for F82H decreases approximately linearly from E_Y=193 to 160 GPa as the temperature is increased from 450 to 700°C. The shear modulus similarly decreases from G= 75 to 60.5 GPa as the temperature is increased from 450 to 700°C. Poisson's ratio (n=(E_Y/2G) - 1) is constant up to 500°C with a value of 0.29, and then slowly increases to 0.31 at 700°C.

8. Stress-rupture

There have been numerous studies of the creep and stress-rupture behavior of unirradiated and irradiated 8-9%Cr steels at temperatures up to 650°C (0.5 T_M) [1,2,12]. Good creep resistance exists for temperatures up to ~550°C (0.45 T_M), but poor creep resistance occurs at 600°C and above. For example, the 10,000 h creep rupture strength of F82H is 200 MPa at 550°C, 120 MPa at 600°C and 50 MPa at 650°C [1]. Improvements in the thermal creep resistance of reduced-activation ferritic steels can be achieved with oxide dispersion strengthened alloys [13].

9. Thermal expansion, specific heat and thermal conductivity

The thermophysical properties for several different heats of F82H (8Cr-2WVTa) have been measured from room temperature to 700°C [1,3,5]. The mean coefficient of thermal expansion (a_th) varied from 10.4 ppm/°C at room temperature to 12.4 ppm/°C at 700°C. The specific heat at constant pressure (C_P) varied from 470 J/kg-K at 20°C to 810 J/kg-K at 700°C. The specific heat increase was approximately linear with temperature between 20 and 500°C, and was strongly nonlinear above 500°C. The thermal conductivity at 20-800°C was determined from thermal diffusivity measurements using laser flash techniques and was found to be nearly independent of temperature, with an average value of 33 W/m-K between 20 and 700°C for two different heats of F82H.

10. Ductile to brittle transition temperature (unirradiated and irradiated)

The measured value of the ductile to brittle transition temperature (DBTT) in body-centered cubic materials depends on numerous experimental parameters, including the specimen geometry, strain rate, and the sharpness of the notch where the crack is initiated (notch acuity) [14]. The measured DBTT in miniature unirradiated F82H machined Charpy vee-notch (MCVN) specimens ranges from about -60°C to -110°C, where the lower DBTT was obtained using 1.5 mm thick Charpy impact specimens [2,9,15]. The DBTT for unirradiated F82H measured on precracked compact tension specimens is near -50°C [16]. Low temperature irradiation causes a moderate increase in the DBTT (DDBTT~20 to 100°C) of advanced reduced activation 8-9Cr steels for the damage levels investigated to date [2,9,10,15,17,18]. Very small changes in the DBTT (DDBTT <50°C) have been observed in 8-9%Cr steels irradiated at temperatures above 400°C. The effect of fusion-relevant helium generation rates on the DBTT in irradiated specimens has not been adequately studied.

11. Recommended reference operating temperature limits

The maximum operating temperature limit for 8-9%Cr reduced activation ferritic/martensitic steels is ~550°C, due to thermal creep considerations. Somewhat higher temperatures could be tolerated for components exposed to low mechanical stresses. Oxide dispersion strengthened ferritic alloys under development may be capable of operation up to temperatures above 550°C [9]. Additional work on irradiated specimens is needed before the minimum operating temperature limit can be established. The reference minimum operating temperature limit will be controlled by radiation hardening, which causes loss of ductility and an increase in the ductile to brittle transition temperature. According to the available irradiation data, the DBTT of 8-9%Cr steels remains near or below room temperature following neutron irradiation at temperatures between 200 and 550°C [2,9,10,15,17,18]. There is some limited evidence that fusion-relevant helium generation rates may cause a further increase in the DBTT beyond that attributable to matrix hardening (defect cluster) effects [8,9,17]. Further work is needed to determine the effect of helium on fracture properties. The proposed reference minimum operating temperature for 8-9%Cr steels is 250°C.

References

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[15] A. Kohyama, in: Proc. IEA Working Group Meeting on Ferritic/Martensitic Steels, Culham, UK, October 1996, ORNL/M-5674, ed. R.L. Klueh, Oak Ridge National Lab, 1996.

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[18] L.E. Schubert, M.L. Hamilton, D.S. Gelles, in: Fusion Materials Semiann. Progress Report for Period ending June 30, 1996, DOE/ER-0313/20 (Oak Ridge National Lab, 1996) p. 171.