Microstructural Precipitation Effects--Gelles a- 123docz.net

7. Gas Effects--Ferritic steels without nickel have lower helium generation rates during FBR irradiation than austenitic steels because nickel has a much higher (n,a) cross section than iron or chromium. All of the studies on the effect of helium [7-10,19,26-- 29] point to its stimulation of void nucleation. Triple ion beam experiments also show

an enhanced effect of helium and hydrogen together [19,27-29].

Most of these mechanisms relate to the vacancy supersaturation available for void nucleation and growth. However, gas effects (g) and the critical size rL for conversion of a stable gas bubble into a bias-driven void [32,40,41] are most relevant to void nucleation, which is the major issue addressed by our data. While most of the above mentioned mechanisms probably contribute in concert to the void swelling resistance observed in these steels during FBR irradiation, our data suggest that mechanisms 3) and 4) are most important to void growth during irradiation at higher helium generation levels.

Our data indicate that void nucleation is very difficult during FBR irradiation with little or no helium and that increased helium generation greatly enhances void nucleation. Void nucleation can occur in austenitic stainless steels even without any helium present due to oxygen effects [42,43]. Elemental nickel used for alloying can have very high oxygen contents, usually higher than the chromium or iron starting stock used to produce austenitic alloys and steels-'. The martenistic/ferritic steels contain little nickel and much higher carbon contents (carbon is a very efficient deoxidizing element, as are silicon and titanium) compared to austenitic steels and alloys, so that they most likely have less oxygen available for void nucleation and are therefore more dependent on helium generated during irradiation.

Macroscopic swelling with increasing fluence can be described in terms of a low swelling transient regime or incubation period, followed by a steady-state regime of much more rapid swelling [6,20]. Voids nucleate and grow during the low swelling transient regime, and usually rapid void growth or coalescence or both cause the onset of rapid swelling. Our data suggest that, while helium appears to have accelerated void nucleation to shorten the incubation period, perhaps by 75 to 100 dpa or more, none of our microstructures appears to be at the end of the transient regime and certainly not fully into the rapid swelling regime.

Two closely related critical parameters from rate theory, the critical radius, r~, and the critical number of gas atoms, n f , whose mathematical equations are described by Odette

[32] and Mansur and Coghlan [4l], help to interpret our microstructural data. Our data suggest empirical estimates for r,; ng* is more difficult to estimate, but can be related to the duration of the incubation period. The rc for the undoped 9Cr and 12Cr steels in this work appears to be <0.75 nm at 400 to 410~ in both FFTF and HFIR, whereas in the nickel doped steels rc increases from that value in FFTF to 2.5 to 3 nm in HFIR. The void formation in 9Cr-IMoVNb in F F T F may also indicate a very low value of n~* in that steel.

At low values of both r, and ng*, more helium would readily lead to more voids, as we observe for the undoped 9Cr and 12Cr steels irradiated in HFIR. Void formation appears more difficult in the nickel-doped steels in HFIR because r, and, presumably ng*, increase.

But void nucleation still increases. The value of r, for the undoped steels appears to be consistent with that expected by others at about 400~ based on ion irradiation data [27,28,40],

but the value of 2.5 to 3 nm in the nickel doped, HFIR irradiated steels seems large, as large as Horton and Mansur [40] find for Fe-10Cr ion irradiated at 580~ to 30 dpa. In SA 316 irradiated at 425 to 450~ voids form easily and r, appears to be - 1 . 5 nm. while in

: Chemical analyses on steels melted at ORNL, 1986 to 1988.

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56 EFFECTS OF RADIATION ON MATERIALS

C W 316 or titanium-modified austenitic stainless steels that are resistant to void formation in H F I R at 400~ r > 2.5 to 5 nm [20.45].

Void formation and growth are also affected by the net bias or vacancy supersaturation and defect partitioning. The H F I R results on the nickel doped steels with large r may suggest a larger vacancy supersaturation than might be anticipated and a more balanced partitioning of point defects among various sinks, despite the high density of bubbles. We can calculate a point defect partitioning factor Q for either vacancies or intersitials, defined for a two sink system as [41]

Q = Z"A/4rrrN, Z' (1)

where Z '~ and Z' are the capture efficiencies of dislocations and cavities, respectively for that particular defect (usually -- l ). X is the dislocation density, r is the average cavity radius.

and N, is the cavity concentration. H o r t o n and Mansur [40] give an expression for the sink strength that includes a bimodal cavity distribution. The sink strength S for vacancies or interstitials can be expressed generally to include other sinks as follo~vs.

S = 4 Z , r , N , , Z ( r , ) + L Z " + I'Z" + 4 Xzrr.,N; Z,,(r.,,,) (2) where the first term sums o v e r a cavity distribution with i size classes, 1" is concentration of subgrain boundaries, i is the n u m b e r of precipitate types with average size r:, and concentration Np, and Z" and Z p are the capture efficiencies of subgrain boundaries and precipitates.

respectively, for vacancies or interstitials. Using the simple form of Eq 1. our steels irradiated in F F T F have Q > 40 so that they appear to be dislocation sink dominated. The values of Z ~ and Z ' are assumed to equal one and values for the other parameters in Eq t are obtained from the experimental data. The undoped steels irradiated in H F I R have Q -- 0.2. while the nickel d o p e d steels have Q - 0.4, including the contribution of the bimodal cavity distribution. Clearly the H F I R irradiated specimens are closer to a balanced situation than the F F T F irradiated steels, and vet do not appear to have a sink structure totally d o m i n a t e d by cavities as do many cold worked austenitic stainless steels irradiated in H F I R . The calculated Q values for the H F I R irradiated steels could be either closer to or further from a balanced sink situation, d e p e n d i n g on how the precipitates and subgrain boundaries com- pete with o t h e r sinks for point defects. The finer subgrain boundary structure may be a factor in reducing void formation and growth in the 1 2 C r - I M o V W - 2 N i irradiated in H F I R . Considering the bias, theoretical work suggests that the bias (B) in the ferritic steels can range from 0.05 to 0.4 [40]. The normal B for austenitic steels is II. 15 to IL2 [32]. We cannot really estimate bias simply from void formation without knowing growth rates. H o w e v e r , enough bias seems to exist for substantial void growth in these ferritic steels. F r o m the various mechanisms that contribute to a low bias in ferritic steels, solute effects and dislocation nature effects appear constant in our comparison of F F T F and H F I R irradiation for each steel and appear small among the various steels in either reactor. Helium generation itself could counter the low intrinsic bias suggested in mechanism l. it helium trapped in vacancies hinders recombination with interstitials as suggested for austenitic steels [6,20].

Helium vacancy complexes may also be m o r e mobile in the ferritic steels if the vacancv self- diffusion rate is higher. Increased helium accumulation at various sink~, governed by D, C, (vacancy diffusivitv times vacancy concentration) in the ferritic steels, could also alter defect capture efficiences. A p a r t from defect-partitioning effects, subgrain boundaries and precipitates could also affect the bias as well. Fine precipitates could contribute to an increase in the bias if they preferentially attract interstitials on the basis of their ~olumetric misfit (undersized misfit would attract interstitials). Fine precipitates do not appear to contribute

MAZIASZ AND KLUEH ON VOID FORMATION AND HELIUM EFFECTS 57

very much to void s~elling resistance from this work, because void formation in HFIR often coincides with their formation. Previous data at 500~ actually showed that fine M~C de~elops cooperatively with voids in HFIR-irradiated 12Cr-1 MoVW [10].

Finally, what do these results mean for the use of ferritic steels for fusion applications?

While swelling in these 9Cr and 12Cr steels is not high and may not increase greatly at higher doses, several factors could still contribute to higher eventual void swelling rates during fusion irradiation than found during FBR irradiation. The voids in the steels with the fusion levels of helium could continue to grow. while new voids still form. These processes could then reduce the number of fine bubbles to lower r, and achieve more balanced defect partitioning, which would further increase swelling. Odette [32] showed that an increase in bias and in the ultimate void density (a normal helium effect in austenitic stainless steelsl could result in the ferritic steels swelling in a manner similar to 2 0 q cold worked Type 316 stainless steel. If voids develop on coarse precipitates, swelling could be enhanced still further. Higher fluence experiments are necessary, but until then, the possibility of helium enhanced void swelling remains a legitimate concern for fusion.

Summary and Conclusions

1. By comparing nickel doped and undoped 9Cr-IMoVNb and 12Cr-IMoVW steels irradiated in FFTF (47 dpa. - 5 appm helium) and HFIR ( - 3 7 dpa, 400 to 430 appm helium), increases in He/dpa ratio were found to cause significant increases in the formation of large (7-30 nm diam.), bias-driven voids at 400 to 410~ Only the nickel doped steels irradiated in HFIR with 400 appm helium or more had visible helium bubbles, which were very fine (2 to 5 nm in diameter) and abundant ( l t o 4 x 1 0 " m ').

2. Irradiation in both reactors produced a spatially uniform network of tangled dislocations and some larger loops, with .'~ being 0.6 to 4 x l i p m ;. Irradiation produced a structure that was different than the initial as-tempered, spatially nonuniform structure in which A varied from <10" to 7 x 10 ~' m 3. Irradiation significantly increased \ only in the 12Cr- 1MoVW steel for which had A was initially very low (<10" m -'1. Irradiation at 400 to 410~

produced almost complete recovery of the as-tempered lath subgrain boundarx structure in the 9Cr-IMoVNb and 9Cr-IMoVNb-2Ni steels. Such boundaries remained stable in similarly irradiated 12Cr-IMoVW and 12Cr-IMoVW-2Ni steels.

3. Irradiation produced some significant changes in the precipitate structure on all of the steels irradiated at 400 to 410~ in both reactors. There was dissolution of many of the coarse, as-tempered M,,C,, particles during irradiation of both 9Cr steels, whereas similar particles were relatively more stable in the 12Cr steels. Irradiation produced an abundant dispersion of new coarse M23C, particles only in the 9Cr-IMoVNb-2Ni steel. Finer MC precipitate particles, present in all the steels, experienced some coarsening and compositional changes (chromium enrichments and ~anadium depletions l during irradiation. Irradiation produced abundant dispersions of fine M,C (q) in all the steels except 9Cr-IMoVNb. These particles were silicon, chromium, and nickel rich in the nickel doped steels, but bad an odd composition with only silicon and chromium in the 12Cr-IMoVW steel. Some coarse M(,C ('q) (silicon, chromium, and nickel rich) was found to replace coarse as-tempered M:,C~, during irradiation in both nickel doped 9Cr and 12Cr steels.

4. Although irradiation at 40(I to 4UI~ bad considerable effects on both the dislocation and subgrain boundary and the precipitate components of the microstructure in all the steels, these changes were nearly the same comparing irradiation in FFTF and HFIR for each heat of steel.

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58 EFFECTS OF RADIATION ON MATERIALS

indicates that there is a sufficient bias for void growth in ferritic, martensitic steels irradiated at about 400~ if there is e n o u g h helium for voids to nucleate. O u r data indicate that the critical size and critical gas content for the conversion of stable gas clusters or bubbles ( r and n,*, respectively) to voids are quite small (r. < 0.75 nm) in the 9 C r - I M o V N b and 12Cr- 1 M o V W steels irradiated in both reactors, because m o r e helium causes more voids without producing resolvable bubbles. The higher helium content of the nickel-doped steels irradiated in H F I R produces many resolvable, but subcritical, helium bubbles, which increases r, to 2.5 to 3 nm: however, voids still form and grow, indicating the presence of some bias. All of the steels irradiated in F F T F a p p e a r to have point defect annihilation d o m i n a t e d by the sink strength of the dislocations. By contrast, the increased density of bubbles to act as sinks in the steels irradiated in H F I R appears to cause m o r e balanced defect partitioning. O u r data suggest that subgrain boundaries, and possibly precipitate particles, may also be important sinks in the system as well.

6. O u r data clearly shows helium-enhanced void formation where the low swelling transient regime could be s h o r t e n e d by as much as 75 to 100 dpa. The levels of cavity swelling observed are small, less than 0.5%. H o w e v e r , microstructural details suggest that void formation is in the earl}' stages of d e v e l o p m e n t and that several mechanisms could easily led to m o r e void nucleation and growth as dose increases. The possibility of h e l i u m - e n h a n c e d void swelling remains a legitimate concern for fusion that higher fluence experiments need to address.

A c k n o w l e d g m e n t s

We would like to thank J. M. Vitek for earlier work on the nickel-doped steels, particularly preparing specimens for the various reactor experiments. We would also like to thank Noble Rouse for excellent preparation of T E M specimens, especialy those that are highly radio- active after H F I R irradiation. We thank D. A. Pedraza and J. M. Vitek at O R N L for reviewing this paper. This research was sponsored bv the Office of Fusion Energy, U.S.

D e p a r t m e n t of Energy, under Contract N O . DE-AC05-84OR214()0 with Martin Marietta Energy Systems, Inc.

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60 EFFECTS OF RADIATION ON MATERIALS

[40] Horton, L. L. and Mansur, L. K.. Effects of Irradiation on Materials: Twelfth International Sym- posium, ASTM STP 870, F. A. Garner and J. S. Perrin, Eds.. American Society for Testing and

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Koichiro Hide, t N a o t o Sekimura, ~- Kofi Fukuya, 3 H i d e o Kusanagi, t M a s a f u m i Taguchi, 4 Tatsuo Satake, 4 Yoshio Arai, 4 Masashi I i m u n a , 3 Hiroshi Takaku, ~ and Siori Ishino 4

Microstructural Change in Ferritic Steels Under Heavy Ion Irradiation

REFERENCE: Hide, K., Sekimura, N., Fukuya, K.. Kusanagi, H., Taguchi. M., Satake, T., Arai, Y., Iimura, M., Takaku, H., and lshino, S., "Microstructural Change in Ferritic Steels Under Heavy Ion Irradiation," Effects o f Radiation on Materials: 14th International Sympo- sium, Volume L A S T M STP 1046. N. H. Packan, R. E. Stoller. and A. S. Kumar, Eds., American Society for Testing and Materials, Philadelphia, 1989, pp. 61-72.

ABSTRACT: Ferritic steel is one of the most promising materials for fuel cladding of fast breeder reactors and fusion first wall. However, not enough data on microstructural evolution under irradiation have been obtained, especially at high doses. In the present study, microstructural changes in several ferritic steels under ion irradiation up to 200 dpa were investigated.

The specimens were irradiated with 200 keV C* or 3 MeV Ni z" ions between 698 and 898 K. Post irradiation microstructures were observed with transmission electron microscopes.

Few voids were detected in MA957. Dislocation structure and dispersed oxide particles (Y_,O3) were found to be stable under irradiation up to 200 dpa. Uniform distributions of fine voids were observed in HT-9, Fe-12Cr-2Mo, and pure Fe-12Cr alloy irradiated with 200 keV C"

ions to 150 dpa. However, 3 MeV Ni:- irradiation of HT-9 between 748 and 848 K caused heterogeneous nucleation of voids near lath boundaries. As-received metallic compound- strengthened steel consisted of martensitic and ferritic phases. Large voids were formed in limited region in martensitic phase during ion irradiation, whereas few voids were found in the ferritic phase. Radiation-enhanced formation of Laves phase (Fe~Mo) was observed in ferritic phase.

KEY WORDS: ferritic steels, radiation damage, microstructural change, voids, ion irradiation, precipitates, M.,~C~, Laves phase

Ferritic steel is one of the attractive candidates for fuel cladding in fast b r e e d e r reactors and first walls in fusion reactors because of its excellent swelling resistance under irradiation.

Various types of ferritic steels with increased strength have been under d e v e l o p m e n t because their high t e m p e r a t u r e strength is not as great as that of austenitic steels. Proposed m e t h o d s to improve high t e m p e r a t u r e strength include (1) the introduction of martensitic phase with carbides or nitrides of vanadium, niobium, and tungsten, or both, (2) the dispersion of fine oxide particles in the matrix; and (3) the precipitation of intermetallic c o m p o u n d s during Central Research Institute of the Electric Power Industry. 2-11-1 Iwato-kita, Komae-shi, Tokyo 201, Japan.

-" Lecturer, Engineering Research Institute, University of Tokyo, 2-11-16 Yayoi, Bunkyo-ku, Tokyo 113, Japan.

3 Nuclear Engineering Laboratory, Toshiba Corporation, 8 Shin-sugita, Isogo-ku, Yokohama 235, Japan.

4 Department of Nuclear Engineering, University of Tokyo, 7-3-1 Hongo, Bunkyo-ku, Tokyo 113, Japan.

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Microstructural Precipitation Effects--Gelles and Thomas [3] suggested that the fine