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martensitic and austenitic steels
Journal of Nuclear Materials 459 (2015) 259–264
Contents lists available at ScienceDirect
Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Gas porosity evolution and ion-implanted helium behavior in reactor ferritic/martensitic and austenitic steels I.I. Chernov a, B.A. Kalin a, M.S. Staltsov a,⇑, Kyi Zin Oo a, S.Yu. Binyukova a, O.S. Staltsova a, A.A. Polyansky a, V.S. Ageev b, A.A. Nikitina b a b
National Research Nuclear University ‘‘Moscow Engineering Physics Institute’’, Moscow, Russia A.A. Bochvar High-Technology Research Institute of Inorganic Materials, Moscow, Russia
a r t i c l e
i n f o
Article history: Received 16 November 2011 Accepted 27 September 2014 Available online 30 January 2015
a b s t r a c t The peculiarities of gas porosity formation and helium retention and release in reactor ferritic/martensitic EP-450 and EP-450-ODS and austenitic ChS-68 steels are investigated by transmission electron microscopy and helium thermal desorption spectrometry (HTDS). The samples were irradiated by 40 keV He+ ions up to a fluence of 5 1020 m2 at 293 and 923 K. An nonuniform distribution of helium bubbles and high-level gas swelling in ferritic/martensitic steels were found at high-temperature helium implantation. The same irradiation conditions result in formation of uniformly distributed helium bubbles and low-level swelling in ChS-68 steel. Temperature range of helium release from EP-450-ODS steel was considerably wider in comparison to HTDS-spectra of the EP-450 steel. A considerable quantity of helium is released from ODS steel in the high-temperature range after the main peak of the HTDSspectrum. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Until recently, the development of structural steels for the core of fast reactors and possibly use in fusion reactors focused on the creation or improvement of austenitic steels. The development of these steels centered on increasing their nickel content and increasing their high-temperature strength and radiation resistance through modified composition. In recent decades, in Russia and other countries in the word, intensive research and development is being carried out in other direction, the development of chromium-based steels with a BCC lattice, because the reduced swelling in such steels compared to that of the austenitic steels with a FCC lattice [1]. However, the main disadvantage of chromium steels is their low high-temperature strength compared to that of austenitic steels. In this regard, increased attention is focused on martensitic and ferritic/martensitic ODS steels as promising materials for fast reactor claddings and the fusion reactor first wall (Table 1). In the structural materials of fast reactor cores (at high burnup of nuclear fuel, 15–20% and more of heavy atoms) and, especially, of the first wall and other components of the fusion reactor discharge chamber, significant amounts of hydrogen and helium will accumulated which have a significant influence on the radiation damage and can lead to the reduced service life of the reactor ⇑ Corresponding author. http://dx.doi.org/10.1016/j.jnucmat.2014.09.086 0022-3115/Ó 2015 Elsevier B.V. All rights reserved.
structural elements [1,2]. However, the radiation resistance of new steels and the behavior of their gaseous impurities have not been sufficiently studied [2]. The purpose of this work was to identify the regularities of trapping, retention, and release of helium in reactor ChS-68 austenitic steel and EP-450 and EP-450-ODS ferritic/martensitic steels by helium thermal desorption spectrometry (HTDS) and transmission electron microscopy (TEM).
2. Experimental procedure Table 2 shows the compositions of the investigated steels and two model alloys that are the base of Russian ferritic and austenitic steels. The samples were implanted by 40 keV He+ ions to a fluence of 51020 m2 at 293 (low-temperature) and 923 K (high-temperature). Ferritic/martensitic steels were normalized followed hightemperature tempering; the austenitic steel was quenched. The gas release activation energy was calculated using the peak temperature shift from changing the rate of uniform heating a1 = 1.2 K/s and a2 = 3 K/s [3,4]:
EðeVÞ ¼ kðT m1 T m2 Þ=ðT m2 T m1 Þ ln½a2 =a1 ðT m1 =T m2 Þ2 ;
ð1Þ
where k is the Boltzmann constant and Tm1 and Tm2 are the peak temperatures for heating rates a1 and a2, respectively.
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– HTDS spectra of 13Cr model alloy and EP-450 steel irradiated by He+ ions at room temperature are similar (see Figs. 2 and 3). At that, the alloy 13Cr has a low-temperature stage (970–1020 K) of helium release. As to the steel EP-450, a small quantity of helium releases at temperatures (1220–1270 K) higher than that of main peak (1110 K); – at the irradiation temperature 923 K, the main peaks of helium release appear at a lower temperature than for samples irradiated at room temperature; – the calculated effective activation energy of gas release in the main HTDS peak for the alloy 13Cr was 2.4 eV, and for ferritic/ martensitic steels it was about 3 eV for all the samples irradiated both at 293 and 923 K (see Table 3).
3. Results and discussion Only primary carbonitrides of the type M(C, N) were observed in the austenized steel ChS-68. The structure of EP-450 was complicated with many secondary precipitates (Fig. 1(a)). TEM and X-ray studies showed the presence of type M23C6 carbides in the steel structure which precipitated and grew mainly along the grain boundaries. In addition, lines of type A2B intermetallic compound (Laves phase) and lines of the type M23C6 carbide were observed. In EP-450-ODS, tiny objects were observed (shown by the arrows in Fig. 1(b)) and found to be dispersed nanosized strengthening particles of yttrium oxide (Y2O3). The structures of ferritic/martensitic steels irradiated by He+ ions prior to HTDS-studies and the typical spectra of helium release at uniform heating are shown in Figs. 2–6. The temperature of the main peaks in the HTDS spectra and the effective activation energies of gas release calculated by formula (1) are presented in Table 3. Bubbles did not form in any of the samples irradiated at room temperature; however, a cluster-loop structure was observed. High-temperature irradiation of ferritic/martensitic steels resulted in formation of bubbles (Figs. 5(a) and 6(a)), the distribution of which is very nonuniform both in size and in the bulk of the material. Though large faceted bubbles were observed in local volumes, some zones were practically free of them. In addition, some interlayers with very small bubbles (shown by the arrows in Fig. 5(a)) were observed. Gas swelling in grains with large bubbles reached 10% at a mean swelling of about 3%. As for other materials [5], the HTDS spectra contain the main peak of gas release and several additional peaks in the low- and high-temperature areas. The location of these peaks on the temperature scale depends on irradiation temperature and the rate of uniform post-radiation heating. The features of thermal desorption parameters (peaks temperatures, gas release activation energy) of ferritic/martensitic steels include the following:
In comparison with the HTDS-spectra of 13Cr alloy and EP-450 steel, the spectra of ODS steel have their own peculiarities. First, the release of helium takes place in a broader temperature interval both at room temperature of helium implantation (see Fig. 4) and, especially, at 923 K (see Fig. 6(b)). Second, additional sufficiently intensive peaks, besides the main peak of helium release, were observed at temperatures higher than the temperature of the main maximum. These peaks have a large amount of released helium, especially when it is implanted at 923 K (see Fig. 6(b)). Figs. 7–9(b) show the typical spectra of gas release from the austenitic materials. As a whole, as seen from Table 3, the HTDS peak temperatures and the gas release activation energies of the austenitic materials are higher than those of the ferritic materials. Unlike the ferritic/martensitic steels, high-temperature irradiation of austenitic steel resulted in formation of small bubbles uniformly distributed in the irradiated volume (see Fig. 9(a)). Gas swelling of ChS-68 steel was about 0.3%. The HTDS spectra of the 16Cr15Ni model alloy and the ChS-68 steel are similar at low-temperature irradiation. At that, a lowtemperature stage of helium release was observed in the range of 970–1170 K and the main maximum at about 1270 K, after which the release of helium was practically absent (see Figs. 7 and 8). At
Table 1 Chemical composition of dispersion-strengthened high-temperature chromium steels [2]. Base of the steel
9Cr-ODS
12Cr-ODS
20Cr-ODS
Content of the main alloying elements and the strengthening phase, wt.% C
Cr
W
Ti
Y2O3
Others
Country
0.13–0.14 0.08–0.12 0.14 0.02–0.06 0.01–0.02 0.02–0.04 0.02 0.02 0.011 0.007 0.05–0.10 0.10 0.02
8.5–9.1 8.5–9.5 9.0 11.72–12.8 13.9–14.2 11.97–13.64 13–14 13.1 13.5 13.4 16.0–22.1 19.02 20
1.91–1.97 1.8–2.2 1.92 1.92–2.75 – 1.65–2.02 3.0 – – 2.1 0.29–1.83 1.85 –
0.18–0.29 60.1 0.2 0.3–0.74 1.0–1.07 0.28–0.30 0.4–1.0 2.2 1.3 1.1 0.27–0.28 0.28 0.5
0.34–0.44 0.20–0.27 0.36 0.24–0.46 0.22–0.27 0.24–0.38 0.25–3.0 0.38 0.39 0.39 0.36–0.37 0.37 –
(0.27–0.23)Y – 0.28Y – 0.3Mo (0.19–0.30)Y – 1.4Mo 0.6Mo; 0.17Al 0.2Mo (0.28–0.29)Y 0.29Y; 1.75Si 5.5Al
Japan Europe Europe Japan Japan Japan USA Russia Russia Russia Japan Japan USA
Table 2 Chemical composition of the investigated materials. Material
ChS-68 EP-450 EP-450-ODS Model austenitic alloy 16Cr15Ni Model ferritic alloy 13Cr
Content of elements, wt.% C
Cr
Ni
Mn
Mo
Nb
Ti
V
B
Others
0.06 0.10–0.15 0.10–0.15 – –
16.3 12–14 12–14 16 13
14.8 60.3 60.3 15 –
1.6 – – – –
2.2 1.2–1.8 1.2–1.8 – –
0.2–0.4 0.2–0.4 – –
0.35 – – – –
0.2 0.1–0.3 0.1–0.3 – –
0.004 0.004 0.004 – –
0.5Si – 0.5Y2O3 – –
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Fig. 1. A microstructure of the ferritic/martensitic (a) EP-450 and (b) EP-450-ODS steels.
Fig. 2. A HTDS-spectrum of the model ferritic 13Cr alloy irradiated by He+ at 293 K at a heating rate a2 = 3 K/s.
Fig. 3. A HTDS-spectrum of EP-450 steel irradiated by He+ ions at 293 K at a heating rate a1 = 1.2 K/s (b).
high-temperature helium implantation, as in ferritic/martensitic steels, gas release occurred over a wider temperature range with no high-temperature stage of helium release (see Fig. 9(b)). The main parameters determining helium behavior in these materials are the energy of formation, migration, bounding, and dissociation of gas atoms in various positions of the crystal lattice. It is difficult to experimentally determine these parameters due to the extremely low solubility of helium in metals. If the concentrations of vacancies and helium are high enough (in practice, when fluence F P 1020 m2 at ion implantation), then helium can be trapped in pure metals at its low-temperature implantation with formation of the type HemVn complexes (He is a helium atom, V is a vacancy). Experimental and theoretical investigations have
Fig. 4. A HTDS-spectrum of EP-450-ODS steel irradiated by He+ ions at 293 K for a heating rate a2 = 3 K/s.
shown that type HemVn complexes can be resistant to temperatures of (0.4–0.6) Tmelting and centers for bubble nucleation in metals at heating [6–10]. In addition, it has been shown in [9–11] that, in solid solution alloys containing carbon and/or substitutional elements M, formation of more complicated complexes of the types HemCkVn, HemMkVn, and others, along with the type HemVn simple complexes, are possible. The complicated complexes demonstrate higher thermal stability than type HemVn complexes [10]. It has been shown previously [3–5,12] that the gas release in the main HTDS peak was realized by migration, coalescence during the migration, and intersection of the sample surface by growing bubbles with formation of a ‘‘pin-hole’’ structure. The activation energy of surface self-diffusion (0.65 eV in a-Fe, 0.68 eV in c-Fe and Ni [13]) is much less than that of the volume self-diffusion (2.5 eV in a-Fe, 2.6–2.9 eV in c-Fe and Ni [9]). Thus, the activation energies of gas release from steels for the main peaks in the HTDS spectra presented in Table 3 testify to the fact that the migration of bubbles occurs by the mechanism of volume diffusion, while in model alloys there is a contribution of the surface diffusion as can be seen by lower values of activation energies. The information above applies to low-temperature ion implantation of helium when no bubbles are formed and helium can be located in the matrix in the form of various above-mentioned complexes. At high-temperature irradiation by He+ ions, the initial structure before the HTDS investigations is a microstructure of small helium bubbles (see Figs. 5(a), 6(a) and 9(a)). In this regard, lower temperatures of peaks (see Table 3) under irradiation at 923 K, which are responsible for the mass release of bubbles from the body of grains in the main maximum of the HTDS spectra, are
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Fig. 5. A (a) typical microstructure of EP-450 steel irradiated by He+ at 923 K and (b) a corresponding spectrum of gas release at a heating rate a1 = 1.2 K/s.
Fig. 6. A (a) typical microstructure of EP-450-ODS steel irradiated by He+ ions at 923 K and (b) a corresponding spectrum of gas release for a heating rate a1 = 1.2 K/s.
Table 3 Temperatures of the main maximums T1 and T2 of the HTDS-spectra at rates of uniform heating a1 = 1.2 and a2 = 3 K/s and effective activation energies of gas release from investigated materials. Material
Tirradiation, K
T1, K
T2, K
E, eV
13Cr EP-450
293 293 923 293 923 293 293 923
1075 1107 1070 1070 1055 1273 1294 1226
1111 1117 1087 1080 1072 1314 1327 1253
2.4 ± 0.2 3.1 ± 0.3 3.0 ± 0.3 3.0 ± 0.3 3.0 ± 0.3 3.0 ± 0.3 3.9 ± 0.4 4.2 ± 0.4
EP-450-ODS 16Cr15Ni ChS-68
Note: Temperature measurement accuracy is T = 10 °C; table shows the average values of T1 b T2 obtained from several experiments.
Fig. 7. A HTDS-spectrum of the model austenitic 16Cr15Ni alloy irradiated by He+ at 293 K at a heating rate a1 = 1.2 K/s.
Fig. 8. A HTDS-spectrum of ChS-68 steel irradiated by He+ ions at 293 K for a heating rate a2 = 3 K/s.
due to the fact that time is not required for dissociation of complexes and formation of bubbles. The regularities of thermal helium desorption from three steels implanted by helium 9Cr and 9Cr-ODS (martensitic steels), and 12Cr-ODS (ferritic/martensitic steel) were studied in [14]. It was shown that the part of helium released in the 9Cr-ODS steel via the bubble migration mechanism was less than in the 9Cr steel. The authors concluded that bubble formation was suppressed in the ODS steels. However, the experimental results obtained in our work (see Figs. 4 and 6(b)) show that the HTDS spectra of the ODS steels contain at least one additional, sufficiently strong, maximum in the high-temperature area after the main peak of gas release. This maximum is presumed to be associated with the accumulation of helium near the oxide particles in the process of uniform heating, because the dispersed Y2O3 particles are incoherent (at least semi-coherent [15,16]) and the helium atoms gather at
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Fig. 9. A (a) typical microstructure of ChS-68 steel irradiated by He+ at 923 K and (b) a corresponding spectrum of gas release at a heating rate a1 = 1.2 K/s.
Fig. 10. The equilibrium of the bubble on a hard spherical particle (a) and dependence of the ratio of the radii of the bubble before (r1) and after (r2) contact with the particle on contact angle with the direct capture of a particle migrating bubble (b) [17].
the ‘‘particle–matrix’’ interfaces and form bubbles [17]. The incoherent particle can directly capture the migrating bubble and thus the bubble loses its mobility [18]. The binding energy (i.e. holding force of a particle which depends on the bubble radius and the contact angle h) of such helium bubbles with incoherent particles is high [17], a large detaching force is necessary to tear the bubbles off the particle (i.e. heating to a high temperature, which can be observed in the HTDS spectra of dispersion-strengthened steel). If we take a number of simplifications, in particular an assumption of isotropic surface properties (i.e., the particles are hard and the gas in the bubble is ideal), then the equilibrium of the bubble and the particle is determined by the equality of the surface tensions (Fig. 10(a)). In this case, the surface energies of the matrix material cm, the inclusions cp, and the boundary energy at the particle–matrix interface cp-m must be balanced [17]:
cp ¼ cpm þ cm cos h;
ð2Þ
where h is an equilibrium contact angle. Using the ideal gas laws, a relationship can be found between the spherical bubble radius (r1) in the matrix (e.g., before the trapping of the bubble by a particle) and the radius (r2) of the bubble with a deviation from the spherical form and located at the particle boundary. Consideration of a change in the surface energy, interface energy, and in the free energy of a gas in the bubble shows that the surface energy change in the new equilibrium configuration is compensated by an opposite change in the interface energy. Thus, a change of the gas free energy takes place, which decreases as a result of the expansion:
DE ¼ NkT lnðV 2 =V 1 Þ;
ð3Þ
where V1 and V2 are volumes of the gas in the spherical bubble in the matrix and after its contact with the particle, respectively.
As a result of this energy decrease, the configuration with a bubble on the surface of a particle is more favorable than the configuration with a bubble remote from the particle surface. At that, the energy change DE is the binding energy Eb of the bubble with the particle [12]: Eb = DE. Therefore, using NkT = p1V1, p1 = 2cm/r1 and V2/V1 = r2/r1, we can write
Eb ¼ ð8=3Þpr 21 cm lnðr 2 =r 1 Þ
ð4Þ
The ratio of bubble radii before and after meeting with the particle surface depends on the angle h. The dependence of the value ln(r2/r1) on the contact angle h is shown in Fig. 10(b), from which it can be observed that Eb increases very rapidly with the angle h and should therefore be higher in case of large tension of the interface. In other words, the reason for decreased helium released in the main peak on the mechanism of bubble migration in such steels is not suppression of bubbles formation, as suggested in [14], but the pumping of a sufficiently large part of implanted helium into hightemperature peaks (i.e., into bubbles growing on the strengthening particle–matrix boundaries). It should be noted that the high-temperature peaks in the HTDS spectrum are also present in the EP-450 steel (see Figs. 3 and 5(b)), but their intensity is much lower than that of the ODS steel (see Figs. 4 and 6(b)). The appearance of such peaks in EP-450 steel can be attributed to the formation of helium bubbles on the precipitate–matrix boundaries with a sufficiently high binding energy of bubbles with the secondary phase particles (see Fig. 1(a)); but, however this binding energy seems less than that of bubbles with incoherent particles of Y2O3. These considerations regarding the nature of high-temperature peaks are indirectly confirmed, for example, by data from [9,19]: in the Ni–7.5% Al alloy irradiated by He+ ions at room temperature, long post-irradiation annealing at 750 °C prior to HTDS-investigation led
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to the formation of incoherent particles (c0 -Ni3Al phase) with large bubbles on their boundaries. In addition a distinct thermal desorption peak appeared at a temperature above the main peak temperature. A similar high-temperature HTDS-peak also appeared in alloys of the system V–Ti, in which nitrides and, perhaps oxides result in additional HTDS-peaks [20]. In spite of the fact that the incoherent particles in Ni–7.5% Al alloy, vanadium alloys, and ODS steels are c0 -Ni3Al (an intermetallic compound), nitrides, and Y2O3, respectively the accumulation of helium with formation at the boundaries of bubbles with a high binding energy with particles seems to occur identically. This can be observed by similar HTDS spectra for the alloy Ni–7.5% Al [9,19], vanadium alloys [20] and, in this work, for the ferritic/martensitic EP-450-ODS steel.
4. Summary Based on the experimental results obtained on the behavior of ion-implanted helium in reactor steels, the following conclusions can be made: 1. The temperature range of helium release from the EP-450-ODS steel is much wider, both at room temperature of helium implantation and, especially, at 923 K in comparison to the HTDS-spectra of EP-450 steel. 2. A significant quantity of helium in ODS steel can be retained in bubbles formed at the incoherent particle–matrix interfaces. At that, the release of helium takes place at temperatures of about 200 K higher than that from the matrix. 3. The HTDS-spectrum of the austenitic ChS-68 steel implanted by helium at 923 K is also much wider than after implantation of helium at room temperature, but the high-temperature stage of helium release is absent in contrast to ferritic/martensitic EP-450 and EP-450-ODS steels.
The work was carried out within the analytical departmental target program ‘‘Development of Scientific Potential of Higher Education (2009–2011)’’. References [1] B.A. Kalin, P.A. Platonov, I.I. Chernov, Ya.I. Shtrombah, Physical Materials Science, vol. 6, Pt. 1: Structural Materials for Nuclear Technology, MEPhI, Moscow, 2008 (in Russian). [2] Phoc. of the 1013th Intern. Conf. on Fusion Reactor Materials, J. Nucl. Mater. 307311 (2002), 329333 (2004), 367370 (2007), 386388 (2009), 417 (2011). [3] V.F. Zelensky, I.M. Nekludov, V.V. Ruzhitsky, J. Nucl. Mater. 151 (1987) 22. [4] I.I. Chernov, B.A. Kalin, Behavior of Helium in Fission and Fusion Structural Materials, MEPhI, Moscow, 2008 (in Russian). [5] S.Yu. Binyukova, I.I. Chernov, B.A. Kalin, At. Energ. 104 (2008) 13 (in Russian). [6] D.J. Reed, Radiat. Eff. 31 (1977) 129. [7] J. Evans, Radiat. Eff. 77 (1983) 105. [8] S.E. Donnelly, Radiat. Eff. 90 (1985) 1. [9] B.A. Kalin, I.I. Chernov, A.N. Kalashnikov, M.N. Esaulov, Probl. Atom. Sci. Tech. Ser.: Radiat. Damage Phys. Radiat. Mater. Sci. 1(65)/2(66) (1997) 53 (in Russian). [10] I.I. Chernov, B.A. Kalin, A.N. Kalashnikov, V.M. Ananin, J. Nucl. Mater. 271&272 (1999) 333. [11] G.J. Van der Kolk, A. Van Veen, L.M. Caspers, Delft. Progr. Rept. Ser.: Phys. Phys. Eng. 4 (1979) 19. [12] I.I. Chernov, M.S. Staltsov, B.A. Kalin, Kyi Zin Oo, Aung Kyaw Zaw, V.I. Statsenko, O.N. Korchagin, O.S. Staltsova, Phys. Chem. Mater. Process. 3 (2012) 22 (in Russian). [13] S.Yu. Davidov, Phys. Solid 41 (1999) 11 (in Russian). [14] A. Kimura, R. Sugano, Y. Matsushita, J. Phys. Chem. Solids 66 (2005) 504. [15] J. Rösler, E. Arzt, Acta Metall. Mater. 38 (1990) 671. [16] M.C. Brandes, L. Kovarik, M.K. Miller, G.S. Daehn, M.J. Mills, Acta Mater. 60 (2012) 1827. [17] D.M. Skorov, A.I. Dashkovsky, V.N. Maskalets, V.K. Hiznij, Surface Energy of the Solid Metallic Phases, Atomizdat, Moscow, 1973 (in Russian). [18] Ya.E. Geguzin, M.A. Krivoglaz, The Motion of Macroscopic Inclusions in Solids, Metallurgita, Moscow, 1971. [19] Kyi Zin Oo, I.I. Chernov, M.S. Staltsov, B.A. Kalin, A.N. Kalashnikov, S.Yu. Binyukova, Aung Kyaw Zaw, V.S. Ageev, A.A. Nikitina, At. Energ. 110 (2011) 130 (in Russian). [20] M.S. Staltsov, I.I. Chernov, B.A. Kalin, Kyi Zin Oo, A.A. Polyansky, O.S. Staltsova, Aung Kyaw Zaw, V.M. Chernov, M.M. Potapenko, in: Proc. of the ICFRM-15.
Contents lists available at ScienceDirect
Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Gas porosity evolution and ion-implanted helium behavior in reactor ferritic/martensitic and austenitic steels I.I. Chernov a, B.A. Kalin a, M.S. Staltsov a,⇑, Kyi Zin Oo a, S.Yu. Binyukova a, O.S. Staltsova a, A.A. Polyansky a, V.S. Ageev b, A.A. Nikitina b a b
National Research Nuclear University ‘‘Moscow Engineering Physics Institute’’, Moscow, Russia A.A. Bochvar High-Technology Research Institute of Inorganic Materials, Moscow, Russia
a r t i c l e
i n f o
Article history: Received 16 November 2011 Accepted 27 September 2014 Available online 30 January 2015
a b s t r a c t The peculiarities of gas porosity formation and helium retention and release in reactor ferritic/martensitic EP-450 and EP-450-ODS and austenitic ChS-68 steels are investigated by transmission electron microscopy and helium thermal desorption spectrometry (HTDS). The samples were irradiated by 40 keV He+ ions up to a fluence of 5 1020 m2 at 293 and 923 K. An nonuniform distribution of helium bubbles and high-level gas swelling in ferritic/martensitic steels were found at high-temperature helium implantation. The same irradiation conditions result in formation of uniformly distributed helium bubbles and low-level swelling in ChS-68 steel. Temperature range of helium release from EP-450-ODS steel was considerably wider in comparison to HTDS-spectra of the EP-450 steel. A considerable quantity of helium is released from ODS steel in the high-temperature range after the main peak of the HTDSspectrum. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Until recently, the development of structural steels for the core of fast reactors and possibly use in fusion reactors focused on the creation or improvement of austenitic steels. The development of these steels centered on increasing their nickel content and increasing their high-temperature strength and radiation resistance through modified composition. In recent decades, in Russia and other countries in the word, intensive research and development is being carried out in other direction, the development of chromium-based steels with a BCC lattice, because the reduced swelling in such steels compared to that of the austenitic steels with a FCC lattice [1]. However, the main disadvantage of chromium steels is their low high-temperature strength compared to that of austenitic steels. In this regard, increased attention is focused on martensitic and ferritic/martensitic ODS steels as promising materials for fast reactor claddings and the fusion reactor first wall (Table 1). In the structural materials of fast reactor cores (at high burnup of nuclear fuel, 15–20% and more of heavy atoms) and, especially, of the first wall and other components of the fusion reactor discharge chamber, significant amounts of hydrogen and helium will accumulated which have a significant influence on the radiation damage and can lead to the reduced service life of the reactor ⇑ Corresponding author. http://dx.doi.org/10.1016/j.jnucmat.2014.09.086 0022-3115/Ó 2015 Elsevier B.V. All rights reserved.
structural elements [1,2]. However, the radiation resistance of new steels and the behavior of their gaseous impurities have not been sufficiently studied [2]. The purpose of this work was to identify the regularities of trapping, retention, and release of helium in reactor ChS-68 austenitic steel and EP-450 and EP-450-ODS ferritic/martensitic steels by helium thermal desorption spectrometry (HTDS) and transmission electron microscopy (TEM).
2. Experimental procedure Table 2 shows the compositions of the investigated steels and two model alloys that are the base of Russian ferritic and austenitic steels. The samples were implanted by 40 keV He+ ions to a fluence of 51020 m2 at 293 (low-temperature) and 923 K (high-temperature). Ferritic/martensitic steels were normalized followed hightemperature tempering; the austenitic steel was quenched. The gas release activation energy was calculated using the peak temperature shift from changing the rate of uniform heating a1 = 1.2 K/s and a2 = 3 K/s [3,4]:
EðeVÞ ¼ kðT m1 T m2 Þ=ðT m2 T m1 Þ ln½a2 =a1 ðT m1 =T m2 Þ2 ;
ð1Þ
where k is the Boltzmann constant and Tm1 and Tm2 are the peak temperatures for heating rates a1 and a2, respectively.
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– HTDS spectra of 13Cr model alloy and EP-450 steel irradiated by He+ ions at room temperature are similar (see Figs. 2 and 3). At that, the alloy 13Cr has a low-temperature stage (970–1020 K) of helium release. As to the steel EP-450, a small quantity of helium releases at temperatures (1220–1270 K) higher than that of main peak (1110 K); – at the irradiation temperature 923 K, the main peaks of helium release appear at a lower temperature than for samples irradiated at room temperature; – the calculated effective activation energy of gas release in the main HTDS peak for the alloy 13Cr was 2.4 eV, and for ferritic/ martensitic steels it was about 3 eV for all the samples irradiated both at 293 and 923 K (see Table 3).
3. Results and discussion Only primary carbonitrides of the type M(C, N) were observed in the austenized steel ChS-68. The structure of EP-450 was complicated with many secondary precipitates (Fig. 1(a)). TEM and X-ray studies showed the presence of type M23C6 carbides in the steel structure which precipitated and grew mainly along the grain boundaries. In addition, lines of type A2B intermetallic compound (Laves phase) and lines of the type M23C6 carbide were observed. In EP-450-ODS, tiny objects were observed (shown by the arrows in Fig. 1(b)) and found to be dispersed nanosized strengthening particles of yttrium oxide (Y2O3). The structures of ferritic/martensitic steels irradiated by He+ ions prior to HTDS-studies and the typical spectra of helium release at uniform heating are shown in Figs. 2–6. The temperature of the main peaks in the HTDS spectra and the effective activation energies of gas release calculated by formula (1) are presented in Table 3. Bubbles did not form in any of the samples irradiated at room temperature; however, a cluster-loop structure was observed. High-temperature irradiation of ferritic/martensitic steels resulted in formation of bubbles (Figs. 5(a) and 6(a)), the distribution of which is very nonuniform both in size and in the bulk of the material. Though large faceted bubbles were observed in local volumes, some zones were practically free of them. In addition, some interlayers with very small bubbles (shown by the arrows in Fig. 5(a)) were observed. Gas swelling in grains with large bubbles reached 10% at a mean swelling of about 3%. As for other materials [5], the HTDS spectra contain the main peak of gas release and several additional peaks in the low- and high-temperature areas. The location of these peaks on the temperature scale depends on irradiation temperature and the rate of uniform post-radiation heating. The features of thermal desorption parameters (peaks temperatures, gas release activation energy) of ferritic/martensitic steels include the following:
In comparison with the HTDS-spectra of 13Cr alloy and EP-450 steel, the spectra of ODS steel have their own peculiarities. First, the release of helium takes place in a broader temperature interval both at room temperature of helium implantation (see Fig. 4) and, especially, at 923 K (see Fig. 6(b)). Second, additional sufficiently intensive peaks, besides the main peak of helium release, were observed at temperatures higher than the temperature of the main maximum. These peaks have a large amount of released helium, especially when it is implanted at 923 K (see Fig. 6(b)). Figs. 7–9(b) show the typical spectra of gas release from the austenitic materials. As a whole, as seen from Table 3, the HTDS peak temperatures and the gas release activation energies of the austenitic materials are higher than those of the ferritic materials. Unlike the ferritic/martensitic steels, high-temperature irradiation of austenitic steel resulted in formation of small bubbles uniformly distributed in the irradiated volume (see Fig. 9(a)). Gas swelling of ChS-68 steel was about 0.3%. The HTDS spectra of the 16Cr15Ni model alloy and the ChS-68 steel are similar at low-temperature irradiation. At that, a lowtemperature stage of helium release was observed in the range of 970–1170 K and the main maximum at about 1270 K, after which the release of helium was practically absent (see Figs. 7 and 8). At
Table 1 Chemical composition of dispersion-strengthened high-temperature chromium steels [2]. Base of the steel
9Cr-ODS
12Cr-ODS
20Cr-ODS
Content of the main alloying elements and the strengthening phase, wt.% C
Cr
W
Ti
Y2O3
Others
Country
0.13–0.14 0.08–0.12 0.14 0.02–0.06 0.01–0.02 0.02–0.04 0.02 0.02 0.011 0.007 0.05–0.10 0.10 0.02
8.5–9.1 8.5–9.5 9.0 11.72–12.8 13.9–14.2 11.97–13.64 13–14 13.1 13.5 13.4 16.0–22.1 19.02 20
1.91–1.97 1.8–2.2 1.92 1.92–2.75 – 1.65–2.02 3.0 – – 2.1 0.29–1.83 1.85 –
0.18–0.29 60.1 0.2 0.3–0.74 1.0–1.07 0.28–0.30 0.4–1.0 2.2 1.3 1.1 0.27–0.28 0.28 0.5
0.34–0.44 0.20–0.27 0.36 0.24–0.46 0.22–0.27 0.24–0.38 0.25–3.0 0.38 0.39 0.39 0.36–0.37 0.37 –
(0.27–0.23)Y – 0.28Y – 0.3Mo (0.19–0.30)Y – 1.4Mo 0.6Mo; 0.17Al 0.2Mo (0.28–0.29)Y 0.29Y; 1.75Si 5.5Al
Japan Europe Europe Japan Japan Japan USA Russia Russia Russia Japan Japan USA
Table 2 Chemical composition of the investigated materials. Material
ChS-68 EP-450 EP-450-ODS Model austenitic alloy 16Cr15Ni Model ferritic alloy 13Cr
Content of elements, wt.% C
Cr
Ni
Mn
Mo
Nb
Ti
V
B
Others
0.06 0.10–0.15 0.10–0.15 – –
16.3 12–14 12–14 16 13
14.8 60.3 60.3 15 –
1.6 – – – –
2.2 1.2–1.8 1.2–1.8 – –
0.2–0.4 0.2–0.4 – –
0.35 – – – –
0.2 0.1–0.3 0.1–0.3 – –
0.004 0.004 0.004 – –
0.5Si – 0.5Y2O3 – –
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Fig. 1. A microstructure of the ferritic/martensitic (a) EP-450 and (b) EP-450-ODS steels.
Fig. 2. A HTDS-spectrum of the model ferritic 13Cr alloy irradiated by He+ at 293 K at a heating rate a2 = 3 K/s.
Fig. 3. A HTDS-spectrum of EP-450 steel irradiated by He+ ions at 293 K at a heating rate a1 = 1.2 K/s (b).
high-temperature helium implantation, as in ferritic/martensitic steels, gas release occurred over a wider temperature range with no high-temperature stage of helium release (see Fig. 9(b)). The main parameters determining helium behavior in these materials are the energy of formation, migration, bounding, and dissociation of gas atoms in various positions of the crystal lattice. It is difficult to experimentally determine these parameters due to the extremely low solubility of helium in metals. If the concentrations of vacancies and helium are high enough (in practice, when fluence F P 1020 m2 at ion implantation), then helium can be trapped in pure metals at its low-temperature implantation with formation of the type HemVn complexes (He is a helium atom, V is a vacancy). Experimental and theoretical investigations have
Fig. 4. A HTDS-spectrum of EP-450-ODS steel irradiated by He+ ions at 293 K for a heating rate a2 = 3 K/s.
shown that type HemVn complexes can be resistant to temperatures of (0.4–0.6) Tmelting and centers for bubble nucleation in metals at heating [6–10]. In addition, it has been shown in [9–11] that, in solid solution alloys containing carbon and/or substitutional elements M, formation of more complicated complexes of the types HemCkVn, HemMkVn, and others, along with the type HemVn simple complexes, are possible. The complicated complexes demonstrate higher thermal stability than type HemVn complexes [10]. It has been shown previously [3–5,12] that the gas release in the main HTDS peak was realized by migration, coalescence during the migration, and intersection of the sample surface by growing bubbles with formation of a ‘‘pin-hole’’ structure. The activation energy of surface self-diffusion (0.65 eV in a-Fe, 0.68 eV in c-Fe and Ni [13]) is much less than that of the volume self-diffusion (2.5 eV in a-Fe, 2.6–2.9 eV in c-Fe and Ni [9]). Thus, the activation energies of gas release from steels for the main peaks in the HTDS spectra presented in Table 3 testify to the fact that the migration of bubbles occurs by the mechanism of volume diffusion, while in model alloys there is a contribution of the surface diffusion as can be seen by lower values of activation energies. The information above applies to low-temperature ion implantation of helium when no bubbles are formed and helium can be located in the matrix in the form of various above-mentioned complexes. At high-temperature irradiation by He+ ions, the initial structure before the HTDS investigations is a microstructure of small helium bubbles (see Figs. 5(a), 6(a) and 9(a)). In this regard, lower temperatures of peaks (see Table 3) under irradiation at 923 K, which are responsible for the mass release of bubbles from the body of grains in the main maximum of the HTDS spectra, are
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Fig. 5. A (a) typical microstructure of EP-450 steel irradiated by He+ at 923 K and (b) a corresponding spectrum of gas release at a heating rate a1 = 1.2 K/s.
Fig. 6. A (a) typical microstructure of EP-450-ODS steel irradiated by He+ ions at 923 K and (b) a corresponding spectrum of gas release for a heating rate a1 = 1.2 K/s.
Table 3 Temperatures of the main maximums T1 and T2 of the HTDS-spectra at rates of uniform heating a1 = 1.2 and a2 = 3 K/s and effective activation energies of gas release from investigated materials. Material
Tirradiation, K
T1, K
T2, K
E, eV
13Cr EP-450
293 293 923 293 923 293 293 923
1075 1107 1070 1070 1055 1273 1294 1226
1111 1117 1087 1080 1072 1314 1327 1253
2.4 ± 0.2 3.1 ± 0.3 3.0 ± 0.3 3.0 ± 0.3 3.0 ± 0.3 3.0 ± 0.3 3.9 ± 0.4 4.2 ± 0.4
EP-450-ODS 16Cr15Ni ChS-68
Note: Temperature measurement accuracy is T = 10 °C; table shows the average values of T1 b T2 obtained from several experiments.
Fig. 7. A HTDS-spectrum of the model austenitic 16Cr15Ni alloy irradiated by He+ at 293 K at a heating rate a1 = 1.2 K/s.
Fig. 8. A HTDS-spectrum of ChS-68 steel irradiated by He+ ions at 293 K for a heating rate a2 = 3 K/s.
due to the fact that time is not required for dissociation of complexes and formation of bubbles. The regularities of thermal helium desorption from three steels implanted by helium 9Cr and 9Cr-ODS (martensitic steels), and 12Cr-ODS (ferritic/martensitic steel) were studied in [14]. It was shown that the part of helium released in the 9Cr-ODS steel via the bubble migration mechanism was less than in the 9Cr steel. The authors concluded that bubble formation was suppressed in the ODS steels. However, the experimental results obtained in our work (see Figs. 4 and 6(b)) show that the HTDS spectra of the ODS steels contain at least one additional, sufficiently strong, maximum in the high-temperature area after the main peak of gas release. This maximum is presumed to be associated with the accumulation of helium near the oxide particles in the process of uniform heating, because the dispersed Y2O3 particles are incoherent (at least semi-coherent [15,16]) and the helium atoms gather at
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Fig. 9. A (a) typical microstructure of ChS-68 steel irradiated by He+ at 923 K and (b) a corresponding spectrum of gas release at a heating rate a1 = 1.2 K/s.
Fig. 10. The equilibrium of the bubble on a hard spherical particle (a) and dependence of the ratio of the radii of the bubble before (r1) and after (r2) contact with the particle on contact angle with the direct capture of a particle migrating bubble (b) [17].
the ‘‘particle–matrix’’ interfaces and form bubbles [17]. The incoherent particle can directly capture the migrating bubble and thus the bubble loses its mobility [18]. The binding energy (i.e. holding force of a particle which depends on the bubble radius and the contact angle h) of such helium bubbles with incoherent particles is high [17], a large detaching force is necessary to tear the bubbles off the particle (i.e. heating to a high temperature, which can be observed in the HTDS spectra of dispersion-strengthened steel). If we take a number of simplifications, in particular an assumption of isotropic surface properties (i.e., the particles are hard and the gas in the bubble is ideal), then the equilibrium of the bubble and the particle is determined by the equality of the surface tensions (Fig. 10(a)). In this case, the surface energies of the matrix material cm, the inclusions cp, and the boundary energy at the particle–matrix interface cp-m must be balanced [17]:
cp ¼ cpm þ cm cos h;
ð2Þ
where h is an equilibrium contact angle. Using the ideal gas laws, a relationship can be found between the spherical bubble radius (r1) in the matrix (e.g., before the trapping of the bubble by a particle) and the radius (r2) of the bubble with a deviation from the spherical form and located at the particle boundary. Consideration of a change in the surface energy, interface energy, and in the free energy of a gas in the bubble shows that the surface energy change in the new equilibrium configuration is compensated by an opposite change in the interface energy. Thus, a change of the gas free energy takes place, which decreases as a result of the expansion:
DE ¼ NkT lnðV 2 =V 1 Þ;
ð3Þ
where V1 and V2 are volumes of the gas in the spherical bubble in the matrix and after its contact with the particle, respectively.
As a result of this energy decrease, the configuration with a bubble on the surface of a particle is more favorable than the configuration with a bubble remote from the particle surface. At that, the energy change DE is the binding energy Eb of the bubble with the particle [12]: Eb = DE. Therefore, using NkT = p1V1, p1 = 2cm/r1 and V2/V1 = r2/r1, we can write
Eb ¼ ð8=3Þpr 21 cm lnðr 2 =r 1 Þ
ð4Þ
The ratio of bubble radii before and after meeting with the particle surface depends on the angle h. The dependence of the value ln(r2/r1) on the contact angle h is shown in Fig. 10(b), from which it can be observed that Eb increases very rapidly with the angle h and should therefore be higher in case of large tension of the interface. In other words, the reason for decreased helium released in the main peak on the mechanism of bubble migration in such steels is not suppression of bubbles formation, as suggested in [14], but the pumping of a sufficiently large part of implanted helium into hightemperature peaks (i.e., into bubbles growing on the strengthening particle–matrix boundaries). It should be noted that the high-temperature peaks in the HTDS spectrum are also present in the EP-450 steel (see Figs. 3 and 5(b)), but their intensity is much lower than that of the ODS steel (see Figs. 4 and 6(b)). The appearance of such peaks in EP-450 steel can be attributed to the formation of helium bubbles on the precipitate–matrix boundaries with a sufficiently high binding energy of bubbles with the secondary phase particles (see Fig. 1(a)); but, however this binding energy seems less than that of bubbles with incoherent particles of Y2O3. These considerations regarding the nature of high-temperature peaks are indirectly confirmed, for example, by data from [9,19]: in the Ni–7.5% Al alloy irradiated by He+ ions at room temperature, long post-irradiation annealing at 750 °C prior to HTDS-investigation led
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to the formation of incoherent particles (c0 -Ni3Al phase) with large bubbles on their boundaries. In addition a distinct thermal desorption peak appeared at a temperature above the main peak temperature. A similar high-temperature HTDS-peak also appeared in alloys of the system V–Ti, in which nitrides and, perhaps oxides result in additional HTDS-peaks [20]. In spite of the fact that the incoherent particles in Ni–7.5% Al alloy, vanadium alloys, and ODS steels are c0 -Ni3Al (an intermetallic compound), nitrides, and Y2O3, respectively the accumulation of helium with formation at the boundaries of bubbles with a high binding energy with particles seems to occur identically. This can be observed by similar HTDS spectra for the alloy Ni–7.5% Al [9,19], vanadium alloys [20] and, in this work, for the ferritic/martensitic EP-450-ODS steel.
4. Summary Based on the experimental results obtained on the behavior of ion-implanted helium in reactor steels, the following conclusions can be made: 1. The temperature range of helium release from the EP-450-ODS steel is much wider, both at room temperature of helium implantation and, especially, at 923 K in comparison to the HTDS-spectra of EP-450 steel. 2. A significant quantity of helium in ODS steel can be retained in bubbles formed at the incoherent particle–matrix interfaces. At that, the release of helium takes place at temperatures of about 200 K higher than that from the matrix. 3. The HTDS-spectrum of the austenitic ChS-68 steel implanted by helium at 923 K is also much wider than after implantation of helium at room temperature, but the high-temperature stage of helium release is absent in contrast to ferritic/martensitic EP-450 and EP-450-ODS steels.
The work was carried out within the analytical departmental target program ‘‘Development of Scientific Potential of Higher Education (2009–2011)’’. References [1] B.A. Kalin, P.A. Platonov, I.I. Chernov, Ya.I. Shtrombah, Physical Materials Science, vol. 6, Pt. 1: Structural Materials for Nuclear Technology, MEPhI, Moscow, 2008 (in Russian). [2] Phoc. of the 1013th Intern. Conf. on Fusion Reactor Materials, J. Nucl. Mater. 307311 (2002), 329333 (2004), 367370 (2007), 386388 (2009), 417 (2011). [3] V.F. Zelensky, I.M. Nekludov, V.V. Ruzhitsky, J. Nucl. Mater. 151 (1987) 22. [4] I.I. Chernov, B.A. Kalin, Behavior of Helium in Fission and Fusion Structural Materials, MEPhI, Moscow, 2008 (in Russian). [5] S.Yu. Binyukova, I.I. Chernov, B.A. Kalin, At. Energ. 104 (2008) 13 (in Russian). [6] D.J. Reed, Radiat. Eff. 31 (1977) 129. [7] J. Evans, Radiat. Eff. 77 (1983) 105. [8] S.E. Donnelly, Radiat. Eff. 90 (1985) 1. [9] B.A. Kalin, I.I. Chernov, A.N. Kalashnikov, M.N. Esaulov, Probl. Atom. Sci. Tech. Ser.: Radiat. Damage Phys. Radiat. Mater. Sci. 1(65)/2(66) (1997) 53 (in Russian). [10] I.I. Chernov, B.A. Kalin, A.N. Kalashnikov, V.M. Ananin, J. Nucl. Mater. 271&272 (1999) 333. [11] G.J. Van der Kolk, A. Van Veen, L.M. Caspers, Delft. Progr. Rept. Ser.: Phys. Phys. Eng. 4 (1979) 19. [12] I.I. Chernov, M.S. Staltsov, B.A. Kalin, Kyi Zin Oo, Aung Kyaw Zaw, V.I. Statsenko, O.N. Korchagin, O.S. Staltsova, Phys. Chem. Mater. Process. 3 (2012) 22 (in Russian). [13] S.Yu. Davidov, Phys. Solid 41 (1999) 11 (in Russian). [14] A. Kimura, R. Sugano, Y. Matsushita, J. Phys. Chem. Solids 66 (2005) 504. [15] J. Rösler, E. Arzt, Acta Metall. Mater. 38 (1990) 671. [16] M.C. Brandes, L. Kovarik, M.K. Miller, G.S. Daehn, M.J. Mills, Acta Mater. 60 (2012) 1827. [17] D.M. Skorov, A.I. Dashkovsky, V.N. Maskalets, V.K. Hiznij, Surface Energy of the Solid Metallic Phases, Atomizdat, Moscow, 1973 (in Russian). [18] Ya.E. Geguzin, M.A. Krivoglaz, The Motion of Macroscopic Inclusions in Solids, Metallurgita, Moscow, 1971. [19] Kyi Zin Oo, I.I. Chernov, M.S. Staltsov, B.A. Kalin, A.N. Kalashnikov, S.Yu. Binyukova, Aung Kyaw Zaw, V.S. Ageev, A.A. Nikitina, At. Energ. 110 (2011) 130 (in Russian). [20] M.S. Staltsov, I.I. Chernov, B.A. Kalin, Kyi Zin Oo, A.A. Polyansky, O.S. Staltsova, Aung Kyaw Zaw, V.M. Chernov, M.M. Potapenko, in: Proc. of the ICFRM-15.