In situ TEM annealing of ion-amorphized Hi Nicalon S and Tyranno SA3 SiC fibers

Nuclear Instruments and Methods in Physics Research B 374 (2016) 76–81

Contents lists available at ScienceDirect

Nuclear Instruments and Methods in Physics Research B journal homepage: www.elsevier.com/locate/nimb

In situ TEM annealing of ion-amorphized Hi Nicalon S and Tyranno SA3 SiC fibers J. Huguet-Garcia a,⇑, A. Jankowiak a, S. Miro b, E. Meslin c, Y. Serruys b, J.-M. Costantini a a

CEA, DEN, Service de Recherches Métallurgiques Appliquées, F-91191 Gif-sur-Yvette, France CEA, DEN, Service de Recherches en Métallurgie Physique, F-91191 Gif-sur-Yvette, France c CEA, DEN, Service de Recherches en Métallurgie Physique, Laboratoire JANNUS, F-91191 Gif-sur-Yvette, France b

a r t i c l e

i n f o

Article history: Received 30 August 2015 Received in revised form 30 November 2015 Accepted 16 December 2015 Available online 23 December 2015 Keywords: SiC fibers Ion-amorphization Recrystallization In situ TEM Thermal annealing

a b s t r a c t In this work, recrystallization of ion-amorphized Hi Nicalon Type S and Tyranno SA3 SiC fibers (4 MeV Au3+, 2  1015 cm 2) has been studied via in situ TEM annealing. Both fibers show a two-step recovery process of the radiation damage. First recovery stage starts at temperatures as low as 250 °C and implies recovery of the radiation swelling. Eventually the amorphous layer recrystallizes with no signs of polytype change (3C-SiC). Recrystallization temperatures yield 900–920 °C and 930 °C for the HNS and the TSA3 respectively. HNS fiber shows columnar recrystallization perpendicular to the amorphous–crystalline interphase with a grain growth rate of 20 nm min 1. On the other hand, recrystallization of TSA3 fiber is rather ‘‘spontaneous” with no preferential growth direction. The different recrystallization is attributed to the different microstructure of the fibers. Ó 2015 Elsevier B.V. All rights reserved.

1. Introduction SiC composites reinforced with nuclear grade SiC fibers, such as Hi Nicalon S (HNS) and Tyranno SA3 (TSA3) are promising candidates for future nuclear applications [1–4]. One of the most detrimental features of SiC for its use under nuclear environments at low temperatures is the degradation of its physico-chemical properties due to irradiation induced amorphization. It has been already demonstrated that, regardless of their microstructure, HNS and TSA3 fibers present similar amorphization threshold conditions than 6H-SiC single crystals [5]. Several studies have reported partial or complete recovery of the crystalline structure of amorphized 6H-SiC by means of thermal annealing [6–9]. In this work, ion-amorphized HNS and TSA3 fibers have been annealed and observed in situ via TEM in order to study the role of their microstructure in their recrystallization behavior.

tured by Nippon Carbon, Japan. Both fibers exhibit a density over 97% of the theoretical density and consist in 3C polytype SiC crystals with turbostratic carbon at grain boundaries [10]. Grain sizes differ from one fiber to the other, ranging from 100 to 300 nm for TSA3 fibers and from 20 to 50 nm for HNS fibers [5,10]. 2.2. Irradiation conditions Hi Nicalon type S (HNS) and Tyranno SA 3 (TSA3) SiC fibers have been irradiated at JANNUS-Saclay irradiation facility [11] with 4 MeV Au3+ ions to 2  1015 cm 2 at room temperature (RT). Ion beam angle of incidence is 15° with respect to the fibers axis and fluence-dose conversion has been estimated with SRIM-2013 [12] using full-cascade calculations. Theoretical density was set to 3.21 g cm 3 and threshold lattice displacement energies of 35 and 20 eV for Si and C sublattices, respectively [13]. At the specified fluence, mean dose yields 4 dpa.

2. Experimental conditions 2.3. In situ TEM 2.1. Materials Materials used are Tyranno SA 3 (TSA3) SiC fibers manufactured by UBE Industries Ltd, Japan, and Hi Nicalon type S (HNS) manufac⇑ Corresponding author. Tel.: +33 1 69 08 12 11; fax: +33 1 69 08 71 67. E-mail address: [email protected] (J. Huguet-Garcia). http://dx.doi.org/10.1016/j.nimb.2015.12.032 0168-583X/Ó 2015 Elsevier B.V. All rights reserved.

TEM annealing and observation of both fibers have been conducted in a conventional CM20 TWIN-FEI (Philips, Amsterdam, Netherlands) operated at 200 kV equipped with a LaB6 crystal as electron source and a Gatan Orius 200 CCD camera (Gatan Inc, Warrendale, PA, USA). Also, this microscope is equipped with a heating specimen holder (25–1000 °C) with manual temperature

J. Huguet-Garcia et al. / Nuclear Instruments and Methods in Physics Research B 374 (2016) 76–81

control. Thin foils for in situ TEM annealing were prepared using the focused ion beam (FIB) technique. The cross-section specimens were extracted from the irradiated fibers using a Helios Nanolab 650 (FEI Co, Hillsboro, OR., USA) equipped with an electronic beam and a Ga ionic beam. Specimens were transferred to a Mo grid in order to withstand high temperatures. Recrystallization process

77

was recorded in real time and quantitative data have been obtained from images extracted from the recorded video using ImageJ software [14]. Thermal annealing tests were conducted as follows. TSA3 fiber was heated at high heating rate to 550 °C, then to 810 °C at 16 °C min 1 and finally to 930 °C at 2 °C min 1. The latter temperature was kept constant during 35 min. As no effects

Fig. 1. In situ TEM observation of the recrystallization of an ion-amorphized HNS fiber. Recovery of the amorphous layer starts by the densification of the amorphous layer followed by its columnar recrystallization.

Fig. 2. In situ TEM observation of the recrystallization of an ion-amorphized TSA3 fiber. Recovery of the amorphous layer starts by the densification of the amorphous layer followed by its spontaneous recrystallization.

78

J. Huguet-Garcia et al. / Nuclear Instruments and Methods in Physics Research B 374 (2016) 76–81

were observed in the microscope under 900 °C for the TSA3 fiber, the HNS fiber was heated to 900 °C at 15 °C min 1 after initial heating at high rate to 800 °C. During the latter test a small temperature excursion to 930 °C happened after 10 min at 900 °C after which the temperature was stabilized at 920 °C. The lack of precision of manual control introduces small deviations in the temperature curves.

3. Results and discussion SiC based composites reinforced with third generation SiC fibers, TSA3 or HNS, are among the most promising materials for future nuclear structural applications [1]. As a nuclear material, these composites, and therefore their fiber reinforcement, will be exposed to neutron and fission products irradiation which can induce the degradation of the physico-chemical properties of SiC hence limiting their in-core lifetime. It is well known that at low

Fig. 3. Kinetics of the thermal annealing induced recrystallization on ion-amorphized HNS fiber: at the top, temperature of the sample during the annealing test, in the middle, densification of the irradiated layer prior to recrystallization and, at the bottom, grain length as a function of time during recrystallization.

temperatures, lattice damage caused by irradiation under ballistic regimes builds-up causing the loss of crystalline structure, i.e. its amorphization, with no significant differences between cubic (3C) and hexagonal (6H) SiC polytypes [15,16]. Though amorphized SiC is highly stable due to the low mobility of point defects [17], several studies have reported partial or complete recovery of the crystalline structure of the irradiated layer by means of thermal annealing [6–9]. However, to the knowledge of the authors, there is no information available about the recrystallization phenomena of ion-amorphized SiC fibers. In order to study how the different microstructures of HNS and TSA3 fibers affects recrystallization, both SiC fibers have been irradiated with 4 MeV Au3+ ions at RT to 4 dpa. This dose is up to ten times the dose to ionamorphization of 6H-SiC under the same irradiation conditions [18] forming an amorphous SiC layer of 1.3 lm in the TSA3 fiber and 1.1 in the HNS fiber [5]. Cross sectional thin foils extracted from the fibers were subsequently annealed and observed via in situ TEM. Figs. 1 and 2 show the evolution of the ion-amorphized HNS and TSA3, respectively. As it can be observed, HNS fibers shows a homogeneous layer of amorphous SiC whereas TSA3 has some inclusions embedded in the amorphous layer. These inclusions are attributed to carbonaceous phases resulting from the intergranular free carbon of the original microstructure as the large size of this carbon pockets hinder ballistic mixing under irradiation [5]. The observed recrystallization differs between the HNS and the TSA3 fibers. As shown in Fig. 1, for the HNS fibers, columnar grains emerge directly from the non-irradiated nanocrystalline substrate perpendicularly to the amorphous–crystalline (a–c) interphase and grow with time. On the other hand, as shown in Fig. 2, recrystallization process of the ion-amorphized TSA3 SiC fiber is rather spontaneous and heterogeneous with no apparent preferential

Fig. 4. Kinetics of the thermal annealing induced structural relaxation on the ionamorphized layer of TSA3 fiber during the temperature transient. Densification starts for temperatures as low as 250 °C and saturates once the layer spontaneously recrystallizes.

J. Huguet-Garcia et al. / Nuclear Instruments and Methods in Physics Research B 374 (2016) 76–81

growth direction with respect the a–c interphase. For both cases, neither surface contamination nor degradation was observed during the test and the irradiated layer was successfully fully recrystallized. Similar experiments have been reported in the literature for ion-amorphized 6H-SiC. At variance with the observed recrystallization of the fibers, thermal annealing of ion-amorphized 6H single crystal induces first the growth of the a–c interphase. Once the saturation of the a–c interphase thickness is achieved, around 100 nm, columnar grains start to emerge from the a–c interphase with an orientation of ±70° with respect to it [19]. Both, growth of the a–c interphase and grain orientation, have been attributed to the epitaxial recrystallization of the amorphous SiC (a-SiC) on top of the (0 0 0 1) oriented 6H SiC substrate. In our case, it is likely that the nanophased microstructure of the HNS hinders solid phase epitaxial recrystallization of the amorphous layer. In the TSA3 case, its spontaneous and non-oriented recrystallization is attributed to the presence of the carbonaceous phases embedded in the amorphous layer which may prevent columnar grain growth while facilitating nucleation of SiC grains near the carbonaceous phases. Recrystallization kinetics of the HNS fiber is shown in Fig. 3. The irradiated layer densifies at a rate of 6.5 nm min 1 during the first annealing stage previous to the recrystallization. Once the thickness of the irradiated layer has already reached the saturation value, columnar recrystallization starts with a constant grain growth rate of 21.2 nm min 1. For the TSA3 only the kinetics of the densification of the irradiated layer can be measured due to the spontaneous recrystallization. As it can be noticed in Fig. 4, densification of the irradiated layer starts for temperatures as low as 250 °C. At the beginning, the irradiated layer densifies at rate of 6.8 nm min 1 to then almost saturate at 580 °C until reaching 860 °C. Once at this temperature, densification restarts at a rate of 4.2 nm min 1 until recrystallization. However, as shown in Figs. 3 and 4, densification of the amorphous layers previous to recrystallization were observed during the temperature ramp hence limiting the physical meaning of the obtained densification rates. Even though the observed densification must be regarded from a qualitative point of view, observations described above suggest that annealing of the ion-amorphized fibers induces a two-step recovery process. First recovery stage, densification, would consist in the structural relaxation of the large swelling resulting from the amorphization process, which is reported to be 11.5–25% [18,20].

79

These observations are consistent with reported structural relaxation of a-SiC prior to recrystallization. For instance, Bohn et al. [21] reported a constant decrease of the amorphous layer width with temperature for annealing of N-implanted 6H-SiC (8  1016 cm 2 62 keV N) between RT and 1450 °C. In agreement with this observation, Snead et al. [22] observed linear densification with annealing temperature for neutron amorphized SiC between 150 °C and 885 °C. Also, Höfgen et al. [6] observed recovery in a temperature range in agreement with the one found for the TSA3 fiber, between 250 °C and 700 °C. Both, Bohn et al. [21] and Snead et al. [22] attributed the structural relaxation to defect annealing processes with low activation energies. This argument is supported by Bae et al., [23] who noticed the presence of a more chemically ordered atomistic structure in a-SiC after 1 h annealing at 890 °C. In a more recent paper, Ishimaru et al. [19] provided in situ TEM observations of a-SiC densification due to structural relaxation during thermal annealing of ion-amorphized 6H-SiC (1015 cm 2 10 MeV Au3+) between 300 °C and 800 °C. In addition, they provided a detailed analysis pointing out that annealing of low energy and short Si–Si bonds (2.32 eV/bond, 2.3 Å) is faster than C–C ones (3.68 eV/bond, 1.5 Å). This unbalanced recovery reduces the average bond length justifying the densification of the a-SiC layer prior to recrystallization. Finally, Miro et al. [9] show that lattice damage fraction recovery and densification of the irradiated layer evolve similarly with the annealing temperature with an onset temperature of 200 °C. Once densification reaches saturation, second recovery stage consist in recrystallization of the amorphous layer resulting in rather complex microstructures. The final microstructure and SAED patterns of the HNS substrate and recrystallized layer are shown in Fig. 5. The recrystallized layer is characterized by a polycrystalline microstructure formed by columnar grains with diffuse boundaries with grain sizes substantially larger than the substrate ones. In addition, there is no noticeable intergranular free carbon. Upper SAED pattern corresponds to the recrystallized layer. The presence of doubled and elongated spots is due to the polycrystalline nature of the recrystallized layer. Interplanar distances have been determined as the inverse of the distance in the reciprocal space between the highest intensity point of each spot with the central one and yield d(hkl)1 = 2.54 and d(hkl)2 = 2.57 Å forming an angle of 55°. Lower SAED shows numerous bright spots disposed

Fig. 5. Detail of the final microstructure of the irradiated HNS fiber after the in situ annealing. The final microstructure and the SAED patterns substantially differ from the substrate. It is not noticeable the presence of free carbon in the recrystallized layer.

80

J. Huguet-Garcia et al. / Nuclear Instruments and Methods in Physics Research B 374 (2016) 76–81

Fig. 6. Detail of the final microstructure of the irradiated TSA3 fiber after the in situ annealing. The final microstructure and the SAED patterns substantially differ from the substrate. It is noticeable the presence of free C in the recrystallized layer.

in concentric rings as a consequence of the nanophased microstructure. Radii of these rings yield r(hkl)1 = 2.55 Å, r(hkl)2 = 1.55 Å and r(hkl)3 = 1.33 Å and can be attributed to (1 1 1), (2 2 0) and (3 1 1) planes of 3C-SiC. The final microstructure of the TSA3 recrystallized layer is presented in Fig. 6 where grains are barely distinguishable. Also, it is remarkable the constant presence in the irradiated layer of what has been attributed to carbonaceous phases during the annealing test. Upper SAED shows, as in the previous case, a hexagonal pattern with interplanar distances d(hkl)1 = 2.55 Å and d(hkl)2 = 2.61 Å forming an angle of 56°. Lower SAED shows the pattern related to the non-irradiated microstructure. Though it is not as clear as for the HNS, concentric rings formed by different spots can also be identified. Radii of these rings yield r(hkl)1 = 2.55 Å, r(hkl)2 = 1.53 Å and r(hkl)3 = 1.33 Å. This values are in agreement with those found for the HNS microstructure and are also characteristic of the 3C-SiC polytype. Also, it is noticeable the presence of a diffuse ring near the transmitted beam spot with radius rc = 3.56 Å attributed to the intergranular free carbon [24]. Polytypic transformations have been observed during solid phase epitaxy recrystallization over 6H single crystalline substrates [7,19,25,26] as interplanar spacing of (1 1 1) reflections of the 3C polytype are equivalent to (1 0 1 2) and (0 0 0 6) reflection of the 6H-SiC polytype allowing a smooth transition between (1 1 1)3C and the (0 0 0 1)6H interface [27]. Regarding the recrystallized layers of HNS and TSA3 SiC fibers, SAED pattern indexation is consistent with (1 1 1) reflections of 3C-SiC, in agreement with Osterberg et al. [28] who did not observe recrystallization induced polytypism in a-SiC over polycrystalline 3C-SiC. Indeed, as polytypic transformations have only been observed for solid phase epitaxy over 6H single crystalline substrates, it is believed that the nanostructured substrate below the amorphous layer of the HNS fiber and the spontaneous recrystallization of the TSA3 fiber prevented polytypic transformation.

4. Conclusions In situ thermal annealing experiments have been successfully conducted in ion-amorphized TSA3 and HNS fibers. Recrystallization characterization reveals a two stage recovery process. First stage consists in the densification of the irradiated layer for temperatures between 250 and 900 °C. Second stage concerns

the recrystallization of the irradiated layer with significant differences between both fibers depending on their microstructure. The nanostructured substrate of the HNS yields columnar recrystallization of the homogeneous amorphous layer perpendicular to the a–c interphase with a grain growth rate of 21 nm min 1. On the other hand, carbonaceous phases embedded in the amorphous layer of the TSA3 yield a rather spontaneous recrystallization with no preferred direction with respect the a–c interphase. Finally, in both cases the recrystallized layer shows a polycrystalline microstructure of 3C-SiC grains without preservation of the original microstructure.

Acknowledgements The authors thank the JANNUS staffs for their technical support and B. Arnal for TEM specimen preparation.

References [1] A. Ivekovic´, S. Novak, G. Drazˇic´, D. Blagoeva, S.G. de Vicente, Current status and prospects of SiCf/SiC for fusion structural applications, J. Eur. Ceram. Soc. 33 (2013) 1577–1589. [2] L.L. Snead, T. Nozawa, M. Ferraris, Y. Katoh, R. Shinavski, M. Sawan, Silicon carbide composites as fusion power reactor structural materials, J. Nucl. Mater. 417 (2011) 330–339. [3] L. Hallstadius, S. Johnson, E. Lahoda, Cladding for high performance fuel, Prog. Nucl. Energy 57 (2012) 71–76. [4] P. Yvon, F. Carré, Structural materials challenges for advanced reactor systems, J. Nucl. Mater. 385 (2009) 217–222. [5] J. Huguet-Garcia, A. Jankowiak, S. Miro, D. Gosset, Y. Serruys, J.-M. Costantini, Study of the ion-irradiation behavior of advanced SiC fibers by Raman spectroscopy and transmission electron microscopy, J. Am. Ceram. Soc. 98 (2015) 675–682. [6] A. Höfgen, V. Heera, F. Eichhorn, W. Skorupa, Annealing and recrystallization of amorphous silicon carbide produced by ion implantation, J. Appl. Phys. 84 (1998) 4769. [7] S. Harada, M. Ishimaru, T. Motooka, T. Nakata, T. Yoneda, M. Inoue, Recrystallization of MeV Si implanted 6H-SiC, Appl. Phys. Lett. 69 (1996) 3534. [8] J. Aihara, K. Hojou, S. Furuno, M. Ishihara, K. Hayashi, Amorphization with ion irradiation and recrystallization by annealing of SiC crystals, Nucl. Instr. Meth. Phys. Res. B 166–167 (2000) 379–384. [9] S. Miro, J.-M. Costantini, J. Huguet-Garcia, L. Thomé, Recrystallization of hexagonal silicon carbide after gold ion irradiation and thermal annealing, Philos. Mag. 94 (2014) 3898–3913. [10] C. Sauder, J. Lamon, Tensile creep behavior of SiC-based fibers with a low oxygen content, J. Am. Ceram. Soc. 90 (2007) 1146–1156.

J. Huguet-Garcia et al. / Nuclear Instruments and Methods in Physics Research B 374 (2016) 76–81 [11] L. Beck, Y. Serruys, S. Miro, P. Trocellier, E. Bordas, F. Leprêtre, et al., Ion irradiation and radiation effect characterization at the JANNUS-Saclay triple beam facility, J. Mater. Res. 30 (2015) 1183–1194. [12] J.F. Ziegler, M.D. Ziegler, J.P. Biersack, SRIM – the stopping and range of ions in matter (2010), Nucl. Instr. Meth. Phys. Res. B 268 (2010) 1818–1823. [13] F. Gao, W.J. Weber, R. Devanathan, Defect production, multiple ion–solid interactions and amorphization in SiC, Nucl. Instr. Meth. Phys. Res. B 191 (2002) 487–496. [14] C.A. Schneider, W.S. Rasband, K.W. Eliceiri, NIH Image to ImageJ: 25 years of image analysis, Nat. Methods 9 (2012) 671–675. [15] A. Debelle, L. Thomé, D. Dompoint, A. Boulle, F. Garrido, J. Jagielski, et al., Characterization and modelling of the ion-irradiation induced disorder in 6HSiC and 3C-SiC single crystals, J. Phys. D Appl. Phys. 43 (2010) 455408. [16] L.L. Snead, Y. Katoh, T. Nozawa, Radiation effects in SiC and SiC–SiC, in: Compr. Nucl. Mater., Elsevier Inc., 2012, pp. 215–240. [17] W. Jiang, H. Wang, I. Kim, I.-T. Bae, G. Li, P. Nachimuthu, et al., Response of nanocrystalline 3C silicon carbide to heavy-ion irradiation, Phys. Rev. B 80 (2009) 161301. [18] X. Kerbiriou, J.-M. Costantini, M. Sauzay, S. Sorieul, L. Thomé, J. Jagielski, et al., Amorphization and dynamic annealing of hexagonal SiC upon heavy-ion irradiation: effects on swelling and mechanical properties, J. Appl. Phys. 105 (2009) 073513. [19] M. Ishimaru, A. Hirata, M. Naito, I.-T. Bae, Y. Zhang, W.J. Weber, Direct observations of thermally induced structural changes in amorphous silicon carbide, J. Appl. Phys. 104 (2008) 033503.

81

[20] W. Jiang, Y. Zhang, W.J. Weber, Temperature dependence of disorder accumulation and amorphization in Au-ion-irradiated 6H-SiC, Phys. Rev. B 70 (16) (2004) 165208. [21] H.G. Bohn, J.M. Williams, C.J. McHargue, G.M. Begun, Recrystallization of ionimplanted alpha-SiC, J. Mater. Res. 2 (1987) 107–116. [22] L.L. Snead, S.J. Zinkle, Structural relaxation of amorphous silicon carbide, Nucl. Instr. Meth. Phys. Res. B 191 (2002) 497–503. [23] I.-T. Bae, M. Ishimaru, Y. Hirotsu, K.E. Sickafus, Solid phase epitaxy of amorphous silicon carbide: ion fluence dependence, J. Appl. Phys. 96 (3) (2004) 1451. [24] Z. Czigány, L. Hultman, Interpretation of electron diffraction patterns from amorphous and fullerene-like carbon allotropes, Ultramicroscopy 110 (2010) 815–819. [25] M. Satoh, Y. Nakaike, T. Nakamura, Solid phase epitaxy of implantation  induced amorphous layer in (1100)and (1120)-oriented 6H-SiC, J. Appl. Phys. 89 (2001) 1986. [26] M. Diani, L. Simon, L. Kubler, D. Aubel, I. Matko, B. Chenevier, et al., Crystal growth of 3C-SiC polytype on 6H-SiC(0 0 0 1) substrate, J. Cryst. Growth 235 (2002) 95–102. [27] P. Pirouz, J.W. Yang, Polytypic transformations of SiC: the role of TEM, Ultramicroscopy 51 (1993) 189–214. [28] D.D. Osterberg, J. Youngsman, R. Ubic, I.E. Reimanis, D.P. Butt, Recrystallization kinetics of 3C silicon carbide implanted with 400 keV cesium ions, J. Am. Ceram. Soc. 96 (2013) 3290–3295.