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On the study of microstructure of annealed C+and H2+ Co_implanted 6H-SiC
Progress in Nuclear Energy 118 (2020) 103143
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Progress in Nuclear Energy journal homepage: http://www.elsevier.com/locate/pnucene
_ On the study of microstructure of annealed Cþand Hþ 2 Co implanted 6H-SiC
Chao Ye a, Guang Ran a, *, Wei Zhou b, **, Qijie Feng b, Jingchi Huang a, Ning Li a a b
College of Energy, Xiamen University, Xiamen, Fujian 361102, China China Academy of Engineering Physics, Mianyang City, Sichuan Province 621900, China
A R T I C L E I N F O
A B S T R A C T
Keywords: SiC Ion irradiation Surface morphology Microstructure Annealing
The microstructure of annealed single crystal 6H-SiC implanted with Cþ and Hþ 2 ions were characterized by glancing incidence X-ray diffraction (GIXRD), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). After annealing, cracks were observed to form in the surface of Cþ ions implanted sample. However, blisters as circular clusters were observed in the surface of the Hþ 2 ions implanted and annealed sample. Compared with the samples implanted with only single type of ions, both cracks and blisters with different morphologies were found to appear in the surface of Cþ and Hþ 2 ions co-implanted sample. The morphology and distribution of blisters and exfoliations in the Cþ and Hþ 2 co-implanted samples were more irregular and inho mogeneous than those in the Hþ 2 ions implanted sample, which can be attributed to the randomly distributed columnar crystals formed during the annealing process. The grain boundaries of these columnar crystals could be used as migration channels for hydrogen, which influence the size evolution of blisters and exfoliations in the sample surface during annealing process.
1. Introduction Due to its perfect high-temperature stability, excellent irradiation resistance and low neutron capture cross-section, SiC is known to be a promising material used as the structural support component in Tristructural isotropic (TRISO) nuclear fuel and the next generation fuel cladding in advanced tolerant fuel (Bosi et al., 2016; Linez et al., 2013). Because the stability of SiC after energetic ion irradiation is a key to its applications in nuclear field, a great amount of efforts have been devoted to understand the ion-irradiation induced damage and its evo lution mechanism in recent years (Deng et al., 2013; Chen et al., 2016). As a key component in nuclear field, SiC will be suffered to neutron irradiation that will induce displacement damage. The irradiated Si atoms will transmute and then produce H atoms in service period (Zhang and Zhao, 2014; Zang et al., 2016). The blisters and exfoliations induced by the interaction between displacement damage and H atoms will be easily formed in the irradiated SiC surface (Zhao et al., 2015; Katharria et al., 2007), which severely influence its service performance (Rebillat et al., 2017). For example, Zhang (Zhang and Li, 2017) observed that many blisters formed in the surface of [0001]-oriented 6H-SiC implan ted with 134 keV Hþ 2 ions and subsequently annealed at the temperature from 973 K to 1373 K. Jung-Wuk Hong (Hong and Cheong, 2006)
considered that the hydrogen implantation-induced interface splitting was due to the growth of cracks induced by the rupture of blisters and he proposed a crack growth model. After being implanted with 100 keV Hþ with the fluences ranged from 4.5 � 1016 Hþ/cm2 to 17 þ 2 1.50 � 10 H /cm , the blisters were formed on the SiC surface after being annealed at 1070 K (Jiang et al., 2000). Actually, after the reaction with neutrons, Si atoms will decay to other atoms. For example, the 30Si will be transformed to 31Si after absorbing a neutron and then decay to 31 P atoms (Thomas et al., 2004). Therefore, that process will induce the decrease of Si atoms in the system and leave a lot of carbon atoms. These redundant carbon atoms should have an obvious effect on the micro structure and properties of SiC. After being implanted with 4 MeV Cþ with the fluences of 1 � 1015 Cþ/cm2 and 1 � 1016 Cþ/cm2 at room temperature, the Si–C bonds became broader or even disappeared. When the SiC was irradiated at a high temperature of 670 K, the irradiated layer was highly heterogeneous and composed of different types of de fects (Cha^ abane et al., 2012). However, the question whether the redundant carbon atoms would affect the formation of surface structures such as blisters, exfoliations and cracks has still not been answered yet. Also the evolution of internal microstructure in the Cþ þ Hþ 2 co-implanted SiC is needed to be further investigated. Therefore, in order to deeply understand the evolution mechanism of
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (G. Ran), [email protected] (W. Zhou). https://doi.org/10.1016/j.pnucene.2019.103143 Received 3 August 2018; Received in revised form 16 August 2019; Accepted 26 August 2019 Available online 5 September 2019 0149-1970/© 2019 Elsevier Ltd. All rights reserved.
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the surface structures such as blisters, exfoliations, cracks and internal microstructure of irradiated SiC during annealing process, Cþ and Hþ 2 ions were implanted into single crystal 6H-SiC and annealed subse quently in the present work. The surface morphology and internal microstructure were investigated by scanning electron microscopy (SEM) and transmission electron microscopy (TEM).
and then mechanically polished to approximately 5 μm thickness using diamond sandpapers. The sample was next glued on a copper grid by G-1 epoxy glue and finally thinned to approximately 100 nm thickness via Arþ ion milling in the Gatan 695 PIPS instrument.
2. Experiments
Crystal SiC is easily changed to amorphization after a low dose irradiation at room temperature, such as approximately 1dpa (Matsu naga, 1991). In the present work, the SRIM simulation results show that the irradiation damage peak of the Cþ ion implanted samples with a fluence of 5 � 1016 Cþ/cm2 is greater than 1 dpa. Therefore, almost the whole irradiated regions have been amorphized from the sample sur faces to the depth of about 700 nm as shown in (Harada et al., 1996; A. Hofgen et al., 1998), when the annealing temperature is higher than 700 � C, the irradiation-induced amorphous SiC can be recrystallized at both sides of the irradiated layer. Our previous experimental results (Ye et al., 2017) indicate that the volume shrinkage of the irradiated SiC layer and the anisotropy of newly born crystals during annealing process would produce the internal stress and then induce not only a large number of dislocation walls in the non-irradiated layer but also the initiation and propagation of the cracks as shown in Fig. 1(b). Fig. 2 shows the surface morphology of Cþ ion implanted and annealed samples. From Fig. 2(a), it can be seen that the width of the crack on the surface of the sample annealed at 900 � C is uniform along the whole length. Compared with the sample annealed at 900 � C, the sample annealed at 1200 � C has a feature of downward depression at the surface side before it cracks as shown in Fig. 2(b), which indicates that the crack initiation of the sample annealed at 1200 � C is more likely to start below the surface. Bright field (BF) cross-sectional TEM images of the 5 � 1016 Cþ/cm2 implanted and annealed for 5 h at 900 � C and 1200 � C samples are shown in Fig. 3(a) and Fig. 3(b), respectively. The direction of ion incidence and irradiated sample surface is marked in Fig. 3(a). The inset in Fig. 3(b) is the concentration distribution of implanted carbon ions along the irradiation depth simulated by SRIM software with quick mode. Fig. 3(a) shows that the crack width in the sample annealed at 900 � C decreases from the sample surface to the internal substrate ma trix, which means that the crack is initiated from the sample surface. While for the irradiated sample annealed at 1200 � C, the widest position of the crack is located at the subsurface as indicated by a white arrow. Meanwhile, at the irradiation surface, a distinct concave shape indicated by a black arrow can be observed, which is consistent with the result of Fig. 2(b). The crack initiation is due to the recrystallization of irradiation-induced amorphous layer (Ye et al., 2017), which means that the recrystallization rate at subsurface region is higher than that in the surface region of the sample in the irradiated layer during 1200 � C annealing process. While for the Cþ irradiated and annealed at 900 � C sample, the result is opposite. Actually, the recrystallization rate is influenced by the annealing temperature, irradiated-induced damage and implantation ion concentration (Li et al., 2015). High implantation-induced damage and high implanted ion concentration lead to a high threshold temperature of the recrystallization. Fig. 3(b) shows that the position of maximum value of crack width is located at the depth of approximately 600 nm, which is in accordance with the depth of peak carbon concentration simulated by SRIM soft ware as shown in the red diagram as an inset in Fig. 3(b). This position is also near the peak area of the irradiation-induced damage. This is similar to the phenomenon mentioned by Bae et al. (Bae et al., 2004), that is the region with the peak carbon concentration can retard the recrystalliza tion process during 900 � C annealing process. Leclerc et al. (Leclerc et al., 2010) proposed that the formation of a thin and deep highly strained region near the surface could result in the accumulation of interstitial atoms to this region and then the recombination of vacancy-type defects. This can be used to explain the fast recrystalliza tion rate of the upper region in irradiated layer. Therefore, during 900 � C
3. Results and discussion
Single crystal 6H-SiC wafers with 4� off-axis [0001] orientation were used in the present work. Ion-irradiation experiments were performed at the NEC 400 kV ion implanter in the College of Energy at Xiamen Uni versity. The samples of 10 � 10 � 0.5 mm in dimensions with their surface polished were pasted on the surface of stainless steel stage with Φ180 mm by carbon paste. The beam spot size of carbon (hydrogen) ions was approximately Φ30 mm. The direction of ion incidence was parallel to the normal direction of polished surface. Stopping and Range of Ions in Matter (SRIM) software with quick mode was used to simulate the distribution characteristics of ion irradiation damage and ion concentration. Samples were first implanted by 400 keV Cþ ions at room tempera ture. The ion flux was kept at 1 � 1014 Cþ/cm2⋅s to prevent excessive beam heating. The displacement energy of Si and C was assumed to be 35 eV and 20 eV (Liu et al., 2016), respectively. The peak displacement damage at 550 nm depth and the peak carbon concentration at 600 nm depth were 0.011 dpa (displacement per atom, dpa) and 0.075% after irradiation with fluence of 1.0 � 1014 Cþ/cm2, respectively. Multiplying these values by implanted Cþ ion fluence gave the peak atomic density of implanted carbon and the peak displacement damage. Therefore, the peak displacement damage was 2.75 dpa and 5.5 dpa after Cþ ion irradiation with fluences of 2.5 � 1016 Cþ/cm2 and 5 � 1016 Cþ/cm2, respectively. Some as-received and Cþ implanted SiC samples were 16 further implanted by 100 keV Hþ 2 ions with a fluence of 1.25 � 10 2 Hþ at room temperature. The hydrogen concentration peak 2 /cm (0.011%) and irradiation damage peak (0.045 dpa) appeared at the depth of 600 nm and 550 nm, respectively. Therefore, there were three types of ion irradiated samples including Cþ implanted, Hþ 2 implanted and Cþ-Hþ 2 co-implanted samples. Finally, some implanted samples were annealed at designed experiment conditions in a tube furnace with argon (99.999% purity) protection. The experimental parameters of ion implantation and annealing were listed in Table 1. The morphology and topography of the ion irradiated and annealed SiC surface were observed by SEM (ZEISSEVO18, Germany). The phases in the Cþ ion irradiated layer were characterized by GIXRD with a fixed angle 1� on a Rigaku D/max-3C X-ray diffractometer with CuKα radia tion (λ ¼ 0.1540598 nm). Under the condition of this fixed incidence angle, the X-ray detection depth was approximately 250 nm, which al ways located in the range of the irradiated region according to SRIM simulation results. TEM was used to analyze the microstructure of the irradiated layers. Cross-sectional TEM samples prepared by a method of mechanical thinning and then ion milling were used to analyze the microstructures of irradiated layers in a JEOL 2100 transmission electron microscope. TEM samples were first cut from SiC bulks along ion incidence direction Table 1 Experimental parameters of Cþ and Hþ 2 co-implantation and then annealing. Ion fluence (ions/cm2) C ions þ
2.5 � 1016 5.0 � 1016 5.0 � 1016 5.0 � 1016 2.5 � 1016
Hþ 2
Annealing temperature (� C)
Annealing time (h)
1200 and 900 1200 900 900 900 900
10 0.5, 5 and 10 0.5, 2, 5 and 10 0.5, 2 and 10 0.5, 2 and 10 0.5, 2 and 10
ions
1.25 � 1016 1.25 � 1016 1.25 � 1016
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Fig. 1. Bright field TEM images of the Cþ ion implanted samples with a fluence of 5 � 1016 Cþ/cm2 (a) and then annealed for 10 h at 900 � C (b).
Fig. 2. SEM images showing the surface morphology of the Cþ ion implanted samples with a fluence of 5 � 1016 Cþ.cm at 1200 � C 5 h (b).
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and then annealed at 900 � C for 5 h (a) and
Fig. 3. TEM BF images of the Cþ ion implanted with a fluence of 5 � 1016 Cþ/cm2 and annealed for 5 h at 900 � C (a) and 1200 � C (b) samples.
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annealing process, the cracks initiate in sample surface region and then propagate. While during 1200 � C annealing process, the recrystalliza tion rate at the peak carbon concentration region is higher than that at the region near surface, which indicates that the retardation effect of high carbon concentration and implantation-induced damage during the recrystallization process is avoided by the high annealing temperature. Meanwhile, as an intrinsic element in SiC, the implanted carbon with high concentration can be rapidly recombined with a large number of vacancies induced by Cþ irradiation to promote the recrystallization process. Therefore, when the annealing temperature is high enough to avoid the retardation effect of impurities and irradiation-induced dam age, the recrystallization rate at the highest carbon concentration region will reach to the largest value. More detail research is needed for un derstanding the mechanism behind in the future work. Surface morphologies of the samples implanted with Cþ and/or Hþ 2 ions and annealed at 900 � C for 2 h are shown in Fig. 4. Only cracks can be observed in the surface of the sample implanted with Cþ ions and only blisters distributing as circular clusters can be seen in the surface of SiC implanted with Hþ 2 ions as shown in Fig. 4(a) and (b), respectively. Fig. 4 (c) is a high magnification image of local region in Fig. 4(b). Similar experimental phenomena have been reported in the literature (Zhang and Li, 2017). However, both cracks and blisters were observed in the surface of the 6H-SiC co-implanted with Cþ and Hþ 2 ions. After having been annealed for 2 h, from Fig. 4(d) and (e), it can be seen that most blisters exfoliated from the sample surface. During the annealing pro cess, with the accumulation of hydrogen, the pressure difference be tween the inside and outside of the surface blisters would exceed its fracture strength that leads to the burst of blister to form exfoliation (Hong and Cheong, 2006). After having been further annealed for 10 h at 900 � C, only a small amounts of exfoliations can be observed on the sample surface as shown in Fig. 4(f). This can be attributed to that the blisters are almost converted to exfoliations that drop out from the sample surface with the increase of annealing time. In order to quantitatively describe the evolution of blister size during
the annealing process, the average blister area (Mex) is used that can be total calculated by Mex ¼ Atotal ex =nex equation, where, Aex is the total blister areas obtained by the sum of the individual blister area; nex is the total number of individual blisters that can be counted from the taken SEM images. The red color in Fig. 4(c) indicates the typical blister areas. Quantitative measurement was conducted using Photoshop software. The counting number is over 50 for one experiment condition. The relationship between the average blister area (Mex) of the samples annealed at 900 � C and annealing time is shown in Fig. 5(a), as well as for 1200 � C annealing experiments. Due to the high annealing temper ature at 1200 � C, the blisters would be quickly fragmented and then separated from the sample surface in a short time, which induces the formation of many small-sized pits on the sample surface. The dash line in Fig. 5(b) indicates the irradiated sample surface and two small-sized pits. Therefore, the blister size at the condition of 1200 � C annealing was not calculated here. Fig. 5(a) shows that the average blister size of the sample implanted þ with only Hþ 2 ions is smaller than that of the samples implanted with H2 and Cþ ions during the 900 � C annealing process at the initial stage of 0.5 h. Meanwhile, the average blister size of the sample implanted with high fluence of Cþ ions is relatively larger than that of the sample implanted with low fluence of Cþ ions. In addition, with the increase of annealing time, the average blister size shows a downward trend for all kinds of ion implanted samples. The mechanisms of the blister size evolution would be analyzed with the TEM results as follows. From the TEM images as shown in Fig. 6(a) and Fig. 6(b), the microcrack indicated by a white arrow distributes at approximately 500 nm irradiation depth in the Hþ 2 implanted SiC sample annealed at 900 � C for 2 h with no macrocrack that can be observed in the sample surface as shown in Fig. 3(b). Fig. 6(b) is a high magnification image of the area indicated by the white arrow in Fig. 6(a). The stripes that can be observed in Fig. 6(b) are the Moire Fringes of 6H-SiC lattice stripes, which is parallel to the (0001) crystal planes. It can be seen that the
2 16 þ 2 Fig. 4. SEM images showing the surface morphology of the samples implanted with fluences of (a) 5 � 1016 Cþ/cm2; (b) 1.25 � 1016 Hþ 2 /cm ; (d) 2.5 � 10 C /cm 2 16 þ 2 16 2 þ � and 1.25 � 1016 Hþ /cm , (e) 5 � 10 C /cm and 1.25 � 10 H /cm and annealed at 900 C for 2 h; (c) High magnification of image (b); (f) The surface 2 2 2 � morphology of the samples implanted with fluences of 5 � 1016 Cþ/cm2 and 1.25 � 1016 Hþ 2 /cm and annealed at 900 C for 10 h.
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Fig. 5. (a) The average area of the blisters of the samples annealed at 900 � C vs. annealing time. The samples previously implanted by one or both of Cþ and Hþ 2 ions. 2 � (b) Cross-sectional BF TEM image of the 6H-SiC implanted with fluences of 5 � 1016 Cþ/cm2 and 1.25 � 1016 Hþ 2 /cm and annealed at 1200 C for 0.5 h.
2 16 þ 2 16 2 Fig. 6. Under-focused BF TEM images of the 6H-SiC implanted with fluences of (a) 1.25 � 1016 Hþ Hþ 2 /cm , (c) and (d) 5 � 10 C /cm and 1.25 � 10 2 /cm ; (b) � � High magnification image of the area indicated by a white arrow in (a); The annealing condition is at (a) and (c) 900 C for 2 h and (d) 1200 C for 2 h.
implanted hydrogen and vacancies induced by ion irradiation were accumulated at the interface between (0001) crystal planes and caused the lattice distortion of some local (0001) planes, which is consistent with the Zhang’s results (Zhang and Zhao, 2014). There is no obvious migration channel for hydrogen in the Hþ 2 implanted SiC under the TEM observation. Actually, at room temperature, SiC will be easily
amorphized when the irradiation damage was up to approximately 1dpa (Matsunaga, 1991). While in the present work, the Hþ 2 ion fluence is 2 1.25 � 1016 Hþ 2 /cm that induces a peak irradiation damage with 0.045 dpa, which means that only a small fraction of SiC has been amorphized in the irradiated layer. After having been further implanted with a flu ence of 5 � 1016 Cþ/cm2, the SiC in the Hþ 2 implanted layer had been 5
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completely amorphized because the peak irradiation damage is up to 5.545 dpa. After being annealed for several hours at 900 � C and 1200 � C, the amorphized SiC would be recrystallized in the form of columnar crystals (Li et al., 2015; Manabu et al., 2008). The grain boundaries of these columnar crystals as indicated by black arrows can be used as the migration channels for hydrogen and vacancies as shown in Fig. 6(c) and (d). These channels allow the implanted hydrogen and irradiation-induce vacancies to migrate and gather much easily. As a result, hydrogen moved much faster in the sample co-implanted with Cþ þ and Hþ 2 ions over in the sample only implanted with H2 ions. Therefore, at the initial stage of 0.5 h annealing process, the average blister size of þ þ the Hþ 2 and C co-implanted sample was larger than that of the only H2 implanted sample. When the annealing temperature increased from 900 � C to 1200 � C, the hydrogen gathered more quickly. Statistical re sults show that the size of hydrogen bubbles in Fig. 6(d) is larger than that in Fig. 6(c). When the annealing time was increased to more than 0.5 h, the average size of blisters has a decreasing trend as shown in Fig. 5(a), which should be due to that most of the large-sized blisters were broken and then exfoliated from sample surface. Therefore, the average size of remnant blisters is relatively small as the amount of hydrogen gathering to the surface became less and less. In addition, the average diameter (Dex) of blister and exfoliation (BE) clusters is used to quantitatively describe the aggregation situation of blisters and exfoliations. And it can be estimated by Dex ¼ Dtotal ex = Nex
þ both of Hþ 2 and C ions implanted and annealed single crystal 6H-SiC wafers with 4� off-axis [0001] orientation were investigated by GIXRD, SEM and TEM. Conclusions are made as follows:
(1) After having been annealed at both of 900 � C and 1200 � C, only cracks could be observed in the surface of sample implanted with Cþ ions and only blisters distributing as circular clusters could be seen in the surface of SiC implanted with Hþ 2 ions. While both of cracks and blisters appeared in the surface of SiC co-implanted with Cþ and Hþ 2 ions. (2) The cracks initiated from the sample interior in the Cþ implanted and annealed at 1200 � C SiC sample, while the cracks initiated from the sample surface after having been annealed at 900 � C. The reason could be attributed that the retardation effect of high carbon concentration and high irradiation-induced damage at 900 � C annealing, which is not applicable for the 1200 � C annealed sample. (3) Compared with the sample implanted with only Hþ 2 ions, the distribution of blisters and exfoliations in the surface of the Cþ and Hþ 2 co-implanted samples are much irregular, which is due to the randomly distributed columnar crystals formed during the annealing process. The grain boundaries of these columnar crystals could be used as the migration channels for hydrogen and vacancies, which led to different levels of hydrogen and vacancy concentration in the sample surface. (4) The average blister size of the sample co-implanted with Cþ and þ Hþ 2 ions is larger than that of the sample only implanted with H2 ions at the initial stage of 0.5 h annealing process with the help of these migration channels. And large blisters broke through the surface to form the exfoliations with the continue of annealing process, which results in the size reduction of the remaining blisters as the amount of implanted-hydrogen and the irradiationinduced vacancies is limited.
equation, where, Dtotal and Nex are the sum of the equivalent Dsingle ex ex diameter and the total number of B-E clusters, respectively. The actual area of B-E cluster measured in Photoshop software was converted to an equivalent circular area that can be used to calculate equivalent Dsingle ex value. The statistical analysis results show that the average diameters (Dex) are 3 μm and 12 μm for the samples implanted with a fluence of 2 � 1.25 � 1016 Hþ 2 /cm and then annealed at 900 C for 0.5 h and 2 h, respectively. Correspondingly, the Dex are approximately 18 μm and 26 μm for the samples annealed at 900 � C for 2 h after having been 2 16 þ 2 implanted with fluences of 1.25 � 1016 Hþ 2 /cm -2.5 � 10 C /cm and 16 þ 2 16 þ 2 1.25 � 10 H2 /cm -5x10 C /cm , respectively. It can be seen that the size of B-E clusters in the Hþ 2 implanted sample surface increases with the increase of Hþ 2 fluence, which is attributed to that the amount of hydrogen migrating to the surface increases with the increase of annealing time in a certain time period. Meanwhile, the size of B-E þ clusters in the Hþ 2 and C co-implanted sample surface increases with the þ increase of C fluence, which is due to that much more vacancies moved to the blisters. The higher the Cþ fluence is, the larger the amount of vacancies will be. Compared with the sample implanted with only Hþ 2 ions, the shape characteristic of blisters and exfoliations of the Cþ and Hþ 2 co-implanted sample is much irregular as shown in Fig. 4. The microcracks filled with hydrogen and vacancies show a homogeneous distribution characteristic in the irradiated region as shown in Fig. 6(a). The size and shape of B-E clusters in the Hþ 2 implanted sample surface also show a homogeneous distribution characteristic. While for the Cþ and Hþ 2 co-implanted SiC, the irradiated region had been completely amorphized after implanta tion and then were recrystallized during the annealing process. The recrystallized columnar crystals are randomly distributed in the irradi ated layer. As a result, the grain boundaries that can be used as the migration channels of hydrogen are also randomly distributed in the irradiated layer. Therefore, a large amount of hydrogen and vacancies could be accumulated at the regions that these channels connect with in the sample surface. While in the rest implanted areas, the amount of hydrogen is relatively small. This is why the distribution characteristics of blisters and exfoliations in the Cþ and Hþ 2 co-implanted sample sur face are so irregular.
Acknowledgement This research was supported by National Natural Science Foundation of China under Grant No. U1832112. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.pnucene.2019.103143. References Bae, In-Tae, Manabu, I., Yoshihiko, H., 2004. Solid phase epitaxy of amorphous silicon carbide: ion fluence dependence. J. Appl. Phys. 96 (3), 1451–1457. Bosi, M., Ferrari, C., Nilssonb, D., Wardb, P.J., 2016. 3C-SiC carbonization optimization and void reduction on misoriented Si substrates: from a research reactor to a production scale reactor. CrystEngComm 18 (18), 7478–7486. Cha^ abane, N., Debell, A., Sattonnay, G., Trocellier, P., Serruys, Y., Thom�e, L., Zhang, Y., Weber, W.J., Meis, C., Gosmain, L., Boulle, A., 2012. Investigation of irradiation effects induced by self-ion in 6H-SiC combining RBS/C, Raman and XRD. Nucl. Instrum. Methods B 286, 108–113. Chen, X.F., Zhou, W., Feng, Q.J., Zheng, J., Liu, X.K., Tang, B., Li, J.B., Xue, J.M., Peng, S. M., 2016. Irradiation effects in 6H–SiC induced by neutron and heavy ions: Raman spectroscopy and high-resolution XRD analysis. J. Nucl. Mater. 478, 215–221. Deng, J.H., Sun, P.C., Cheng, G.A., Zheng, R.T., 2013. Improved field electron emission from SiC assisted carbon nanorod/nanotube heterostructured arrays by using energetic Si ion irradiation. J. Nucl. Mater. 228, 323–327. Harada, S., Ishimaru, M., Motooka, T., Nakata, T., Yoneda, T., Inoue, M., 1996. Recrystallization of MeV Si implanted 6H-SiC. Appl. Phys. Lett. 69 (23), 3534–3536. Hofgen, A., Heera, V., Eichhorn, F., Skorupa, W., 1998. Annealing and recrystallization of amorphous silicon carbide produced by ion implantation. J. Appl. Phys. 84 (9), 4769–4774. Jiang, W., Thevuthasan, S., Grotzschel, R., Weber, W.J., 2000. Irradiation effects and thermal annealing behavior in H-2(þ)-implanted 6H-SiC. Nucl. Instrum. Methods B 166, 374–378. Hong, Jung-Wuk, Cheong, Soonwuk, 2006. A crack model for the onset of blisters using finite surface thicknesses. J. Appl. Phys. 100 (9), 094322-P1-P4.
4. Conclusions The surface morphology and the internal microstructure of one or 6
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Katharria, Y.S., Kumar, S., Lakshmy, P.S., Kanjilal, D., 2007. Self-organization of 6H-SiC (0001) surface under keV ion irradiation. J. Appl. Phys. 102 (4), 44301–44306. Leclerc, S., Beaufort, M.F., Declemy, A., Barbot, J.F., 2010. Strain-induced drift of interstitial atoms in SiC implanted with helium ions at elevated temperature. J. Nucl. Mater. 397 (1–3), 132–134. Li, B.S., Du, Y.Y., Wang, Z.G., 2015. Recrystallization of He-ion implanted 6H-SiC upon annealing. Nucl. Instrum. Methods B 345, 53–57. Linez, F., Gilabert, E., Debelle, A., Desgardin, P., Barthe, M.-F., 2013. Helium interaction with vacancy-type defects created in silicon carbide single crystal. J. Nucl. Mater. 436 (1–3), 150–157. Liu, W., Cheng, L.F., Wang, Y.G., Ma, H.J., 2016. Investigation of the residual stress in reaction-bonded SiC under irradiation. J. Eur. Ceram. Soc. 36, 3901–3907. Manabu, I., Akihiko, H., Muneyuki, N., In-Tae, B., Zhang, Yanwen, Weber, William J., 2008. Direct observations of thermally induced structural changes in amorphous silicon carbide. J. Appl. Phys. 104 (3), 33503–33507. Matsunaga, A., 1991. Radiation-induced amorphization and swelling in ceramics. J. Nucl. Mater. 179–181 (1), 457–460. Rebillat, F., Bouillon, F., Camus, G., Simon, C., 2017. Modeling of simultaneous oxidation and volatilization phenomena along a crack in a self-healing multi-constituent material. Oxid. Metals 1–12.
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Progress in Nuclear Energy journal homepage: http://www.elsevier.com/locate/pnucene
_ On the study of microstructure of annealed Cþand Hþ 2 Co implanted 6H-SiC
Chao Ye a, Guang Ran a, *, Wei Zhou b, **, Qijie Feng b, Jingchi Huang a, Ning Li a a b
College of Energy, Xiamen University, Xiamen, Fujian 361102, China China Academy of Engineering Physics, Mianyang City, Sichuan Province 621900, China
A R T I C L E I N F O
A B S T R A C T
Keywords: SiC Ion irradiation Surface morphology Microstructure Annealing
The microstructure of annealed single crystal 6H-SiC implanted with Cþ and Hþ 2 ions were characterized by glancing incidence X-ray diffraction (GIXRD), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). After annealing, cracks were observed to form in the surface of Cþ ions implanted sample. However, blisters as circular clusters were observed in the surface of the Hþ 2 ions implanted and annealed sample. Compared with the samples implanted with only single type of ions, both cracks and blisters with different morphologies were found to appear in the surface of Cþ and Hþ 2 ions co-implanted sample. The morphology and distribution of blisters and exfoliations in the Cþ and Hþ 2 co-implanted samples were more irregular and inho mogeneous than those in the Hþ 2 ions implanted sample, which can be attributed to the randomly distributed columnar crystals formed during the annealing process. The grain boundaries of these columnar crystals could be used as migration channels for hydrogen, which influence the size evolution of blisters and exfoliations in the sample surface during annealing process.
1. Introduction Due to its perfect high-temperature stability, excellent irradiation resistance and low neutron capture cross-section, SiC is known to be a promising material used as the structural support component in Tristructural isotropic (TRISO) nuclear fuel and the next generation fuel cladding in advanced tolerant fuel (Bosi et al., 2016; Linez et al., 2013). Because the stability of SiC after energetic ion irradiation is a key to its applications in nuclear field, a great amount of efforts have been devoted to understand the ion-irradiation induced damage and its evo lution mechanism in recent years (Deng et al., 2013; Chen et al., 2016). As a key component in nuclear field, SiC will be suffered to neutron irradiation that will induce displacement damage. The irradiated Si atoms will transmute and then produce H atoms in service period (Zhang and Zhao, 2014; Zang et al., 2016). The blisters and exfoliations induced by the interaction between displacement damage and H atoms will be easily formed in the irradiated SiC surface (Zhao et al., 2015; Katharria et al., 2007), which severely influence its service performance (Rebillat et al., 2017). For example, Zhang (Zhang and Li, 2017) observed that many blisters formed in the surface of [0001]-oriented 6H-SiC implan ted with 134 keV Hþ 2 ions and subsequently annealed at the temperature from 973 K to 1373 K. Jung-Wuk Hong (Hong and Cheong, 2006)
considered that the hydrogen implantation-induced interface splitting was due to the growth of cracks induced by the rupture of blisters and he proposed a crack growth model. After being implanted with 100 keV Hþ with the fluences ranged from 4.5 � 1016 Hþ/cm2 to 17 þ 2 1.50 � 10 H /cm , the blisters were formed on the SiC surface after being annealed at 1070 K (Jiang et al., 2000). Actually, after the reaction with neutrons, Si atoms will decay to other atoms. For example, the 30Si will be transformed to 31Si after absorbing a neutron and then decay to 31 P atoms (Thomas et al., 2004). Therefore, that process will induce the decrease of Si atoms in the system and leave a lot of carbon atoms. These redundant carbon atoms should have an obvious effect on the micro structure and properties of SiC. After being implanted with 4 MeV Cþ with the fluences of 1 � 1015 Cþ/cm2 and 1 � 1016 Cþ/cm2 at room temperature, the Si–C bonds became broader or even disappeared. When the SiC was irradiated at a high temperature of 670 K, the irradiated layer was highly heterogeneous and composed of different types of de fects (Cha^ abane et al., 2012). However, the question whether the redundant carbon atoms would affect the formation of surface structures such as blisters, exfoliations and cracks has still not been answered yet. Also the evolution of internal microstructure in the Cþ þ Hþ 2 co-implanted SiC is needed to be further investigated. Therefore, in order to deeply understand the evolution mechanism of
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (G. Ran), [email protected] (W. Zhou). https://doi.org/10.1016/j.pnucene.2019.103143 Received 3 August 2018; Received in revised form 16 August 2019; Accepted 26 August 2019 Available online 5 September 2019 0149-1970/© 2019 Elsevier Ltd. All rights reserved.
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the surface structures such as blisters, exfoliations, cracks and internal microstructure of irradiated SiC during annealing process, Cþ and Hþ 2 ions were implanted into single crystal 6H-SiC and annealed subse quently in the present work. The surface morphology and internal microstructure were investigated by scanning electron microscopy (SEM) and transmission electron microscopy (TEM).
and then mechanically polished to approximately 5 μm thickness using diamond sandpapers. The sample was next glued on a copper grid by G-1 epoxy glue and finally thinned to approximately 100 nm thickness via Arþ ion milling in the Gatan 695 PIPS instrument.
2. Experiments
Crystal SiC is easily changed to amorphization after a low dose irradiation at room temperature, such as approximately 1dpa (Matsu naga, 1991). In the present work, the SRIM simulation results show that the irradiation damage peak of the Cþ ion implanted samples with a fluence of 5 � 1016 Cþ/cm2 is greater than 1 dpa. Therefore, almost the whole irradiated regions have been amorphized from the sample sur faces to the depth of about 700 nm as shown in (Harada et al., 1996; A. Hofgen et al., 1998), when the annealing temperature is higher than 700 � C, the irradiation-induced amorphous SiC can be recrystallized at both sides of the irradiated layer. Our previous experimental results (Ye et al., 2017) indicate that the volume shrinkage of the irradiated SiC layer and the anisotropy of newly born crystals during annealing process would produce the internal stress and then induce not only a large number of dislocation walls in the non-irradiated layer but also the initiation and propagation of the cracks as shown in Fig. 1(b). Fig. 2 shows the surface morphology of Cþ ion implanted and annealed samples. From Fig. 2(a), it can be seen that the width of the crack on the surface of the sample annealed at 900 � C is uniform along the whole length. Compared with the sample annealed at 900 � C, the sample annealed at 1200 � C has a feature of downward depression at the surface side before it cracks as shown in Fig. 2(b), which indicates that the crack initiation of the sample annealed at 1200 � C is more likely to start below the surface. Bright field (BF) cross-sectional TEM images of the 5 � 1016 Cþ/cm2 implanted and annealed for 5 h at 900 � C and 1200 � C samples are shown in Fig. 3(a) and Fig. 3(b), respectively. The direction of ion incidence and irradiated sample surface is marked in Fig. 3(a). The inset in Fig. 3(b) is the concentration distribution of implanted carbon ions along the irradiation depth simulated by SRIM software with quick mode. Fig. 3(a) shows that the crack width in the sample annealed at 900 � C decreases from the sample surface to the internal substrate ma trix, which means that the crack is initiated from the sample surface. While for the irradiated sample annealed at 1200 � C, the widest position of the crack is located at the subsurface as indicated by a white arrow. Meanwhile, at the irradiation surface, a distinct concave shape indicated by a black arrow can be observed, which is consistent with the result of Fig. 2(b). The crack initiation is due to the recrystallization of irradiation-induced amorphous layer (Ye et al., 2017), which means that the recrystallization rate at subsurface region is higher than that in the surface region of the sample in the irradiated layer during 1200 � C annealing process. While for the Cþ irradiated and annealed at 900 � C sample, the result is opposite. Actually, the recrystallization rate is influenced by the annealing temperature, irradiated-induced damage and implantation ion concentration (Li et al., 2015). High implantation-induced damage and high implanted ion concentration lead to a high threshold temperature of the recrystallization. Fig. 3(b) shows that the position of maximum value of crack width is located at the depth of approximately 600 nm, which is in accordance with the depth of peak carbon concentration simulated by SRIM soft ware as shown in the red diagram as an inset in Fig. 3(b). This position is also near the peak area of the irradiation-induced damage. This is similar to the phenomenon mentioned by Bae et al. (Bae et al., 2004), that is the region with the peak carbon concentration can retard the recrystalliza tion process during 900 � C annealing process. Leclerc et al. (Leclerc et al., 2010) proposed that the formation of a thin and deep highly strained region near the surface could result in the accumulation of interstitial atoms to this region and then the recombination of vacancy-type defects. This can be used to explain the fast recrystalliza tion rate of the upper region in irradiated layer. Therefore, during 900 � C
3. Results and discussion
Single crystal 6H-SiC wafers with 4� off-axis [0001] orientation were used in the present work. Ion-irradiation experiments were performed at the NEC 400 kV ion implanter in the College of Energy at Xiamen Uni versity. The samples of 10 � 10 � 0.5 mm in dimensions with their surface polished were pasted on the surface of stainless steel stage with Φ180 mm by carbon paste. The beam spot size of carbon (hydrogen) ions was approximately Φ30 mm. The direction of ion incidence was parallel to the normal direction of polished surface. Stopping and Range of Ions in Matter (SRIM) software with quick mode was used to simulate the distribution characteristics of ion irradiation damage and ion concentration. Samples were first implanted by 400 keV Cþ ions at room tempera ture. The ion flux was kept at 1 � 1014 Cþ/cm2⋅s to prevent excessive beam heating. The displacement energy of Si and C was assumed to be 35 eV and 20 eV (Liu et al., 2016), respectively. The peak displacement damage at 550 nm depth and the peak carbon concentration at 600 nm depth were 0.011 dpa (displacement per atom, dpa) and 0.075% after irradiation with fluence of 1.0 � 1014 Cþ/cm2, respectively. Multiplying these values by implanted Cþ ion fluence gave the peak atomic density of implanted carbon and the peak displacement damage. Therefore, the peak displacement damage was 2.75 dpa and 5.5 dpa after Cþ ion irradiation with fluences of 2.5 � 1016 Cþ/cm2 and 5 � 1016 Cþ/cm2, respectively. Some as-received and Cþ implanted SiC samples were 16 further implanted by 100 keV Hþ 2 ions with a fluence of 1.25 � 10 2 Hþ at room temperature. The hydrogen concentration peak 2 /cm (0.011%) and irradiation damage peak (0.045 dpa) appeared at the depth of 600 nm and 550 nm, respectively. Therefore, there were three types of ion irradiated samples including Cþ implanted, Hþ 2 implanted and Cþ-Hþ 2 co-implanted samples. Finally, some implanted samples were annealed at designed experiment conditions in a tube furnace with argon (99.999% purity) protection. The experimental parameters of ion implantation and annealing were listed in Table 1. The morphology and topography of the ion irradiated and annealed SiC surface were observed by SEM (ZEISSEVO18, Germany). The phases in the Cþ ion irradiated layer were characterized by GIXRD with a fixed angle 1� on a Rigaku D/max-3C X-ray diffractometer with CuKα radia tion (λ ¼ 0.1540598 nm). Under the condition of this fixed incidence angle, the X-ray detection depth was approximately 250 nm, which al ways located in the range of the irradiated region according to SRIM simulation results. TEM was used to analyze the microstructure of the irradiated layers. Cross-sectional TEM samples prepared by a method of mechanical thinning and then ion milling were used to analyze the microstructures of irradiated layers in a JEOL 2100 transmission electron microscope. TEM samples were first cut from SiC bulks along ion incidence direction Table 1 Experimental parameters of Cþ and Hþ 2 co-implantation and then annealing. Ion fluence (ions/cm2) C ions þ
2.5 � 1016 5.0 � 1016 5.0 � 1016 5.0 � 1016 2.5 � 1016
Hþ 2
Annealing temperature (� C)
Annealing time (h)
1200 and 900 1200 900 900 900 900
10 0.5, 5 and 10 0.5, 2, 5 and 10 0.5, 2 and 10 0.5, 2 and 10 0.5, 2 and 10
ions
1.25 � 1016 1.25 � 1016 1.25 � 1016
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Fig. 1. Bright field TEM images of the Cþ ion implanted samples with a fluence of 5 � 1016 Cþ/cm2 (a) and then annealed for 10 h at 900 � C (b).
Fig. 2. SEM images showing the surface morphology of the Cþ ion implanted samples with a fluence of 5 � 1016 Cþ.cm at 1200 � C 5 h (b).
2
and then annealed at 900 � C for 5 h (a) and
Fig. 3. TEM BF images of the Cþ ion implanted with a fluence of 5 � 1016 Cþ/cm2 and annealed for 5 h at 900 � C (a) and 1200 � C (b) samples.
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annealing process, the cracks initiate in sample surface region and then propagate. While during 1200 � C annealing process, the recrystalliza tion rate at the peak carbon concentration region is higher than that at the region near surface, which indicates that the retardation effect of high carbon concentration and implantation-induced damage during the recrystallization process is avoided by the high annealing temperature. Meanwhile, as an intrinsic element in SiC, the implanted carbon with high concentration can be rapidly recombined with a large number of vacancies induced by Cþ irradiation to promote the recrystallization process. Therefore, when the annealing temperature is high enough to avoid the retardation effect of impurities and irradiation-induced dam age, the recrystallization rate at the highest carbon concentration region will reach to the largest value. More detail research is needed for un derstanding the mechanism behind in the future work. Surface morphologies of the samples implanted with Cþ and/or Hþ 2 ions and annealed at 900 � C for 2 h are shown in Fig. 4. Only cracks can be observed in the surface of the sample implanted with Cþ ions and only blisters distributing as circular clusters can be seen in the surface of SiC implanted with Hþ 2 ions as shown in Fig. 4(a) and (b), respectively. Fig. 4 (c) is a high magnification image of local region in Fig. 4(b). Similar experimental phenomena have been reported in the literature (Zhang and Li, 2017). However, both cracks and blisters were observed in the surface of the 6H-SiC co-implanted with Cþ and Hþ 2 ions. After having been annealed for 2 h, from Fig. 4(d) and (e), it can be seen that most blisters exfoliated from the sample surface. During the annealing pro cess, with the accumulation of hydrogen, the pressure difference be tween the inside and outside of the surface blisters would exceed its fracture strength that leads to the burst of blister to form exfoliation (Hong and Cheong, 2006). After having been further annealed for 10 h at 900 � C, only a small amounts of exfoliations can be observed on the sample surface as shown in Fig. 4(f). This can be attributed to that the blisters are almost converted to exfoliations that drop out from the sample surface with the increase of annealing time. In order to quantitatively describe the evolution of blister size during
the annealing process, the average blister area (Mex) is used that can be total calculated by Mex ¼ Atotal ex =nex equation, where, Aex is the total blister areas obtained by the sum of the individual blister area; nex is the total number of individual blisters that can be counted from the taken SEM images. The red color in Fig. 4(c) indicates the typical blister areas. Quantitative measurement was conducted using Photoshop software. The counting number is over 50 for one experiment condition. The relationship between the average blister area (Mex) of the samples annealed at 900 � C and annealing time is shown in Fig. 5(a), as well as for 1200 � C annealing experiments. Due to the high annealing temper ature at 1200 � C, the blisters would be quickly fragmented and then separated from the sample surface in a short time, which induces the formation of many small-sized pits on the sample surface. The dash line in Fig. 5(b) indicates the irradiated sample surface and two small-sized pits. Therefore, the blister size at the condition of 1200 � C annealing was not calculated here. Fig. 5(a) shows that the average blister size of the sample implanted þ with only Hþ 2 ions is smaller than that of the samples implanted with H2 and Cþ ions during the 900 � C annealing process at the initial stage of 0.5 h. Meanwhile, the average blister size of the sample implanted with high fluence of Cþ ions is relatively larger than that of the sample implanted with low fluence of Cþ ions. In addition, with the increase of annealing time, the average blister size shows a downward trend for all kinds of ion implanted samples. The mechanisms of the blister size evolution would be analyzed with the TEM results as follows. From the TEM images as shown in Fig. 6(a) and Fig. 6(b), the microcrack indicated by a white arrow distributes at approximately 500 nm irradiation depth in the Hþ 2 implanted SiC sample annealed at 900 � C for 2 h with no macrocrack that can be observed in the sample surface as shown in Fig. 3(b). Fig. 6(b) is a high magnification image of the area indicated by the white arrow in Fig. 6(a). The stripes that can be observed in Fig. 6(b) are the Moire Fringes of 6H-SiC lattice stripes, which is parallel to the (0001) crystal planes. It can be seen that the
2 16 þ 2 Fig. 4. SEM images showing the surface morphology of the samples implanted with fluences of (a) 5 � 1016 Cþ/cm2; (b) 1.25 � 1016 Hþ 2 /cm ; (d) 2.5 � 10 C /cm 2 16 þ 2 16 2 þ � and 1.25 � 1016 Hþ /cm , (e) 5 � 10 C /cm and 1.25 � 10 H /cm and annealed at 900 C for 2 h; (c) High magnification of image (b); (f) The surface 2 2 2 � morphology of the samples implanted with fluences of 5 � 1016 Cþ/cm2 and 1.25 � 1016 Hþ 2 /cm and annealed at 900 C for 10 h.
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Fig. 5. (a) The average area of the blisters of the samples annealed at 900 � C vs. annealing time. The samples previously implanted by one or both of Cþ and Hþ 2 ions. 2 � (b) Cross-sectional BF TEM image of the 6H-SiC implanted with fluences of 5 � 1016 Cþ/cm2 and 1.25 � 1016 Hþ 2 /cm and annealed at 1200 C for 0.5 h.
2 16 þ 2 16 2 Fig. 6. Under-focused BF TEM images of the 6H-SiC implanted with fluences of (a) 1.25 � 1016 Hþ Hþ 2 /cm , (c) and (d) 5 � 10 C /cm and 1.25 � 10 2 /cm ; (b) � � High magnification image of the area indicated by a white arrow in (a); The annealing condition is at (a) and (c) 900 C for 2 h and (d) 1200 C for 2 h.
implanted hydrogen and vacancies induced by ion irradiation were accumulated at the interface between (0001) crystal planes and caused the lattice distortion of some local (0001) planes, which is consistent with the Zhang’s results (Zhang and Zhao, 2014). There is no obvious migration channel for hydrogen in the Hþ 2 implanted SiC under the TEM observation. Actually, at room temperature, SiC will be easily
amorphized when the irradiation damage was up to approximately 1dpa (Matsunaga, 1991). While in the present work, the Hþ 2 ion fluence is 2 1.25 � 1016 Hþ 2 /cm that induces a peak irradiation damage with 0.045 dpa, which means that only a small fraction of SiC has been amorphized in the irradiated layer. After having been further implanted with a flu ence of 5 � 1016 Cþ/cm2, the SiC in the Hþ 2 implanted layer had been 5
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completely amorphized because the peak irradiation damage is up to 5.545 dpa. After being annealed for several hours at 900 � C and 1200 � C, the amorphized SiC would be recrystallized in the form of columnar crystals (Li et al., 2015; Manabu et al., 2008). The grain boundaries of these columnar crystals as indicated by black arrows can be used as the migration channels for hydrogen and vacancies as shown in Fig. 6(c) and (d). These channels allow the implanted hydrogen and irradiation-induce vacancies to migrate and gather much easily. As a result, hydrogen moved much faster in the sample co-implanted with Cþ þ and Hþ 2 ions over in the sample only implanted with H2 ions. Therefore, at the initial stage of 0.5 h annealing process, the average blister size of þ þ the Hþ 2 and C co-implanted sample was larger than that of the only H2 implanted sample. When the annealing temperature increased from 900 � C to 1200 � C, the hydrogen gathered more quickly. Statistical re sults show that the size of hydrogen bubbles in Fig. 6(d) is larger than that in Fig. 6(c). When the annealing time was increased to more than 0.5 h, the average size of blisters has a decreasing trend as shown in Fig. 5(a), which should be due to that most of the large-sized blisters were broken and then exfoliated from sample surface. Therefore, the average size of remnant blisters is relatively small as the amount of hydrogen gathering to the surface became less and less. In addition, the average diameter (Dex) of blister and exfoliation (BE) clusters is used to quantitatively describe the aggregation situation of blisters and exfoliations. And it can be estimated by Dex ¼ Dtotal ex = Nex
þ both of Hþ 2 and C ions implanted and annealed single crystal 6H-SiC wafers with 4� off-axis [0001] orientation were investigated by GIXRD, SEM and TEM. Conclusions are made as follows:
(1) After having been annealed at both of 900 � C and 1200 � C, only cracks could be observed in the surface of sample implanted with Cþ ions and only blisters distributing as circular clusters could be seen in the surface of SiC implanted with Hþ 2 ions. While both of cracks and blisters appeared in the surface of SiC co-implanted with Cþ and Hþ 2 ions. (2) The cracks initiated from the sample interior in the Cþ implanted and annealed at 1200 � C SiC sample, while the cracks initiated from the sample surface after having been annealed at 900 � C. The reason could be attributed that the retardation effect of high carbon concentration and high irradiation-induced damage at 900 � C annealing, which is not applicable for the 1200 � C annealed sample. (3) Compared with the sample implanted with only Hþ 2 ions, the distribution of blisters and exfoliations in the surface of the Cþ and Hþ 2 co-implanted samples are much irregular, which is due to the randomly distributed columnar crystals formed during the annealing process. The grain boundaries of these columnar crystals could be used as the migration channels for hydrogen and vacancies, which led to different levels of hydrogen and vacancy concentration in the sample surface. (4) The average blister size of the sample co-implanted with Cþ and þ Hþ 2 ions is larger than that of the sample only implanted with H2 ions at the initial stage of 0.5 h annealing process with the help of these migration channels. And large blisters broke through the surface to form the exfoliations with the continue of annealing process, which results in the size reduction of the remaining blisters as the amount of implanted-hydrogen and the irradiationinduced vacancies is limited.
equation, where, Dtotal and Nex are the sum of the equivalent Dsingle ex ex diameter and the total number of B-E clusters, respectively. The actual area of B-E cluster measured in Photoshop software was converted to an equivalent circular area that can be used to calculate equivalent Dsingle ex value. The statistical analysis results show that the average diameters (Dex) are 3 μm and 12 μm for the samples implanted with a fluence of 2 � 1.25 � 1016 Hþ 2 /cm and then annealed at 900 C for 0.5 h and 2 h, respectively. Correspondingly, the Dex are approximately 18 μm and 26 μm for the samples annealed at 900 � C for 2 h after having been 2 16 þ 2 implanted with fluences of 1.25 � 1016 Hþ 2 /cm -2.5 � 10 C /cm and 16 þ 2 16 þ 2 1.25 � 10 H2 /cm -5x10 C /cm , respectively. It can be seen that the size of B-E clusters in the Hþ 2 implanted sample surface increases with the increase of Hþ 2 fluence, which is attributed to that the amount of hydrogen migrating to the surface increases with the increase of annealing time in a certain time period. Meanwhile, the size of B-E þ clusters in the Hþ 2 and C co-implanted sample surface increases with the þ increase of C fluence, which is due to that much more vacancies moved to the blisters. The higher the Cþ fluence is, the larger the amount of vacancies will be. Compared with the sample implanted with only Hþ 2 ions, the shape characteristic of blisters and exfoliations of the Cþ and Hþ 2 co-implanted sample is much irregular as shown in Fig. 4. The microcracks filled with hydrogen and vacancies show a homogeneous distribution characteristic in the irradiated region as shown in Fig. 6(a). The size and shape of B-E clusters in the Hþ 2 implanted sample surface also show a homogeneous distribution characteristic. While for the Cþ and Hþ 2 co-implanted SiC, the irradiated region had been completely amorphized after implanta tion and then were recrystallized during the annealing process. The recrystallized columnar crystals are randomly distributed in the irradi ated layer. As a result, the grain boundaries that can be used as the migration channels of hydrogen are also randomly distributed in the irradiated layer. Therefore, a large amount of hydrogen and vacancies could be accumulated at the regions that these channels connect with in the sample surface. While in the rest implanted areas, the amount of hydrogen is relatively small. This is why the distribution characteristics of blisters and exfoliations in the Cþ and Hþ 2 co-implanted sample sur face are so irregular.
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