⇦Chinese Physics Letters, 2020, Vol. 37, No. 7, Article code 076102 Comparison of Cavities Formed in Single Crystalline and Polycrystalline $\alpha$-SiC after H Implantation Qing Liao (廖庆)1, Long Kang (康龙)2, Tong-Min Zhang (张桐民)2, Hui-Ping Liu (刘会平)2, Tao Wang (王韬)3, Xiao-Gang Li (李小刚)2, Jin-Yu Li (李锦钰)2, Zhen Yang (杨振)4, and Bing-Sheng Li (李炳生)1* Affiliations 1State Key Laboratory for Environment-Friendly Energy Materials, Southwest University of Science and Technology, Mianyang 621010, China 2Institute of Modern Physics, Chinese Academy of Sciences, Lanzhou 730000, China 3Institute of Fluid Physics, China Academy of Engineering Physics, Mianyang 621900, China 4Sino-French Institute of Nuclear Engineering and Technology, Sun Yat-Sen University, Zhuhai 519082, China Received 19 March 2020; accepted 19 May 2020; published online 21 June 2020 Supported by the National Natural Science Foundation of China (Grant Nos. U1832133 and 11905206).
*Corresponding author. Email: libingshengmvp@163.com Citation Text: Liao Q, Kang L, Zhang T M, Liu H P and Wang T et al. 2020 Chin. Phys. Lett. 37 076102    Abstract Cavities and extended defects formed in single crystalline and polycrystalline $\alpha$-SiC implanted with H$^{+}$ ions are compared. The samples are investigated by cross-sectional transmission electron microscopy. H$_{2}$ bubbles are formed during H implantation and H$_{2}$ molecules escape the sample to form cavities during thermal annealing at 1100℃. Microcracks and the extended defects prefer to nucleate in single crystalline $\alpha$-SiC, but not polycrystalline $\alpha$-SiC. Grain boundaries can account for the experimental results. The formation of cavities on grain boundaries is investigated. DOI:10.1088/0256-307X/37/7/076102 PACS:61.80.Jh, 61.82.Fk, 68.37.Lp, 81.40.Wx © 2020 Chinese Physics Society Article Text Silicon carbide (SiC) is an important functional and structural material used in semiconductor devices and in nuclear power plants. For example, hexagonal single-crystalline SiC, such as 4H-SiC and 6H-SiC, has excellent optical and electronic properties, so great efforts have been taken in the development of the potential of SiC acting as a primary candidate material used in high-temperature, high-power, high-frequency and high-voltage switching devices.[1,2] Similarly, polycrystalline SiC, such as hot-pressed SiC, can be regarded as a pressure vessel material for the tristructural-isotropic (TRISO) coated fuel particles for high temperature gas-cooled reactors, due to the very low diffusion rate of fission products in SiC. In addition, SiC has good properties of thermal diffusion, liquid metal corrosion, low neutron capture, as well as a mature industry technology adopted to large area production. SiC/SiC composites can be regarded as a component used in flow channel inserts (FCls), structural material for the first wall in fusion systems, coating of fuel cladding acting as accident tolerant fuel concepts for light water reactors, and a primary candidate material of a pump blade used in an accelerator driven system (ADS) and a lead-cooled fast reactor system.[3] Understanding hydrogen irradiation damage and defect evolution in SiC are very important whatever SiC materials are used in semiconductor industries or in nuclear power plants. For example, SiC-on-insulator (SiCOI) structures have many good properties, i.e., low power dissipation and high resistance to radiation, so it is worth investigating how to fabricate SiCOI structures. An efficient and mature technology is the Smart-cut technology,[4] which consists of three steps: (1) hydrogen (H) ion implantation, (2) wafer bonding, and (3) fracture to achieve thin layer transfer. The critical step is to form platelets through H implantation into single crystalline SiC. After annealing, these platelets can migrate and coalesce into microcracks that are parallel to the sample surface, leading to the surface wafer exfoliation. It is important to choose ideal irradiation energy, dose and temperature to achieve wafer cutting. To date, the implantation damage and platelet nucleation in H-implanted 6H-SiC have been investigated.[5–11] The nucleation and growth of platelets relay on implantation dose and annealing temperature. Our recent results show that the best threshold condition of 134 keV H$_{2}^{+}$ ion-implanted 6H-SiC at room temperature is of $1.5 \times 10^{16}$ H$_{2}^{+}$/cm$^{2}$ after 1100℃ annealing for 15 min.[6] When polycrystalline SiC is used in the field of nuclear power plants, it will suffer from intense neutron irradiation. It has been reported that hydrogen production in SiC can reach 35 appm H/dpa in LiPb-magnetic confinement power plant ($\sim $20 MW$\cdot$y$^{-1}$m$^{-2}$, over 100 dpa per full power year as structural material for the first wall).[12] Energetic H atoms can be formed through nuclear reaction within SiC. These H atoms will accumulate to form platelets, even microcracks, resulting in the degradation of material properties, such as exfoliation and embrittlement along grain boundaries.[13] Therefore, the study of H-implantation-induced defects and cavity distribution in polycrystalline SiC can provide important scientific data. The influence of grain boundaries on cavity nucleation in H-implanted polycrystalline SiC is less studied. How about the cavity growth in the grain boundary plane and the relationship of implantation fluence with the defect sink ability of grain boundaries? These questions are still unanswered. In this letter, we make a comparison study of cavities formed in single-crystalline and polycrystalline $\alpha$-SiC after H implantation followed by thermal annealing. The research is important to understand the role of grain boundaries in the cavity nucleation and defect evolution in grain interiors. The research results will provide some references for material design, in order to improve radiation resistance. In the present study, the $\langle 0001\rangle$-oriented single crystalline 6H-SiC wafers were obtained from the MIT company with dimensions of 10 $\times 10 \times 0.35$ mm$^{3}$. The hot-pressed SiC polycrystallines were supplied by the Saint-Gobain company, with a density of 3.1 g/cm$^{3}$ and grain sizes of 4–10 µm using boron as the sintered aid in the form of B$_{4}$C. SiC polycrystalline is composed of approximately 90% 6H-SiC and 10% 4H-SiC.[14] The sample has dimensions of 15 $\times 15 \times 0.6$ mm$^{3}$. Hydrogen implantation experiments were performed at the 320 kV high-voltage platform equipped with an electron cyclotron resonance (ECR) ion source in Institute of Modern Physics, Chinese Academy of Sciences. Hydrogen molecular beams (H$_{2}^{+}$) with an energy of 134 keV were chosen to implant 6H-SiC samples at room temperature. Three different implantation fluences were set, i.e., $1.5 \times 10^{16}\,{\rm H}_{2}^{+}$/cm$^{2}$, $2.5 \times 10^{16}\,{\rm H}_{2}^{+}$/cm$^{2}$ and $5 \times 10^{16}\,{\rm H}_{2}^{+}$/cm$^{2}$. Hydrogen atom beams (H$^{+}$) with an energy of 120 keV were chosen to implant SiC polycrystalline samples at room temperature. The three different implantation fluences of $3 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$, $5 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$ and $1 \times 10^{17}\,{\rm H}^{+}$/cm$^{2}$ were set. According to Ref. [15], the hydrogen molecular beam can produce the same irradiation effects as the hydrogen atom beams. In order to conveniently compare, two hydrogen atoms instead of one hydrogen molecule are used. The beam was magnetically scanned against a stationary ion beam to provide uniform ion fluence across the sample. The scanned area is $16 \times 17$ mm$^{2}$, which ensures the entire sample surface irradiated. The implantation damage in displacements per atom (dpa) and concentration of 67 keV H$^{+}$ implanted 6H-SiC were simulated by SRIM-2010 with quick cascade calculation,[16] and the man-projected range $R_{\rm p}$ is approximately 393 nm with a straggling $\Delta R_{\rm p}$ of approximately 46 nm. As for 120 keV H$^{+}$ implanted SiC polycrystalline, the $R_{\rm p}$ is approximately 710 nm with $\Delta R_{\rm p}$ of approximately 80 nm. In order to investigate the cavity formation and defect recovery in the damaged layer, isothermal annealing at temperature 1100℃ for 15 min in a vacuum (approximately 10$^{-3}$ Pa) is performed. The majority of the implanted hydrogen is released in the grain after annealing at 1100℃, leaving cavities behind.
cpl-37-7-076102-fig1.png
Fig. 1. Bight-field under-focused XTEM micrographs of cavities formed in 6H-SiC [(a), (b), (c)] and polycrystalline SiC [(d), (e), (f)] with H$^{+}$ ion implantation to fluences of [(a), (d)] $3 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$, [(b), (e)] $5 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$ and [(c), (f)] $1 \times 10^{17}\,{\rm H}^{+}$/cm$^{2}$ at RT followed by thermal annealing at 1100℃ for 15 min. Note that the original implanted surface is located to the left micrograph. Defect-induced lattice strain exhibits black contrast and cavities or microcracks exhibit white contrast. Selected area electron diffraction taken from the zone A noted by a white square and a magnified image taken from the zone B noted by a white square are shown as insets in (c). Nanometric bubbles ($\le $1 nm) exhibiting bright spots are indicated by white arrows.
The microstructural evolution as a function of the implantation fluence was investigated with a Tecnai G20 transmission electron microscope (TEM) equipped with a double-tilt goniometer stage. In order to study the depth distribution of cavities and lattice defects, cross-sectional TEM (XTEM) images were obtained at 200 kV. The fabrication process of XTEM samples was described in our recent papers.[17,18] The cavities were imaged under under-focused and over-focused conditions to highlight the crack edges with Fresnel contrast. The observed extended defects were examined in weak-beam dark-field (WBDF) mode with different ${\boldsymbol g}$ where ${\boldsymbol g}$ is the diffraction vector. Figure 1 presents the under-focused XTEM bright field micrographs of cavities formed in 6H-SiC and SiC polycrystalline. A buried microcrack layer was observed in 6H-SiC implanted with a fluence of $3 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$. In comparison, isolated cavities that have an elliptical shape were observed beneath the sample surface of SiC polycrystalline, as shown in Fig. 1(d). Cavity-induced strain contrast around the cavity can be clearly visible. The cavity preferentially grown in a plane parallel to the sample surface due to the plastic deformation of the subsurface layer is found, such as blisters on the sample surface.[6] With increasing fluence to $5 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$, broken microcracks and dense lattice defects exhibiting black contrast are visible beneath the sample surface of 6H-SiC, compared with faceted cavities that elongate along the $a$ direction parallel to the sample surface formed in polycrystalline SiC. With continual increasing fluence to $1 \times 10^{17}$ H$^{+}$/cm$^{2}$, instead of microcracks, many white spots and lattice defects are visible beneath the sample surface of 6H-SiC. In the middle of the damaged layer, no contrast was observed due to amorphous structure confirmed by selected area electron diffraction, as shown in inset of Fig. 1(c). However, for the polycrystalline SiC implanted with $1 \times 10^{17}$ H$^{+}$/cm$^{2}$, many nanometric spherical cavities are observed in the damaged layer and some large faceted cavities are observed in the middle of the damaged layer, as shown in Fig. 1(f). Compared with 6H-SiC, the number of observed cavities in polycrystalline SiC is significantly larger than that in 6H-SiC. Meanwhile, no amorphous structure was observed in the damaged layer.
cpl-37-7-076102-fig2.png
Fig. 2. WBDF XTEM micrographs showing the extended defects formed in 6H-SiC [(a), (b), (c)] and SiC polycrystalline [(d), (e), (f)] with H$^{+}$ ion implantation to fluences of [(a), (d)] $3 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$, [(b), (e)]$ 5 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$, and [(c), (f)] 1 $\times 10^{17}\,{\rm H}^{+}$/cm$^{2}$ at RT followed by thermal annealing at 1100℃ for 15 min. Diffraction vector was captioned in each image. Defect-induced lattice strain exhibits white contrast and the original implanted surface is located in the left micrograph.
To investigate the extended defects formed in the damaged layer, WBDF images are presented in Fig. 2. For the 6H-SiC implanted with a fluence of $3 \times 10^{16}$ H$^{+}$/cm$^{2}$, it can be seen in Fig. 2(a) that the extended defects exhibiting white contrasts along the microcrack were visible under ${\boldsymbol g} = \bar{2}110$, corresponding to the Burgers vector of $b = \frac{1}{6}\langle 20\bar{2}3\rangle $. The width of the damaged layer ranges from 11 nm to 43 nm. For the SiC polycrystalline implanted with the same fluence, a strong white contrast is seen in the periphery of the cavity, corresponding to strain contrast. With increasing fluence to $5 \times 10^{16}$ H$^{+}$/cm$^{2}$, in the 6H-SiC observed under ${\boldsymbol g} = 0002$, many white dots due to defect clusters and Frank loops are visible. The width of the damaged layer ranges from 71 nm to 85 nm. In comparison, for the SiC polycrystalline observed under ${\boldsymbol g} = 2\bar{1}\bar{1}1$, many white lines parallel to the sample surface are visible. These observed white lines have a Burgers vector of $b = \frac{1}{6}\langle 20\bar{2}3\rangle$. In addition, many white dots that have a Burgers vector of $b = \frac{1}{2}\langle 000\bar{1}\rangle$ were observed. The width of the damaged layer ranges from 24 nm to 44 nm. With continual increasing fluence to $1 \times 10^{17}$ H$^{+}$/cm$^{2}$, in the 6H-SiC observed under ${\boldsymbol g} = 0002$, many white dots due to defect clusters and Frank loops are visible. The observed Frank loops are larger compared to 6H-SiC implanted to a fluence of $5 \times 10^{16}$ H$^{+}$/cm$^{2}$. Meanwhile, the width of the observed damaged layer increased to a range of 151–168 nm. In the middle of the damaged layer, only isolated and slightly white dots are visible, indicating the loss of the long-range order and short-range order kept in the amorphous structure. As for SiC polycrystalline implanted to a fluence of $1 \times 10^{17}$ H$^{+}$/cm$^{2}$, many white dots can be observed under ${\boldsymbol g} = 0\bar{1}12$. These white dots can be regarded as point defects and the extended defects that have Burgers vectors of $b =\frac{1}{6}\langle 20\bar{2}3\rangle$, $\frac{1}{3}\langle 0\bar{1}10\rangle$ and $\frac{1}{2}\langle 0001\rangle$. The width of the damaged layer ranges from 130 nm to 142 nm.
cpl-37-7-076102-fig3.png
Fig. 3. Cavity formation on grain boundaries in SiC polycrystalline implanted with H$^{+}$ ions to fluences of (a) $3 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$, (b) $5 \times 10^{16}\,{\rm H}^{+}$/cm$^{2}$ and (c) $1 \times 10^{17}\,{\rm H}^{+}$/cm$^{2}$ at RT followed by thermal annealing at 1100℃ for 15 min. Note that grain boundaries are tilted in (a) and (b) to show cavity morphology in the grain boundary plane. Cavities were observed under the under-focused condition with $\Delta f = -512$ nm. A grain boundary is strongly inclined to present the morphology of cavities similar to the coral shape indicated by red lines. It indicates a strong wettability of hydrogen deposited in the grain boundary plane.
In comparison with H implanted 6H-SiC, the observed lattice defects and damaged layer are significantly less in the H implanted polycrystalline SiC, in particular for the low and intermediate fluences. It is indicated that grain boundaries play an important role in the evolution of implantation-induced defects. Figure 3 presents the cavity formation on grain boundaries. A triple grain junction was presented in polycrystalline SiC implanted to a fluence of $3 \times 10^{16}$ H$^{+}$/cm$^{2}$. It can be seen that some isolated and faceted cavities were visible on grain boundaries. The formation of faceted cavities demonstrates that cavity growth is rapid as compared to grain interiors, in order to reduce the surface energy of the cavity by diverging from a spherical geometry. Tilting the grain boundary the inner distribution of cavities are visible. It can be seen that the cavity has a coral shape. This result indicates that the complexes of deposited hydrogen and vacancies, similar to water, migrate in the grain boundary plane that has a strong solubility of H atoms. It may be that many available vacancies in the grain boundary plane can account for the observed cavity shape. In addition, the strong solubility of H atoms in the grain boundary plane would contribute to the preferential occurrence of hydrogen embrittlement. The surface free energy decides the morphology of cavities, and the cavities are preferential to grow on a plane with a low surface free energy. It should be noted that there is a triple grain junction and no cavities can be observed on one of grain boundaries. It is demonstrated that the type of grain boundary can affect the cavity nucleation, consistent with our recent theory calculation of different sink strength of grain boundaries for H atoms. Our recent study shows that $\varSigma 3 (2\bar{1}\bar{1})$ and $\varSigma 5(120), (130)$ have strength trapping.[19] In detail, the sink strength is correlated with hydrogen cluster size, grain boundary formation energy, and grain boundary tilt angle. With increasing fluence to $5 \times 10^{16}$ H$^{+}$/cm$^{2}$, two cavity layers were observed and one of cavity layer is located at the end of projected range and the other is located on grain boundary. Cavities located in the grain boundary plane also have a coral shape, and the number density of observed cavities increases significantly. With continually increasing fluence to $1 \times 10^{17}$ H$^{+}$/cm$^{2}$, cavity nucleation on grain boundary is not significant. The present experimental results show that grain boundaries act as strong sinks for mobile H atoms. The decrease in number density of cavity nucleation in grain interior can be explained by the high efficiency of grain boundaries for trapping interstitial H atoms. The width $W$ of the denuded zone along the grain boundary can be expressed as $W^{2}-W_{0}^{2}\approx 2\int_0^{t_{\rm s}} D_{\rm H}[C_{\rm 0H}(t)/C_{\rm H}]dt$, where $W_{0}$ is the width of the denuded zone after cavity nucleation, $D_{\rm H}$ is the H diffusivity, $C_{\rm 0H}$ is the concentration of interstitial H in solution within the grain interior, $C_{\rm H}$ is the concentration of implanted H, and $t_{\rm s}$ is time for the continual increase of fluence.[20] It can be obtained that $W^{2}$ increases with $D_{\rm H}\cdot C_{\rm 0H}$ that is proportional to annealing temperature, and decreases with the implanted H concentration. It is consistent with the experimental results that cavity nucleation is significant after low- and intermediate-fluence implantation, but not at high-fluence implantation. Because grain boundaries acting as defect sinks can efficiently trap implanted H atoms and implantation-induced Frenkel pairs, the residual number of existed H atoms and vacancies that participates in bubble nucleation in grain interior is decreased. As a result, the residual extended defects and cavities after thermal annealing decrease significantly, compared to H implanted 6H-SiC, in particular at low and intermediate fluences. It is well known that single-crystalline SiC can be easily amorphized with H implantation dose up to 0.2 dpa at RT.[6] According to SRIM-2010 simulation, the peak damage level reaches 0.37 dpa for $1 \times 10^{17}$ H$^{+}$/cm$^{2}$ implantation. Our previous research found that the complete recrystallization of amorphous SiC can occur during thermal annealing at 900℃.[21] However, amorphous structure was still visible in the damaged layer of the 6H-SiC sample after the high fluence implantation. It has been reported that the recrystallization of amorphous materials is related to mutual annihilation of point defects and point defect complexes. Muto et al.[22] investigated the chemical bonds in H$^{+}$ implanted 6H-SiC, and they found that hydrogen atoms preferentially react with C atoms to form CH$_{4}$, besides H$_{2}$. CH$_{4}$ and H$_{2}$ are inclined to occupy vacancy sites to form clusters. It is well recognized that gas atoms can stabilize vacancies to decrease migration rate of vacancies. Therefore, the amorphous layer contains a high density of H$_{2}$ and CH$_{4}$ bubbles after annealing. As shown in the inset of Fig. 1(c), some bubbles with sizes less than 1 nm are observed under the under-focused condition. Because grain boundaries trap implanted H atoms and implantation-induced Frenkel pairs to enhance defect recovery in grain interiors, increasing the density of grain boundaries can improve SiC radiation resistance. However, cavities are inclined to nucleate on grain boundaries to deteriorate SiC mechanical properties, such as embrittlement. It is necessary to increase diffusion rate of H atoms towards surface, to decrease bubble nucleation on grain boundaries. Setting the grain boundaries perpendicular to the irradiation surface and fabricating special grain boundaries, such as $\varSigma\,13, 27$, in order to reduce bubble nucleation, would be a good choice for SiC materials in future. Otherwise, setting nanochannels in SiC, similar to the case in He implanted W reported by Qin et al.,[23] can increase H desorption and defect recovery. In conclusion, we have studied the formation of cavities and the extended defects in the H-implanted single-crystalline and polycrystalline $\alpha$-SiC at RT with fluences of $3 \times 10^{16}$, $5 \times 10^{16}$ and $1 \times 10^{17}$ H ions/cm$^{2}$ followed by annealing at 1100℃ for 15 min. A continual microcrack is formed in 6H-SiC implanted to $3 \times 10^{16}$ H ions/cm$^{2}$, while only isolated cavities are formed in poly-crystalline $\alpha$-SiC. With increasing fluence to $5 \times 10^{16}$ H ions/cm$^{2}$, broken microcracks and dense lattice defects are formed in 6H-SiC, while faceted cavities are formed in polycrystalline $\alpha$-SiC. After H$^{+}$ ion implantation to a fluence of $1 \times 10^{17}$ H ions/cm$^{2}$, nanometric cavities and amorphous layers are formed in 6H-SiC, while only nanometric cavities are formed in polycrystalline $\alpha$-SiC. Grain boundaries can trap interstitial H atoms and Frenkel pairs, and trap strength is related to implanted H deposition concentration. Increasing H implanted fluence will decrease the trap strength. The grown cavities in the grain boundary plane exhibit a coral shape, similar to the water mark. It demonstrates the strong wettability of hydrogen in the grain boundary plane. 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