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Recrystallization of He-ion implanted 6H-SiC upon annealing
Nuclear Instruments and Methods in Physics Research B 345 (2015) 53–57
Contents lists available at ScienceDirect
Nuclear Instruments and Methods in Physics Research B journal homepage: www.elsevier.com/locate/nimb
Recrystallization of He-ion implanted 6H-SiC upon annealing B.S. Li a,⇑, Y.Y. Du a,b, Z.G. Wang a a b
Institute of Modern Physics, Chinese Academy of Sciences, Lanzhou, Gansu 730000, China Graduate University of Chinese Academy of Sciences, Beijing 100049, China
a r t i c l e
i n f o
Article history: Received 8 October 2014 Received in revised form 1 December 2014 Accepted 17 December 2014
Keywords: Ion implantation Annealing Recrystallization Bubbles Transmission electron microscopy
a b s t r a c t Solid phase epitaxial growth of amorphous 6H-SiC created by 15 keV He ion implantation to doses of 1.5 1016, 5 1016 and 1 1017 cm2 at room temperature (RT) followed by annealing ranging from 600 °C to 900 °C for 30 min was investigated. The recrystallization process was investigated via crosssectional transmission electron microscopy (XTEM). Recrystallization initially nucleates and grows at the interface between the amorphous layer and 6H-SiC substrate. In the middle of the amorphous layer, recrystallization nucleation is inhibited. Recrystallization rate is related to the implantation-induced damage and concentration of He impurity. The Fourier transformed images denote that the region of recrystallization contains 3C-SiC and 6H-SiC with different crystalline orientations. Besides, for the 1 1017 cm2 implanted sample, partial areas are kept amorphous in the damaged layer. The threshold temperature of full recrystallization of He ion-implantation-induced amorphization in 6H-SiC and the previous observations on other ions implantation, such as Ne, Ar, Xe etc is compared. The possible reasons are discussed. Ó 2014 Elsevier B.V. All rights reserved.
1. Introduction SiC is a wide band-gap semiconductor material with high breakdown voltage, large electron saturation migration rate, small dielectric coefficient and good chemical stability which is of great interest for high-speed communication and high-power control devices [1–3]. For SiC microelectronics devices, ion-implantation is a kind of great ways to dope impurities [4,5]. However, high dose energetic ions could create radiation damages which may cause amorphization at RT [6–9]. Thermally annealing could cause recrystallization of SiC amorphous layers which significantly affects electric and optical properties of SiC material [10–14]. Recrystallization of amorphous SiC occurs in a wide window of annealing temperature. A partial recrystallization of amorphous SiC occurs upon annealing at a low temperature, whereas complete recrystallization of amorphous SiC occurs upon annealing at a high temperature. Implanted impurities would possibly migrate towards the crystalline surface to decrease the concentration of impurities in SiC for doping under a high temperature annealing. Therefore it’s a great importance of studying the complete recrystallization threshold temperature of SiC materials. Recrystallization of SiC amorphous layers upon annealing has been extensively investigated [15–20]. Bae et al. [16] studied ⇑ Corresponding author. Tel.: +86 931 496 9646. E-mail address: [email protected] (B.S. Li). http://dx.doi.org/10.1016/j.nimb.2014.12.049 0168-583X/Ó 2014 Elsevier B.V. All rights reserved.
microstructural evolution of amorphous 6H-SiC created by Xe-ion implantation upon annealing. They found that the recrystallization rate is related to the implantation-induced damage and Xe impurity concentration. Nakamura et al. [17] reported recrystallized orientation caused by 100 keV 2 1015 cm2 Ar ions implantation upon annealing below 1000 °C via (0 0 0 1)- and (1 1 0 0)- oriented 6H-SiC. For (1 1 0 0)-oriented 6H-SiC, a uniform recrystallized rate is 4.4 nm/min at 770 °C; for (0 0 0 1)-oriented 6H-SiC, the recrystallized rate is 0.019 nm/min at 770 °C and an activation energy is 3.4 eV. Heera et al. [15] reported that almost single-crystal 3C-SiC is formed due to recrystallization on the top of the 6H-SiC substrate. Moreover, the complete recrystallization threshold temperature of amorphous SiC materials has also been widely studied [15,18,20–22]. Gorelik et al. [21] reported that the complete recrystallization threshold temperature of keV Ge-ion implantation in 6H-SiC is higher than 1000 °C. Ishimaru et al. [22] reported that the complete recrystallization threshold temperature of MeV Au-ion implantation in 4H-SiC is around 1193 K. Harada et al. [18] reported that the complete recrystallization threshold temperature of MeV Si-ion implantation in 6H-SiC is as low as 1000 °C. He implantation-induced damage in SiC has been extensively investigated because He implantation can introduce nanoscale cavities proposed for the gettering of metallic impurities in electronic devices and for an economical method of preparing SiCOI (silicon carbide-on-insulator) materials by ‘‘smart-cut’’ [23]. However, the complete recrystallization
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threshold temperature of keV He-ion implantation has not been systematically studied. In the present work, we have studied the microstructural evolution of recrystallization process of amorphous layers created by 15 keV He-implanted 6H-SiC to doses of 1.5 1016, 5 1016 and 1 1017 cm2 followed by annealing ranging from 600 °C to 900 °C for 30 min, and discussed the difference between He+implantation and with the previous observation on the Ne+implantation [29]. 2. Experimental procedures 6H-SiC samples oriented h0 0 0 1i surface supplied by the MIT company, were implanted by 15 keV He ions to doses of 1.5 1016, 5 1016 and 1 1017 cm2 at RT. According to the Monte-Carlo code SRIM2008 [24], the implantation doses correspond to the peak damage varying from 0.4 to 2.7 dpa (displacement per atom), and to the peak helium concentration varying from 1.2 to 8 at.%. The implantation experiment was performed in the 320 kV Multi-discipline Research Platform for Highly Charged Ions of the Institute of Modern Physics, Chinese Academy of Sciences (CAS). Post-implantation, wafers were isochronally annealed in a tube furnace ranging from 600 °C to 900 °C for 30 min in vacuum (61 10-3Pa) . Recrystallization of 6H-SiC samples amorphous layers was investigated via TEM using a Tecnai G20 operated at 200 kV and equipped with a double tilt goniometer stage. The helium bubbles were observed by the bright-field through-focal series imaging technique. XTEM samples were prepared by mechanical thinning up to approximately 40 lm in thickness and then ion milling with Ar ions by two steps. At the first step, ion milling energy was 5 kV with a glancing angle of ±5° until optically controlled perforation occurred in the middle of the XTEM sample. At the second step, ion milling energy was 2 kV with a glancing angle of ±3° for 1 h to minimize radiation damage induced by the Ar ions. 3. Results and discussion Fig. 1 shows low-magnification under-focus XTEM bright images of He-implanted 6H-SiC to doses of [(a)–(c)] 1.5 1016, [(d)–(f)] 5 1016 and [(g)–(i)] 1 1017 cm2, [(a), (d) and (g)] asimplantation;; [(b), (e) and (h)] at 800 °C annealing; [(c), (f) and (i)] at 900 °C annealing. The microstructures of He-implanted 6HSiC after 600 °C annealing were not shown here, because there are no apparent microstructural changes as compared to those of 800 °C annealing. In the present study, 1.5 1016, 5 1016 and 1 1017 cm2 are termed as ‘‘low dose’’, ‘‘moderate dose’’ and ‘‘high dose’’, respectively. In the TEM image, the dark contrasts exhibiting dark spots located in the upper and lower interface of the damage band are observed. This dark contrast originates from a high density of stacking faults and lattice strain caused by implantation-induced damage. In the middle of the damage band, no diffraction contrast is observed after annealing temperatures 6800 °C. Besides, the selected-area diffraction patterns taken from the implantation area over 200 nm of the corresponding samples are shown as insets (see Fig. 1). It can be seen that a ring pattern is exhibited in this area. These results suggest that complete amorphization occurred in the middle of the damage band. In the asimplantation, the images show that the thickness of amorphous layers corresponds to 125 nm in Fig. 1(a), 140 nm in Fig. 1(d) and 167 nm in Fig. 1(g). Through under-focus and over-focus transformation, nanoscale helium bubbles with 1–2 nm in diameter are observed in the deeper zone of the damage band for the moderate and high dose implanted samples. After annealing up to 800 °C, no significantly structural evolution is found in these samples. Only a
small zone of recrystallization occurred near the amorphous/crystalline (a/c) interface, and most of the amorphous layer remained. The density and size of helium bubbles observed in the damage band did not change. After annealing at 900 °C, substantial microstructural features changed compared to the cases of the lower annealing temperature. In the insets of Fig. 1(c), (f) and (i), the selected-area diffraction patterns taken from the damage band exhibit strong Bragg reflections, suggesting a full epitaxial recrystallization of the amorphized layer at annealing temperature 900 °C. It should be noted that besides the diffraction spots corresponding to the [11–20] diffraction pattern from 6H-SiC, the extra spots due to the [0 1 1] diffraction pattern from 3C-SiC (marked in the inset, note that the 1 1 1 reflections of 3C-SiC are equivalent to the 10–12 and 0 0 0 6 reflections of 6H-SiC) are superimposed, suggesting that the polycrystalline structures were formed in the damage band. It is consistent with the report of Harada et al. [18]. Besides, helium bubbles grown significantly for the moderate and high dose implanted samples. In parallel, helium bubbles are found in the damage band of the low dose implanted sample. Thermal annealing gives rise to epitaxial recrystallization of the amorphous layer and an epitaxial regrowth starts from the a/c interface rather than the inner of the amorphous layer. The width of the defected crystalline layer from the surface to the upper a/c interface and from the lower a/c interface to the crystalline substrate, and the width of the amorphous layer changed with annealing temperature were analyzed, and the results are given in Table 1. The results suggest that the recrystallization rate increases with increasing annealing temperature, because of the increase in the diffusion rates of implantation-induced defects at high temperature. It should be noticed that the recrystallization rates at the upper and lower a/c interfaces are related to the ion dose. The high dose implantation has a lowest recrystallization rate as compared to the cases of the low and moderate doses implantation. The effects of implantation-induced damage and impurity concentration on solid phase epitaxial growth in amorphous 6H-SiC have been studied, and results of some of previous studies related solid phase epitaxial growth of 6H-SiC are summarized in Table 2, compared to that of the present study. In Table 2, one can see that high implantation-induced damage and high impurity concentration lead to a high threshold temperature of epitaxial recrystallization. In our previous work [25], using Raman spectroscopy, it was shown that the peak intensities of 6H-SiC decreased, while the peak intensities of Si–Si and C–C homonuclear bonds increased with increasing ion dose. The results imply that a high dose implantation-induced atomistic structure has more chemical disorder than that of a low dose implantation, consistent with the report of Bae et al. [16]. Bae et al. [16] mentioned that chemically disordered atomistic structure caused by Xe implantation in 6HSiC would retard the recrystallization. Therefore, the recrystallization rate is lowest for the high dose implantation. Another reason why the recrystallization rate is lowest for the high dose implantation is impurity concentration. Bae et al. [16] investigated the recrystallization proceeds in the Xe-implanted 6H-SiC, and they found that implanted Xe impurity easily accumulated at a/c interface to retard movement of the a/c interface during recrystallization. This phenomenon could hold for the present results, because a high density of helium bubbles was formed near the a/ c interface in the 5 1016 and 1 1017 cm2 implanted samples after a thermal annealing. In particularly, after annealing at 800 °C, more helium bubbles were formed at the a/c interface due to a high diffusion rate of helium atoms at high temperature. The formation mechanism of helium bubbles in 6H-SiC has been well studied [30–33]. The formation of helium bubbles is related to helium concentration in 6H-SiC [31,32]. When the local concentration of helium atoms is over the threshold concentration for forming bubbles, helium bubbles can be formed. A high density
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Fig. 1. Under-focused XTEM bright field micrographs of 15 keV He-implanted 6H-SiC to doses of [(a)–(c)] 1.5 1016, [(d)–(f)] 5 1016 and [(g)–(i)] 1 1017 cm2, [(a), (d) and (g)] as-implantation; [(b), (e) and (h)] at 800 °C; [(c), (f) and (i)] at 900 °C. The sample surfaces were indicated by arrows.
Table 1 The thickness values of the amorphous layers, the upper and lower layers of recrystallization as a function of annealing temperatures ranging from RT to 900 °C. Dose Layers
1.5 1016 cm2
5 1016 cm2
1 1017 cm2
Thickness As-implanted
Amorphous layers Upper layers Lower layers
125 nm 15 nm 12.5 nm
140 nm 3 nm 10 nm
167 nm 0 nm 5.5 nm
600 °C
Amorphous layers Upper layers Lower layers
105 nm 20.5 nm 14 nm
130 nm 15 nm 11 nm
155.5 nm 0 nm 13.5 nm
800 °C
Amorphous layers Upper layers Lower layers
93 nm 32 nm 22.5 nm
101 nm 30 nm 20 nm
140 nm 0 nm 21.5 nm
900 °C
Amorphous layers Upper layers Lower layers
0 nm – –
0 nm – –
0 nm – –
Table 2 The changes of threshold temperature of complete recrystallization of amorphous 6H-SiC with impurity atoms, dpa, and impurity concentration. Impurity atom
Acceleration voltage
Ion fluence (cm2)
dpa
Concentration of impurity atoms (%)
Temperature of recrystallization (°C)
Refs.
N Cr Si Ar Xe Kr Ne Ge Au He
62 keV 260 keV 2 MeV 100 keV 150 keV 5.3 MeV 2.3 MeV 250 keV 10 MeV 15 keV
8.0 1016 1.0 1016 2 1016 2 1015 1 1015 1 1015 3.75 1015 1.5 1016 1 1015 1.5 1016
16 10 6.5 1.5 2.4 1 0.9 24 2.8 0.6
15 1.0 0.8 0.4 0.4 0.02 0.17 1.7 0.025 1.8
1500 1500 1000 800 890 1000 1000 >1000 920 800–900
[26] [26] [27] [28] [16] [29] [29] [21] [22] Present study
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Fig. 2. High resolution XTEM images of recrystallization regions observed in Fig. 1(c), (f) and (i) corresponding to (a)–(c), respectively.
of helium bubbles at the a/c interface indicates a higher concentration of helium atoms at this zone than other zones in the damage band. In Table 2, a comparison of the present result to literature data after Ne implantation with a similar displacement damage level below 1 dpa shows that the threshold temperature of epitaxial recrystallization is lower after He implantation. Motooka et al. [34] investigated structural relaxation processes in amorphous Si and regarded that large vacancy-type defects can retard solid phase epitaxial growth in amorphous Si. They proposed that existence of the inner free surfaces of the large vacancies would stabilize amorphous structure. It is considered that the same mechanism is associated with the reason for the low threshold temperature of epitaxial recrystallization in the present study. According to the elastic collision theory, Ne ion implantation in 6H-SiC produce more vacancies per ion than that of He ion implantation. Moreover, atomic He easily diffuses into vacancies to form bubbles due to the low solubility of He atoms in 6H-SiC. When vacancy-type defects contain He atoms, the diffusion rate of vacancy-type defects would be reduced. Hence, the agglomeration and coalescence of vacancy-type defects would be retarded during a thermal annealing. The present results have shown that helium bubbles have no significant growth after thermal annealing at 6800 °C. Helium bubbles have rapidly growth when the thermal annealing was done at 900 °C. It has been reported that the silicon vacancy in SiC becomes mobile at 800–900 °C and helium atoms dissociate from small vacancy-type defects above 800 °C [35,36]. Helium bubbles rapidly grown via vacancy accumulation during annealing at 900 °C. XTEM image shows that some helium bubbles with diameters above 20 nm were formed in the high dose implanted 6H-SiC after annealing at 900 °C, as shown in Fig. 1(i). Helium bubbles growth via vacancy accumulation and some vacancies recombination with interstitial-type defects lead to the decrease of vacancy concentration. The decrease of vacancy concentration and lack of the inner free surface of the bubbles lead to the rapid solid phase epitaxial growth of the amorphous layer induced by He implanted-6H-SiC after annealing between 800 and 900 °C. It should be noted that the recrystallization rate at the upper a/c interface is faster than that at the lower a/c interface except the
cases of the high dose implantation, because the upper a/c interface is not found in the high dose as-implanted sample. Because recrystallization process is related to mutual annihilation of vacancy-type defects and interstitial-type defects, the recrystallization rate depends on the diffusion rates of defects formed by He implantation. Leclerc et al. [37] investigated strain-induced drift of interstitial atoms in SiC and mentioned that the formation of a thin and deep highly strain region in the near surface region leads to the accumulation of interstitial atoms drifted to the deep highly strained region. The migration of interstitial-type defects into the highly strained region and then the recombination of vacancy-type defects can explain why the recrystallization rate in the upper a/c interface is faster. To elucidate the microstructures of the recrystallization region, high resolution XTEM was employed. Fig. 2(a)–(c) present high resolution XTEM images of the recrystallization regions observed in Fig. 1(c), (f) and (i), respectively. The low dose implanted sample after annealing at 900 °C, lattice fringes are clearly observed in Fig. 2(a). Fourier transformed images taken from A–C zones in the Fig. 2(a) were obtained, as shown in the insets. The pattern taken from the zone A denotes that this zone is 6H-SiC. The pattern is same as that of the substrate, indicating that zone A is epitaxial 6H-SiC layer-by-layer grown from the upper a/c interface. One can see that the fringes indicated by the circle I are wavy, suggesting that this zone contains numerous stacking faults. It is consistent with the diffraction pattern which exhibits diffuse streaks marked by arrows in the inset of Fig. 1(c). These stacking faults located in the basal plane were formed by intrinsic or extrinsic type defects. The pattern taken from the zone B is same as that of A, but its [0 0 0 1] direction is different from the c-axis of the substrate. Columnar growth in different directions was previous reported [15,18]. Zone B also contains numerous stacking faults indicated by the circle II. The pattern taken from the zone C denotes that this zone is 3C-SiC. At the perfect lattice zone, the measured lattice plane spacings of 0.25(3) nm in zone A, 0.25(2) nm in zone B and 0.25(1) nm in zone C referring to the lattice plane spacings of the SiC matrix are in consistent with the theoretical lattice plane spacings of 0.252 nm for both c plane in 6H-SiC and for (0 0 2) plane in 3C-SiC. Fig. 2(b) shows the microstructural changes of
B.S. Li et al. / Nuclear Instruments and Methods in Physics Research B 345 (2015) 53–57
the moderate dose implanted sample after annealing at 900 °C. No apparent structural difference is found between the low dose implanted sample and the moderate dose implanted sample, except some highly disordered areas indicated by the circles in Fig. 2(b). Fig. 2(c) shows the microstructural changes of the high dose implanted sample after annealing at 900 °C. Slight microstructural difference is found between the low dose implanted sample and moderate dose implanted sample. The recrystallized area shown in Fig. 2(c) can be divided into five different zones, named as A–E. Fourier transformed images taken from A to E zones in the Fig. 2(c) were obtained, as shown in the insets. The pattern taken from the zone A denotes that this zone is amorphous. The pattern taken from the zone B denotes that this zone is 6H-SiC. The pattern is same as that of the substrate, indicating that zone B is epitaxial 6H-SiC layer-by-layer grown from the lower a/c interface. The patterns taken from the zone C and E denote that these zones are 6H-SiC, but their [0 0 0 1] directions are different from the c-axis of the substrate. The pattern taken from the zone D denotes that this zone is 3C-SiC. The lattice plane spacings of these regions were measured and the results are same as those of the low dose implanted sample. It is interesting that the angle between the c-axes of the 6H-SiC zones B and C/D shown in Fig. 2(c) is 72°, which is good consistent with previous reports [18,20]. According to the high resolution XTEM images, complete recrystallization occurred in the low dose implanted sample after annealing at 900 °C for 30 min. It needs higher annealing temperatures than 900 °C for complete recrystallization in the amorphous 6H-SiC created by the moderate dose and high dose implantation.
structure, resulting in the increase of the threshold temperature of complete recrystallization. Acknowledgments The work was supported by National Nature Science Foundation of China (Grant Nos. 11005130, 11475229). The authors appreciate the Laboratory of 320 kV High-voltage Platform in Institute of Modern Physics, CAS for helium implantation. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
4. Conclusions We have investigated the recrystallization process of 15 keV He-implanted 6H-SiC to doses of 1.5 1016, 5 1016 and 1 1017 cm2 followed by annealing ranging from 600 °C to 900 °C for 30 min. XTEM was employed to investigate the microstructural evolution upon annealing. The recrystallization rate is related to annealing temperature, implantation-induced damage and He impurity concentration. Recrystallization initially nucleates and grows at the a/c interface, and epitaxial 6H-SiC layer-by-layer grows from the a/c interface. Because the formation of a thin and deep highly strain region in the near surface region leads to the accumulation of interstitial atoms drifted to the deep highly strained region, the recrystallization rate at the upper a/c interface is faster than that at the lower a/c interface. The region of recrystallization contains 3C-SiC and 6H-SiC with different crystalline orientations. Many stacking faults are in these recrystallization zones due to the bond mismatches existing at the interface between the columnar and layered region. Complete recrystallization occurred in the 1.5 1016 cm2 implanted sample after annealing at 900 °C for 30 min. Compared with Ne-ion implantation with the similar implantation-induced damage, large vacancy-type defects could enhance the stability of the amorphous
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[20] [21] [22] [23] [24] [25]
[26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37]
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Nuclear Instruments and Methods in Physics Research B journal homepage: www.elsevier.com/locate/nimb
Recrystallization of He-ion implanted 6H-SiC upon annealing B.S. Li a,⇑, Y.Y. Du a,b, Z.G. Wang a a b
Institute of Modern Physics, Chinese Academy of Sciences, Lanzhou, Gansu 730000, China Graduate University of Chinese Academy of Sciences, Beijing 100049, China
a r t i c l e
i n f o
Article history: Received 8 October 2014 Received in revised form 1 December 2014 Accepted 17 December 2014
Keywords: Ion implantation Annealing Recrystallization Bubbles Transmission electron microscopy
a b s t r a c t Solid phase epitaxial growth of amorphous 6H-SiC created by 15 keV He ion implantation to doses of 1.5 1016, 5 1016 and 1 1017 cm2 at room temperature (RT) followed by annealing ranging from 600 °C to 900 °C for 30 min was investigated. The recrystallization process was investigated via crosssectional transmission electron microscopy (XTEM). Recrystallization initially nucleates and grows at the interface between the amorphous layer and 6H-SiC substrate. In the middle of the amorphous layer, recrystallization nucleation is inhibited. Recrystallization rate is related to the implantation-induced damage and concentration of He impurity. The Fourier transformed images denote that the region of recrystallization contains 3C-SiC and 6H-SiC with different crystalline orientations. Besides, for the 1 1017 cm2 implanted sample, partial areas are kept amorphous in the damaged layer. The threshold temperature of full recrystallization of He ion-implantation-induced amorphization in 6H-SiC and the previous observations on other ions implantation, such as Ne, Ar, Xe etc is compared. The possible reasons are discussed. Ó 2014 Elsevier B.V. All rights reserved.
1. Introduction SiC is a wide band-gap semiconductor material with high breakdown voltage, large electron saturation migration rate, small dielectric coefficient and good chemical stability which is of great interest for high-speed communication and high-power control devices [1–3]. For SiC microelectronics devices, ion-implantation is a kind of great ways to dope impurities [4,5]. However, high dose energetic ions could create radiation damages which may cause amorphization at RT [6–9]. Thermally annealing could cause recrystallization of SiC amorphous layers which significantly affects electric and optical properties of SiC material [10–14]. Recrystallization of amorphous SiC occurs in a wide window of annealing temperature. A partial recrystallization of amorphous SiC occurs upon annealing at a low temperature, whereas complete recrystallization of amorphous SiC occurs upon annealing at a high temperature. Implanted impurities would possibly migrate towards the crystalline surface to decrease the concentration of impurities in SiC for doping under a high temperature annealing. Therefore it’s a great importance of studying the complete recrystallization threshold temperature of SiC materials. Recrystallization of SiC amorphous layers upon annealing has been extensively investigated [15–20]. Bae et al. [16] studied ⇑ Corresponding author. Tel.: +86 931 496 9646. E-mail address: [email protected] (B.S. Li). http://dx.doi.org/10.1016/j.nimb.2014.12.049 0168-583X/Ó 2014 Elsevier B.V. All rights reserved.
microstructural evolution of amorphous 6H-SiC created by Xe-ion implantation upon annealing. They found that the recrystallization rate is related to the implantation-induced damage and Xe impurity concentration. Nakamura et al. [17] reported recrystallized orientation caused by 100 keV 2 1015 cm2 Ar ions implantation upon annealing below 1000 °C via (0 0 0 1)- and (1 1 0 0)- oriented 6H-SiC. For (1 1 0 0)-oriented 6H-SiC, a uniform recrystallized rate is 4.4 nm/min at 770 °C; for (0 0 0 1)-oriented 6H-SiC, the recrystallized rate is 0.019 nm/min at 770 °C and an activation energy is 3.4 eV. Heera et al. [15] reported that almost single-crystal 3C-SiC is formed due to recrystallization on the top of the 6H-SiC substrate. Moreover, the complete recrystallization threshold temperature of amorphous SiC materials has also been widely studied [15,18,20–22]. Gorelik et al. [21] reported that the complete recrystallization threshold temperature of keV Ge-ion implantation in 6H-SiC is higher than 1000 °C. Ishimaru et al. [22] reported that the complete recrystallization threshold temperature of MeV Au-ion implantation in 4H-SiC is around 1193 K. Harada et al. [18] reported that the complete recrystallization threshold temperature of MeV Si-ion implantation in 6H-SiC is as low as 1000 °C. He implantation-induced damage in SiC has been extensively investigated because He implantation can introduce nanoscale cavities proposed for the gettering of metallic impurities in electronic devices and for an economical method of preparing SiCOI (silicon carbide-on-insulator) materials by ‘‘smart-cut’’ [23]. However, the complete recrystallization
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threshold temperature of keV He-ion implantation has not been systematically studied. In the present work, we have studied the microstructural evolution of recrystallization process of amorphous layers created by 15 keV He-implanted 6H-SiC to doses of 1.5 1016, 5 1016 and 1 1017 cm2 followed by annealing ranging from 600 °C to 900 °C for 30 min, and discussed the difference between He+implantation and with the previous observation on the Ne+implantation [29]. 2. Experimental procedures 6H-SiC samples oriented h0 0 0 1i surface supplied by the MIT company, were implanted by 15 keV He ions to doses of 1.5 1016, 5 1016 and 1 1017 cm2 at RT. According to the Monte-Carlo code SRIM2008 [24], the implantation doses correspond to the peak damage varying from 0.4 to 2.7 dpa (displacement per atom), and to the peak helium concentration varying from 1.2 to 8 at.%. The implantation experiment was performed in the 320 kV Multi-discipline Research Platform for Highly Charged Ions of the Institute of Modern Physics, Chinese Academy of Sciences (CAS). Post-implantation, wafers were isochronally annealed in a tube furnace ranging from 600 °C to 900 °C for 30 min in vacuum (61 10-3Pa) . Recrystallization of 6H-SiC samples amorphous layers was investigated via TEM using a Tecnai G20 operated at 200 kV and equipped with a double tilt goniometer stage. The helium bubbles were observed by the bright-field through-focal series imaging technique. XTEM samples were prepared by mechanical thinning up to approximately 40 lm in thickness and then ion milling with Ar ions by two steps. At the first step, ion milling energy was 5 kV with a glancing angle of ±5° until optically controlled perforation occurred in the middle of the XTEM sample. At the second step, ion milling energy was 2 kV with a glancing angle of ±3° for 1 h to minimize radiation damage induced by the Ar ions. 3. Results and discussion Fig. 1 shows low-magnification under-focus XTEM bright images of He-implanted 6H-SiC to doses of [(a)–(c)] 1.5 1016, [(d)–(f)] 5 1016 and [(g)–(i)] 1 1017 cm2, [(a), (d) and (g)] asimplantation;; [(b), (e) and (h)] at 800 °C annealing; [(c), (f) and (i)] at 900 °C annealing. The microstructures of He-implanted 6HSiC after 600 °C annealing were not shown here, because there are no apparent microstructural changes as compared to those of 800 °C annealing. In the present study, 1.5 1016, 5 1016 and 1 1017 cm2 are termed as ‘‘low dose’’, ‘‘moderate dose’’ and ‘‘high dose’’, respectively. In the TEM image, the dark contrasts exhibiting dark spots located in the upper and lower interface of the damage band are observed. This dark contrast originates from a high density of stacking faults and lattice strain caused by implantation-induced damage. In the middle of the damage band, no diffraction contrast is observed after annealing temperatures 6800 °C. Besides, the selected-area diffraction patterns taken from the implantation area over 200 nm of the corresponding samples are shown as insets (see Fig. 1). It can be seen that a ring pattern is exhibited in this area. These results suggest that complete amorphization occurred in the middle of the damage band. In the asimplantation, the images show that the thickness of amorphous layers corresponds to 125 nm in Fig. 1(a), 140 nm in Fig. 1(d) and 167 nm in Fig. 1(g). Through under-focus and over-focus transformation, nanoscale helium bubbles with 1–2 nm in diameter are observed in the deeper zone of the damage band for the moderate and high dose implanted samples. After annealing up to 800 °C, no significantly structural evolution is found in these samples. Only a
small zone of recrystallization occurred near the amorphous/crystalline (a/c) interface, and most of the amorphous layer remained. The density and size of helium bubbles observed in the damage band did not change. After annealing at 900 °C, substantial microstructural features changed compared to the cases of the lower annealing temperature. In the insets of Fig. 1(c), (f) and (i), the selected-area diffraction patterns taken from the damage band exhibit strong Bragg reflections, suggesting a full epitaxial recrystallization of the amorphized layer at annealing temperature 900 °C. It should be noted that besides the diffraction spots corresponding to the [11–20] diffraction pattern from 6H-SiC, the extra spots due to the [0 1 1] diffraction pattern from 3C-SiC (marked in the inset, note that the 1 1 1 reflections of 3C-SiC are equivalent to the 10–12 and 0 0 0 6 reflections of 6H-SiC) are superimposed, suggesting that the polycrystalline structures were formed in the damage band. It is consistent with the report of Harada et al. [18]. Besides, helium bubbles grown significantly for the moderate and high dose implanted samples. In parallel, helium bubbles are found in the damage band of the low dose implanted sample. Thermal annealing gives rise to epitaxial recrystallization of the amorphous layer and an epitaxial regrowth starts from the a/c interface rather than the inner of the amorphous layer. The width of the defected crystalline layer from the surface to the upper a/c interface and from the lower a/c interface to the crystalline substrate, and the width of the amorphous layer changed with annealing temperature were analyzed, and the results are given in Table 1. The results suggest that the recrystallization rate increases with increasing annealing temperature, because of the increase in the diffusion rates of implantation-induced defects at high temperature. It should be noticed that the recrystallization rates at the upper and lower a/c interfaces are related to the ion dose. The high dose implantation has a lowest recrystallization rate as compared to the cases of the low and moderate doses implantation. The effects of implantation-induced damage and impurity concentration on solid phase epitaxial growth in amorphous 6H-SiC have been studied, and results of some of previous studies related solid phase epitaxial growth of 6H-SiC are summarized in Table 2, compared to that of the present study. In Table 2, one can see that high implantation-induced damage and high impurity concentration lead to a high threshold temperature of epitaxial recrystallization. In our previous work [25], using Raman spectroscopy, it was shown that the peak intensities of 6H-SiC decreased, while the peak intensities of Si–Si and C–C homonuclear bonds increased with increasing ion dose. The results imply that a high dose implantation-induced atomistic structure has more chemical disorder than that of a low dose implantation, consistent with the report of Bae et al. [16]. Bae et al. [16] mentioned that chemically disordered atomistic structure caused by Xe implantation in 6HSiC would retard the recrystallization. Therefore, the recrystallization rate is lowest for the high dose implantation. Another reason why the recrystallization rate is lowest for the high dose implantation is impurity concentration. Bae et al. [16] investigated the recrystallization proceeds in the Xe-implanted 6H-SiC, and they found that implanted Xe impurity easily accumulated at a/c interface to retard movement of the a/c interface during recrystallization. This phenomenon could hold for the present results, because a high density of helium bubbles was formed near the a/ c interface in the 5 1016 and 1 1017 cm2 implanted samples after a thermal annealing. In particularly, after annealing at 800 °C, more helium bubbles were formed at the a/c interface due to a high diffusion rate of helium atoms at high temperature. The formation mechanism of helium bubbles in 6H-SiC has been well studied [30–33]. The formation of helium bubbles is related to helium concentration in 6H-SiC [31,32]. When the local concentration of helium atoms is over the threshold concentration for forming bubbles, helium bubbles can be formed. A high density
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Fig. 1. Under-focused XTEM bright field micrographs of 15 keV He-implanted 6H-SiC to doses of [(a)–(c)] 1.5 1016, [(d)–(f)] 5 1016 and [(g)–(i)] 1 1017 cm2, [(a), (d) and (g)] as-implantation; [(b), (e) and (h)] at 800 °C; [(c), (f) and (i)] at 900 °C. The sample surfaces were indicated by arrows.
Table 1 The thickness values of the amorphous layers, the upper and lower layers of recrystallization as a function of annealing temperatures ranging from RT to 900 °C. Dose Layers
1.5 1016 cm2
5 1016 cm2
1 1017 cm2
Thickness As-implanted
Amorphous layers Upper layers Lower layers
125 nm 15 nm 12.5 nm
140 nm 3 nm 10 nm
167 nm 0 nm 5.5 nm
600 °C
Amorphous layers Upper layers Lower layers
105 nm 20.5 nm 14 nm
130 nm 15 nm 11 nm
155.5 nm 0 nm 13.5 nm
800 °C
Amorphous layers Upper layers Lower layers
93 nm 32 nm 22.5 nm
101 nm 30 nm 20 nm
140 nm 0 nm 21.5 nm
900 °C
Amorphous layers Upper layers Lower layers
0 nm – –
0 nm – –
0 nm – –
Table 2 The changes of threshold temperature of complete recrystallization of amorphous 6H-SiC with impurity atoms, dpa, and impurity concentration. Impurity atom
Acceleration voltage
Ion fluence (cm2)
dpa
Concentration of impurity atoms (%)
Temperature of recrystallization (°C)
Refs.
N Cr Si Ar Xe Kr Ne Ge Au He
62 keV 260 keV 2 MeV 100 keV 150 keV 5.3 MeV 2.3 MeV 250 keV 10 MeV 15 keV
8.0 1016 1.0 1016 2 1016 2 1015 1 1015 1 1015 3.75 1015 1.5 1016 1 1015 1.5 1016
16 10 6.5 1.5 2.4 1 0.9 24 2.8 0.6
15 1.0 0.8 0.4 0.4 0.02 0.17 1.7 0.025 1.8
1500 1500 1000 800 890 1000 1000 >1000 920 800–900
[26] [26] [27] [28] [16] [29] [29] [21] [22] Present study
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Fig. 2. High resolution XTEM images of recrystallization regions observed in Fig. 1(c), (f) and (i) corresponding to (a)–(c), respectively.
of helium bubbles at the a/c interface indicates a higher concentration of helium atoms at this zone than other zones in the damage band. In Table 2, a comparison of the present result to literature data after Ne implantation with a similar displacement damage level below 1 dpa shows that the threshold temperature of epitaxial recrystallization is lower after He implantation. Motooka et al. [34] investigated structural relaxation processes in amorphous Si and regarded that large vacancy-type defects can retard solid phase epitaxial growth in amorphous Si. They proposed that existence of the inner free surfaces of the large vacancies would stabilize amorphous structure. It is considered that the same mechanism is associated with the reason for the low threshold temperature of epitaxial recrystallization in the present study. According to the elastic collision theory, Ne ion implantation in 6H-SiC produce more vacancies per ion than that of He ion implantation. Moreover, atomic He easily diffuses into vacancies to form bubbles due to the low solubility of He atoms in 6H-SiC. When vacancy-type defects contain He atoms, the diffusion rate of vacancy-type defects would be reduced. Hence, the agglomeration and coalescence of vacancy-type defects would be retarded during a thermal annealing. The present results have shown that helium bubbles have no significant growth after thermal annealing at 6800 °C. Helium bubbles have rapidly growth when the thermal annealing was done at 900 °C. It has been reported that the silicon vacancy in SiC becomes mobile at 800–900 °C and helium atoms dissociate from small vacancy-type defects above 800 °C [35,36]. Helium bubbles rapidly grown via vacancy accumulation during annealing at 900 °C. XTEM image shows that some helium bubbles with diameters above 20 nm were formed in the high dose implanted 6H-SiC after annealing at 900 °C, as shown in Fig. 1(i). Helium bubbles growth via vacancy accumulation and some vacancies recombination with interstitial-type defects lead to the decrease of vacancy concentration. The decrease of vacancy concentration and lack of the inner free surface of the bubbles lead to the rapid solid phase epitaxial growth of the amorphous layer induced by He implanted-6H-SiC after annealing between 800 and 900 °C. It should be noted that the recrystallization rate at the upper a/c interface is faster than that at the lower a/c interface except the
cases of the high dose implantation, because the upper a/c interface is not found in the high dose as-implanted sample. Because recrystallization process is related to mutual annihilation of vacancy-type defects and interstitial-type defects, the recrystallization rate depends on the diffusion rates of defects formed by He implantation. Leclerc et al. [37] investigated strain-induced drift of interstitial atoms in SiC and mentioned that the formation of a thin and deep highly strain region in the near surface region leads to the accumulation of interstitial atoms drifted to the deep highly strained region. The migration of interstitial-type defects into the highly strained region and then the recombination of vacancy-type defects can explain why the recrystallization rate in the upper a/c interface is faster. To elucidate the microstructures of the recrystallization region, high resolution XTEM was employed. Fig. 2(a)–(c) present high resolution XTEM images of the recrystallization regions observed in Fig. 1(c), (f) and (i), respectively. The low dose implanted sample after annealing at 900 °C, lattice fringes are clearly observed in Fig. 2(a). Fourier transformed images taken from A–C zones in the Fig. 2(a) were obtained, as shown in the insets. The pattern taken from the zone A denotes that this zone is 6H-SiC. The pattern is same as that of the substrate, indicating that zone A is epitaxial 6H-SiC layer-by-layer grown from the upper a/c interface. One can see that the fringes indicated by the circle I are wavy, suggesting that this zone contains numerous stacking faults. It is consistent with the diffraction pattern which exhibits diffuse streaks marked by arrows in the inset of Fig. 1(c). These stacking faults located in the basal plane were formed by intrinsic or extrinsic type defects. The pattern taken from the zone B is same as that of A, but its [0 0 0 1] direction is different from the c-axis of the substrate. Columnar growth in different directions was previous reported [15,18]. Zone B also contains numerous stacking faults indicated by the circle II. The pattern taken from the zone C denotes that this zone is 3C-SiC. At the perfect lattice zone, the measured lattice plane spacings of 0.25(3) nm in zone A, 0.25(2) nm in zone B and 0.25(1) nm in zone C referring to the lattice plane spacings of the SiC matrix are in consistent with the theoretical lattice plane spacings of 0.252 nm for both c plane in 6H-SiC and for (0 0 2) plane in 3C-SiC. Fig. 2(b) shows the microstructural changes of
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the moderate dose implanted sample after annealing at 900 °C. No apparent structural difference is found between the low dose implanted sample and the moderate dose implanted sample, except some highly disordered areas indicated by the circles in Fig. 2(b). Fig. 2(c) shows the microstructural changes of the high dose implanted sample after annealing at 900 °C. Slight microstructural difference is found between the low dose implanted sample and moderate dose implanted sample. The recrystallized area shown in Fig. 2(c) can be divided into five different zones, named as A–E. Fourier transformed images taken from A to E zones in the Fig. 2(c) were obtained, as shown in the insets. The pattern taken from the zone A denotes that this zone is amorphous. The pattern taken from the zone B denotes that this zone is 6H-SiC. The pattern is same as that of the substrate, indicating that zone B is epitaxial 6H-SiC layer-by-layer grown from the lower a/c interface. The patterns taken from the zone C and E denote that these zones are 6H-SiC, but their [0 0 0 1] directions are different from the c-axis of the substrate. The pattern taken from the zone D denotes that this zone is 3C-SiC. The lattice plane spacings of these regions were measured and the results are same as those of the low dose implanted sample. It is interesting that the angle between the c-axes of the 6H-SiC zones B and C/D shown in Fig. 2(c) is 72°, which is good consistent with previous reports [18,20]. According to the high resolution XTEM images, complete recrystallization occurred in the low dose implanted sample after annealing at 900 °C for 30 min. It needs higher annealing temperatures than 900 °C for complete recrystallization in the amorphous 6H-SiC created by the moderate dose and high dose implantation.
structure, resulting in the increase of the threshold temperature of complete recrystallization. Acknowledgments The work was supported by National Nature Science Foundation of China (Grant Nos. 11005130, 11475229). The authors appreciate the Laboratory of 320 kV High-voltage Platform in Institute of Modern Physics, CAS for helium implantation. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
4. Conclusions We have investigated the recrystallization process of 15 keV He-implanted 6H-SiC to doses of 1.5 1016, 5 1016 and 1 1017 cm2 followed by annealing ranging from 600 °C to 900 °C for 30 min. XTEM was employed to investigate the microstructural evolution upon annealing. The recrystallization rate is related to annealing temperature, implantation-induced damage and He impurity concentration. Recrystallization initially nucleates and grows at the a/c interface, and epitaxial 6H-SiC layer-by-layer grows from the a/c interface. Because the formation of a thin and deep highly strain region in the near surface region leads to the accumulation of interstitial atoms drifted to the deep highly strained region, the recrystallization rate at the upper a/c interface is faster than that at the lower a/c interface. The region of recrystallization contains 3C-SiC and 6H-SiC with different crystalline orientations. Many stacking faults are in these recrystallization zones due to the bond mismatches existing at the interface between the columnar and layered region. Complete recrystallization occurred in the 1.5 1016 cm2 implanted sample after annealing at 900 °C for 30 min. Compared with Ne-ion implantation with the similar implantation-induced damage, large vacancy-type defects could enhance the stability of the amorphous
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[26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37]
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