As one of the most common surface analytical techniques, Rutherford backscattering spectrometry along a low index crystallographic direction or channel (RBS/C) is sensitive to very small atomic displacements from the crystalline lattice sites. The RBS/C interaction yield is strongly influenced by the crystal disorder and crystal orientation, which makes the technique unique for the analysis of ion implantation damage in single-crystal films and wafers. As shown in Fig. 2, A, B, and C represent three relationships between location and the status of damages with the RBS/C from surface to inner region. The energy of backscattering atoms of A have little energy loss and most of ion-beams can be channeled, indicating A is near the surface and has good order lattice. B has an obvious peak, which indicates that many dislocation atoms or disorder lattices exist under the surface; C is the inner part of the crystal, which contains the damage status of C and the effect of B. The high yield in C is mainly due to a mass of dislocation atoms, containing the distortion lattice of C and the little random atoms of B.

Fig. 2.

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	Fig. 2. Relationship between the location and the status of damages with RBS/C.

The X_min refers to the ratio of the channeled intensity to the random intensity in the region just below the surface scattering peak. It yields information about the crystallinity in the near surface of the sample (the first 3 nm–15 nm in depth), which is related to the surface damage, contamination and the existence of oxides. Figure 3 shows the X_min values of samples before and after thermal implantation annealing at various temperatures (1600 °C–1800 °C) for 15 min. Each of all the annealed samples has a smaller X_min value than that of the as-implanted (X_min =0.403) and the X_min value decreases with annealing temperature rising up. With the annealing temperature higher than 1700 °C, the sample’s surface damage can be effectively reduced.

Fig. 3.

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	Fig. 3. Variation of Xmin (160° detector ) with temperature of implanted samples before and after activation annealing for 15 min.

From Figs. 4(a)–4(c), all the samples of SiC samples can be channeled, and the channeling yield increases with channel number decreasing without obvious damage peaks, which do not reach the random level. The crystal has no obvious amorphous layer and the experimentally obtained RBS spectra after annealing are closely identical to the fresh virgin spectrum. It can be concluded that the use of hot temperature (500 °C) implantation can avoid forming the high-density extended defects, dopant redistribution, and surface decomposition in the implanted layer. However, corresponding to the random spectra, a defective layer containing a low level of lattice displacement or distortion is formed on the implanted sample in Fig. 4(a).

Fig. 4.

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	Fig. 4. Overlays of implanted samples (a) before and (b) after activation annealing at various temperatures for 15 min and (c) at 1800 °C and different annealing times.

In the amplified regions on the right in Figs. 4(b) and 4(c), the densities of defects near surface under all the annealing conditions except for the sample of 1600 °C for 15 min, are lower than that of the as-implanted one, which proves that the defects near-surface region are greatly affected by annealing temperature. However, for the annealing time of 15 min, the defects near the bottom of the ion-implanted layer for each of all annealing samples are more than those of as-implanted sample at the different temperatures (from 1600 °C–1800 °C), indicated by the through RBS/C analysis results as shown in Fig. 4(b).

In order to reduce the defects both near the surface and bottom of the ion-implanted layer in high temperature annealing process, the annealing is carried under the conditions of 1800 °C for 5 min and 15 min. It is shown in the amplified region on the right in Fig. 4(c) that the damages near the surface and the bottom of the sample annealed for 1800 °C and 15 min are more pronounced than for 1800 °C and 5 min. Furthermore, as seen in Fig. 4(c), the density of defects near the bottom region under the annealing condition of 1800 °C for 15 min is higher than that of the as-implanted one, but not under the condition of 1800 °C and 5 min. It may be because the longer time thermal annealing induces new extended defects near the bottom of the ion-implanted layers. Therefore, the annealing time is another key factor affecting the generation of extended defects.

In Fig. 5, all samples are tested after removing protective carbon capping layer and sacrificial oxidation. Profiles of implanted atoms before and after annealing at 1800 °C, 15 min, and 5 min by SIMS do not coincide with each other near the bottom of the ion-implanted layer. Possible reasons are that the defect agglomeration and dopant redistribution occur near the bottom of the ion-implanted layer in the annealing process. This is consistent with the conclusion that the defects near the bottom of the ion-implanted layer will be dynamically restored and generated in the activation annealing process. Therefore, it is necessary to optimize the annealing conditions to minimize the defects near the bottom of the ion-implanted layer. As shown in the amplified area of Fig. 5, the curve of the sample obtained under the activation annealing condition (1800 °C/5 min) is highly coincident with that near the bottom of the as-implanted sample.

Fig. 5.

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	Fig. 5. Profiles of implanted atoms before and after annealing at 1800 °C, 15 min/5 min by SIMS after removing protective carbon capping layer and 60-nm-thick sacrificial oxidation SiO2 layer.

From the above results of the RBS-C technique, we find quite a low level of the residual disorders and extended defects in the implanted and annealed SiC layers. As a complement to RBS/C, cross section samples in the implanted region are prepared. And the lattice structure of the implanted layer and the annealed layerare also measured with TEM, which is equipped with an FEI Talos F200 transmission electron microscope at 200 kV. The samples for TEM are cross-sectioned, prepared by focused ion beam (FIB).

Figure 6(a) and 6(b) show cross-sectional bright field-electron (BF) micrographs of identically Al-implanted samples with the annealing temperature at 1800 °C for 15 min and 5 min, respectively. As seen from Figs. 6(a) and 6(b), the visible depths of the damage areas are 445 nm and 467 nm respectively. From the above SIMS test data, it can be inferred that the damage areas include ion implantation region and part of trailing region. In Fig. 6, all of the extended defects are visible in the Al-implanted region after 1800 °C/5 min annealing and 1800 °C/15 min annealing, especially near the bottom of Al-implanted region. These defects are lettered from A to C, which are polygonal loops differing in size. These loops form a network of almost orthogonal trace. Defect A is smallest, next is defect B. In the case of defect C, it is possible to demonstrate that it has a complex structure related to dislocations or carbon vacancy agglomerates.^[19–24] The distribution of these defects is not uniform and appears to peak in a buried region close to the end of the Al-implanted region, which is between the bottom of implanted Al region and the unimplanted substrate. The defects of 1800 °C/15 min sample (Fig. 6(a)) are obviously more than those of the sample with 1800 °C/5 min (Fig. 6(b)). This indicates that the longer annealing time will lead to more extended defects at the same temperature, which is in good agreement with the RBS analysis results. As shown in the amplified defects area in Figs. 7(b) and 7(c), there are many polygonal dislocation loops with different sizes. The micrograph indicates that most of the defects at the bottom of implanted Al layers created during annealing are dislocations as indicated by arrows in Fig. 7(c). The formation of these dislocations may be related to a cluster of C-related defects, such as carbon vacancies.^[19–26]

Fig. 6.

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	Fig. 6. Cross-sectional micrograph in Al-implanted region by (a) 1800 °C/5 min annealing and (b) 1800 °C/15 min annealing.

Fig. 7.

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	Fig. 7. Cross-sectional micrograph in Al-implanted region by 1800 °C/5 min annealing.

To obtain the mechanisms responsible for the formation and elimination of extended defects, according to the previous defect theory,^[15,21–23] the defect model in Al-implanted 4H-SiC layer is also investigated. The high temperature activation annealing is the process of replacing the implanted Al atoms with Si. During the activation annealing, most of dopants are activated, and the ion-induced damages are recovered because of the recrystallization of the annealed samples. As seen from Fig. 8, in the high temperature activation annealing process, the persistent carbon vacancies in the drift layers after epitaxy can be effectively recombined by the excess carbon interstitials C_i created by implation. At the same time, new carbon vacancies will be formed in the high temperature process. So, the optimization of annealing conditions can be realized by controlling the generation and elimination of carbon vacancy (V_C) density in the ion implanted layers.

Fig. 8.

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	Fig. 8. Defect model in Al-implanted 4H-SiC layer.

As seen from Table 1, the Hall hole density P, Hall mobility μ and electrical activation ratio A ( )are improved by increasing annealing temperature through using Pensl et al.’s Hall scattering factor. Sheet resistance R and specific contact resistance (SCR) decrease with annealing temperature increasing. As the temperature increases, the rate of annealing activation increases. At the same annealing temperature, the longer time annealing can improve the electrical activation properties effectively. However, the consideration of the lower defects, the sample annealed at 1800 °C/5 min shows relatively good electrical properties.

Table 1.

Table 1.

Table 1. Values obtained by Hall-effect and TLM measurements for p-type 4H-SiC samples at room temperature (RT). . Sample-post implantation annealing/(°C/min) Measured (SIMS) implanted Al density Ct/1019 cm−3 Sheet resistance R/(k ) RT Hall hole density P/1017 cm−3 Hall mobility μ/(m2/ ) Electrical activation ratio/A% Specific contact resistance (SCR)/(Ω/cm2) 1600/15 2 47.80 1.38 23.7 0.69 4.06×10−4 1650/15 2 34.71 1.81 24.8 0.91 3.15×10−4 1700/15 2 25.62 2.29 26.6 1.15 8.50×10−5 1800/15 2 15.62 3.15 31.3 1.58 6.46×10−6 1800/5 2 19.90 2.8 28.1 1.40 6.59×10−5

Table 1. Values obtained by Hall-effect and TLM measurements for p-type 4H-SiC samples at room temperature (RT). .

Furthermore, the surface roughness values of Al-implanted samples before and after activation annealing at various temperatures are measured by AFM. We use the 5-point average of the large area implantation and activation annealing sample. The results in Table 2 show that the average surface roughness values increase slightly from 0.058 nm to 0.110 nm after Al hot implantation, the surface deterioration may be due to the Al implantation. After annealing at 1600 °C–1800 °C for 15 min, the discontinuity of surface is not obviously raised, which is below 0.304 nm. The surface step-bunching is effectively suppressed and the surface roughness values of all the samples are not obviously deteriorated, which is due to the existence of an effective protective carbon capping layer.

Table 2.

Table 2.

Table 2. Average surface roughness values of Al-implanted 4H-SiC followed by annealing at various temperatures. . Condition Non-implanted As-implanted Annealing temperature/(°C) 1600 1650 1700 1800 RMS roughness/nm 0.058 0.110 0.177 0.270 0.227 0.304

Table 2. Average surface roughness values of Al-implanted 4H-SiC followed by annealing at various temperatures. .