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Low temperature growth of Ge1 − xSnx buffer layers for tensile–strained Ge layers
Thin Solid Films 518 (2010) S2–S5
Contents lists available at ScienceDirect
Thin Solid Films j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / t s f
Low temperature growth of Ge1 − xSnx buffer layers for tensile–strained Ge layers Yosuke Shimura a,⁎, Norimasa Tsutsui a, Osamu Nakatsuka a, Akira Sakai b, Shigeaki Zaima a a b
Department of Crystalline Materials Science, Graduate School of Engineering, Nagoya University, Furo-cho, Chikusa-ku, Nagoya 464-8603, Japan Department of Systems Innovation, Graduate School of Engineering Science, Osaka University, 1-3 Machikaneyama-cho, Toyonaka, Osaka 560-8531, Japan
a r t i c l e
i n f o
Available online 18 October 2009 Keywords: Silicon Germanium Tin Tensile strain Epitaxial growth Point defect
a b s t r a c t We have investigated the dependence of Sn precipitation and crystallinity of Ge1 − xSnx layers on the growth temperature. We also demonstrated a growth of a tensile–strained Ge layer on strain-relaxed Ge1 − xSnx buffer layers. In order to suppress Sn precipitation in Ge1 − xSnx layers and improve the crystalline quality, we strongly suggest that point defects have to be introduced by using low temperature growth MBE. The point defects effectively contribute to the lateral propagation of misfit dislocations at the Ge1 − xSnx/virtual Ge substrate. The point defects would be also effective to stabilize substitutional Sn atoms in Ge1 − xSnx layers due to the formation of Sn-vacancy pairs. As a result, Sn precipitation was suppressed in the Ge1 − xSnx layer by low temperature growth, and we realized the Ge1 − xSnx layer with a Sn content of 7.1%. We also achieved the formation of the Ge layer with a tensile–strain value of 0.71%. © 2009 Elsevier B.V. All rights reserved.
1. Introduction The performance of the conventional Si and strained-Si metaloxide-semiconductor field effect transistors (MOSFETs) can be improved by the shrinking dimension of transistors. However, it has now faced with its physical limit beyond the 32 nm technology node. One effective approach independent of the device scaling is the introduction of new materials and strain in channel region. A tensile– strained Ge is one of the alternative channel materials to replace Si for the application to high speed MOSFETs because of its higher carrier mobility than strained-Si for not only holes but also electrons [1]. In order to obtain a high enough effective electron mobility compared with strained-Si channel MOSFETs, at least, 1.0% biaxial tensile strain to bulk Ge must be induced. Several groups have reported the methods to induce tensile strain into Ge such as the use of the difference of the thermal expansion coefficient between Si and Ge [2], and the use of InxGa1 − xAs buffer layers on a GaAs substrate [3]. By using these methods, the strain value has been reached 0.2% and 1.37%, respectively. In the first case, however, the realizable tensile strain value is limited to 0.25% which is not enough to realize higher effective electron mobility than strainedSi. In the latter case, because Ga and As can diffuse into Ge layer, and they will behave as impurities for Ge as acceptor and donor, respectively; there are some issues for engineering carrier concentration in tensile–strained Ge layer. On the other hand, Ge1 − xSnx is a promising material as a buffer layer for realizing a tensile–strained Ge layer, since Sn is a group IV material and Ge1 − xSnx has a larger lattice
⁎ Corresponding author. Tel.: +81 52 789 3819 ; fax: +81 52 789 2760. E-mail address: [email protected] (Y. Shimura). 0040-6090/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2009.10.044
constant than Ge [4–8]. We have to develop the process technology of a tensile–strained Ge and its buffer layer on the conventional Si platform. Therefore, we have focused strain-relaxed Ge1 − xSnx layers which can be grown on a Si substrate for buffer layers to induce tensile strain to a Ge layer. A lattice constant of the Ge1 − xSnx layer can be controlled by the content of substitutional Sn atoms. Previously, we reported the tensile–strained Ge layer with a strain value of 0.68% on a compositionally step-graded structure consisting of three Ge1 − xSnx buffer layers [5]. Fang et al. also reported the formation of a Ge1 − xSnx layer with a Sn content of 3.5%, and a Ge layer with tensile strain value of 0.43% grown on the Ge1 − xSnx layer by using ultrahigh vacuum CVD (UHV-CVD) [7]. They also reported the formation of a Ge1 − xSnx layer with a Sn content of 5% for application of optical devices [8]. However, the Sn contents of Ge1 − xSnx and the strain values of Ge are insufficient to overcome the effective electron mobility of strainedSi. It is necessary to establish the growth process of Ge1 − xSnx layers with Sn content over 6.8% to induce a tensile strain value of 1.0% to a Ge layer. Ge1 − xSnx layers are compressively strained by a Si substrate just after the growth generally, then, we have to relax the strain of Ge1 − xSnx layers in order to realize a large in-plane lattice constant of Ge1 − xSnx for a large tensile–strain value of a Ge layer. We have reported that the threading dislocations preexisting in a virtual Ge substrate (v-Ge) effectively enhance the strain relaxation by propagation of misfit dislocations [4]. On the other hand, the Sn precipitation from a Ge1 − xSnx, due to the low solid solubility limit of Sn in Ge, reduces the strain of a Ge1 − xSnx layer and leads to decreasing in a lattice constant of a Ge1 − xSnx layer. Therefore, it is very important to enhance the strain relaxation by enhancing lateral propagation of misfit dislocations with suppressing Sn precipitation to realize a Ge1 − xSnx layer with a large lattice constant.
Y. Shimura et al. / Thin Solid Films 518 (2010) S2–S5
S3
Previously, we reported that there is the critical misfit strain for the Sn precipitation [5]. This fact indicates that the Sn precipitation strongly relates with the strain. However, it is difficult to control a behavior of dislocations in case of increasing the number of stacks of Ge1 − xSnx layer [6]. Therefore, we have to increase Sn content in each Ge1 − xSnx layers, rather, we have to fabricate a greatly tensile– strained Ge layer on only one Ge1 − xSnx layer, it is also desired in the viewpoint of simplifying the growth process. We reported that long time annealing at low temperature is effective to realize Ge1 − xSnx layers with high Sn content and degree of strain relaxation (DSR) [6]. In this study, we focus on lowering the growth temperature to control the dislocation behavior and the Sn precipitation during post deposition annealing (PDA). As well known, point defects are induced by lowering the growth temperature. Kasper et al. reported that point defects play a role to form new dislocations which can enhance strain relaxation [9]. However, the effect of the low temperature growth of Ge1 − xSnx layer has not been clear yet. 2. Experimental Ge and Ge1 − xSnx layers were grown by using a molecular beam epitaxy (MBE) system whose base pressure less than 1 × 10− 8 Pa. After a chemical and in situ thermal cleaning performed at 850 °C for a Si(001) substrate, a Ge layer was epitaxially grown on a Si substrate. Then, the Ge layer was, followed by ex situ rapid thermal annealing at 700 °C for 1 min in N2 ambient to fully relax the strain in the Ge layer. We call this structure consisting of a fully-strain-relaxed Ge layer on a Si substrate the “virtual Ge substrate (v-Ge)” [10]. Ge and Sn were deposited on v-Ge by using a Knudsen cell and an arc-plasma gun, respectively, after chemical and in situ thermal cleaning performed at 430 °C for v-Ge. The thickness of the Ge1 − xSnx layer was 100 nm, and the Sn content ranged from 5.3% to 8.1%. Growth temperatures of 100, 150, and 200 °C were chosen to change the density of point defects. The PDA at 500 °C for 60 min in N2 ambient was performed after the Ge1 − xSnx layer growth for the strain relaxation. Four-crystal X-ray diffraction two dimensional reciprocal space mapping (XRD-2DRSM) and cross sectional transmission electron microscopy (XTEM) were used to estimate the content of substitutional Sn atoms in Ge1 − xSnx layers and to characterize the crystallinity and the dislocation structure of epitaxial Ge and Ge1 − xSnx layers. 3. Results and discussion
Fig. 1. The summary of Ge1 − xSnx 224 reciprocal lattice points for Ge1 − xSnx layer grown at 100, 150, and 200 °C. Corresponding symbols to the growth condition are summarized in Table 1.
tionally, these Ge1 − xSnx layers grown at 150 and 100 °C have a potential to induce the tensile strain of 0.86% and 0.82% in a Ge layer, respectively. It is expected that there are two reasons with which Sn precipitation in Ge1 − xSnx can be suppressed by the low temperature growth. One is the enhancement of propagation of dislocation by point defects as mentioned above. As a result, the strain in a Ge1 − xSnx layer is reduced, and then, it is supposed that the Sn precipitation is suppressed. Another is the reduction of the local strain around the Sn atom-vacancy pair generated by a point defect. Sn atoms can be efficient traps for vacancies because it is energetically stable [11]. If there are no point defects, the lattice near a Sn atom is strained compressively from Ge atoms around a Sn atom as shown in Fig. 2(a). In contrast, point defects in a Ge1 − xSnx layer grown at a low temperature effectively contribute to the reduction of the local strain around a Sn atom as shown in Fig. 2(b), and the stabilization of Sn atoms. Therefore, engineering of the local strain in Ge1 − xSnx layers is very important to control the Sn precipitation.
3.1. Dependence of low temperature growth on Sn precipitation 3.2. Crystallinity of Ge1 − xSnx layer As the growth temperature becomes lower, a density of point defect is expected to be higher. The strain relaxation behavior and the Sn content in these samples were compared after PDA. Fig. 1 shows the summary of peak positions of Ge1 − xSnx224 reciprocal lattice points as measured by XRD-2DRSM for Ge1 − xSnx samples grown with various growth temperatures. The symbols and details of the experimental conditions for these samples are summarized in Table 1. In Fig. 1, the vertical and diagonal lines indicate reciprocal lattice peak positions for pseudomorphic and fullystrain-relaxed Ge1 − xSnx layers, respectively, with various Sn contents. Then, we can estimate the Sn content and the DSR of the Ge1 − xSnx layer from the peak position obtained by XRD-2DRSM. For all Ge1 − xSnx layers grown at 200 °C, the substitutional Sn content doesn't exceed 5.5% due to the Sn precipitation after PDA. On the other hand, for Ge1 − xSnx layers grown at 150 and 100 °C, substitutional Sn contents of 6.8% and 7.1%, respectively, can be achieved while the Sn precipitation occurs. This result suggests that there is a limit of the substitutional Sn content in Ge1 − xSnx layers for each growth temperature, and the limit increases with decrease of the growth temperature. In other words, the Sn precipitation is effectively suppressed by lowering the growth temperature. Addi-
We also estimate the full width at half maximum (FWHM) of the Ge1 − xSnx 224 Bragg reflection peak profile along the long axis of the ellipsoidal peak in XRD-2DRSM for the samples listed in Table 1 in order to evaluate the mosaicity in the Ge1 − xSnx layer. Fig. 3 shows the
Table 1 The detail of growth and PDA conditions of prepared Ge1 − xSnx/Ge/Si samples. Sample
Growth temperature (°C)
Sn content (%)
1 2 3
200
5.4 6.3 6.9
4 5 6
150
7.0 7.4 8.1
7 8
100
7.4 8.0
Symbol in Fig. 1
As-grown
After PDA
○ □
♢
● ■ ♦
△ ▽
▲ ▼
Ge1 − xSnx thickness (nm)
Annealing condition
100
500 °C 60 min
S4
Y. Shimura et al. / Thin Solid Films 518 (2010) S2–S5
3.3. Fabrication of tensile–strained Ge layer Finally, we prepared the tensile–strained Ge layer with a thickness of 20 nm on the Ge1 − xSnx layer grown at 150 °C, whose contents of substitutional Sn atoms and the DSR were 6.8% and 87%, respectively (indicated as the rhombus with diagonal stripe in Fig. 1). Fig. 4 shows the XRD-2DRSM result around the Ge 224 reciprocal lattice point for the Ge/Ge1 − xSnx/v-Ge sample. Because of overlapping the diffraction peak attributed to the tensile–strained Ge layer and that of v-Ge, we estimated the in-plane lattice constant of the top Ge layer by the following process: first, we performed the peak deconvolution on the vertical A–A' line drawn to pass the reciprocal lattice peak from the Ge1 − xSnx layer. Then, we defined the B–B' line passing the reciprocal lattice peaks attributed to v-Ge and the tensile–strained Ge layer obtained from the first peak deconvolution, and finally, the peak position of the tensile–strained Ge layer was determined by the peak deconvolution as shown in Fig. 4(b). As a result, the in-plane lattice constant of the tensile– strained Ge layer was estimated to be 0.56978 nm, which is equivalent to an in-plane tensile strain of 0.71% for bulk Ge. This tensile–strain value of Ge layer grown on Si substrate is the largest one in reported values [5].
4. Summary Fig. 2. The schematic diagram of a Sn atom in Ge1 − xSnx; (a) without vacancy, (b) with vacancy.
summary of these FWHMs plotted against the 220 lattice spacing of Ge1 − xSnx layers. The FWHMs for the as-grown samples increase with the lattice spacing. However, they don't depend on the growth temperature. This result suggests that the growth temperature does not influence on the crystallinity of as-grown Ge1 − xSnx layers. On the other hand, the value of the FWHM of all samples after PDA is smaller than that before PDA. This result suggests that propagation of misfit dislocation results in increasing the characteristic domain size in a Ge1 − xSnx layer effectively. Additionally, the FWHM of samples after PDA also depends on the lattice spacing, and it also seems to depend on the growth temperature. Clarifying the detailed dependence of the domain structure related to the strain relaxation on the growth temperature is a future study.
Fig. 3. The summary of the FWHMs of the Ge1 − xSnx 224 Bragg reflection peak profile along the long axis of the ellipsoidal peak estimated from XRD-2DRSM for Ge1 − xSnx layers grown at 100, 150, and 200 °C.
We investigated the dependence of Sn precipitation in Ge1 − xSnx layers on the growth temperature. We found that Sn precipitation was suppressed by lowering the growth temperature. We deduce the reasons why the Sn precipitation was suppressed as follows. One is the reduction of the strain in Ge1 − xSnx layers due to enhancement of lateral propagation of misfit dislocations by the introduction of point defects in the Ge1 − xSnx layers with lowering the growth temperature. Another is the reduction of the local strain around a Sn atom to stabilize with binding a vacancy introduced by lowering the growth temperature. As a result, we realized the Ge1 − xSnx layer with a high substitutional Sn content and the DSR of 7.1% and 79%, respectively.
Fig. 4. XRD-2DRSM result for asymmetric 224 Bragg reflections in the Ge/Ge1 − xSnx/v-Ge sample and diffraction intensity profile; (a) along the A–A' line and (b) B–B' line.
Y. Shimura et al. / Thin Solid Films 518 (2010) S2–S5
Additionally, we prepared a Ge layer with the largest tensile–strain value of 0.71% on the Ge1 − xSnx layer grown at 150 °C. Acknowledgements This work was partly supported by a Grant-in-Aid for Scientific Research on Priority Areas (No. 18063012) from the Ministry of Education, Culture, Sports, Science and Technology in Japan. The authors would like to thank Dr. S. Takeuchi (Nagoya University) for useful discussion. References [1] M.V. Fischetti, S.E. Laux, J. Appl. Phys. 80 (1996) 2234. [2] Y. Ishikawa, K. Wada, D.D. Cannon, J. Liu, H.-C. Luan, L.C. Kimerling, Appl. Phys. Lett. 82 (2003) 2044.
S5
[3] Y. Bai, K.E. Lee, C. Cheng, M.L. Lee, E.A. Fitzgerald, J. Appl. Phys. 104 (2008) 084518. [4] S. Takeuchi, A. Sakai, K. Yamamoto, O. Nakatsuka, M. Ogawa, S. Zaima, Semicond. Sci. Technol. 22 (2007) S231. [5] S. Takeuchi, Y. Shimura, O. Nakatsuka, S. Zaima, M. Ogawa, A. Sakai, Appl. Phys. Lett. 92 (2008) 231916. [6] Y. Shimura, N. Tsutsui, O. Nakatsuka, A. Sakai, S. Zaima, Jpn. J. Appl. Phys. 48 (2009) 04C130. [7] Y.-Y. Fang, J. Tolle, J. Tice, A.V.G. Chizmeshya, J. Kouvetakis, V.R. D'Costa, J. Menéndez, Chem. Mater. 19 (2007) 5910. [8] J. Kouvetakis, J. Tolle, R. Roucka, V.R. D'Costa, Y.-Y. Fang, A.V.G. Chizmeshya, J. Menendez, ECS Trans. 16 (10) (2008) 807. [9] E. Kasper, K. Lyutovich, M. Bauer, M. Oehme, Thin Solid Films 336 (1998) 319. [10] A. Sakai, N. Taoka, O. Nakatsuka, S. Zaima, Y. Yasuda, Appl. Phys. Lett. 86 (2005) 221916. [11] A. Chroneos, Phys. Stat. Sol. (b) 244 (2007) 3206.
Contents lists available at ScienceDirect
Thin Solid Films j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / t s f
Low temperature growth of Ge1 − xSnx buffer layers for tensile–strained Ge layers Yosuke Shimura a,⁎, Norimasa Tsutsui a, Osamu Nakatsuka a, Akira Sakai b, Shigeaki Zaima a a b
Department of Crystalline Materials Science, Graduate School of Engineering, Nagoya University, Furo-cho, Chikusa-ku, Nagoya 464-8603, Japan Department of Systems Innovation, Graduate School of Engineering Science, Osaka University, 1-3 Machikaneyama-cho, Toyonaka, Osaka 560-8531, Japan
a r t i c l e
i n f o
Available online 18 October 2009 Keywords: Silicon Germanium Tin Tensile strain Epitaxial growth Point defect
a b s t r a c t We have investigated the dependence of Sn precipitation and crystallinity of Ge1 − xSnx layers on the growth temperature. We also demonstrated a growth of a tensile–strained Ge layer on strain-relaxed Ge1 − xSnx buffer layers. In order to suppress Sn precipitation in Ge1 − xSnx layers and improve the crystalline quality, we strongly suggest that point defects have to be introduced by using low temperature growth MBE. The point defects effectively contribute to the lateral propagation of misfit dislocations at the Ge1 − xSnx/virtual Ge substrate. The point defects would be also effective to stabilize substitutional Sn atoms in Ge1 − xSnx layers due to the formation of Sn-vacancy pairs. As a result, Sn precipitation was suppressed in the Ge1 − xSnx layer by low temperature growth, and we realized the Ge1 − xSnx layer with a Sn content of 7.1%. We also achieved the formation of the Ge layer with a tensile–strain value of 0.71%. © 2009 Elsevier B.V. All rights reserved.
1. Introduction The performance of the conventional Si and strained-Si metaloxide-semiconductor field effect transistors (MOSFETs) can be improved by the shrinking dimension of transistors. However, it has now faced with its physical limit beyond the 32 nm technology node. One effective approach independent of the device scaling is the introduction of new materials and strain in channel region. A tensile– strained Ge is one of the alternative channel materials to replace Si for the application to high speed MOSFETs because of its higher carrier mobility than strained-Si for not only holes but also electrons [1]. In order to obtain a high enough effective electron mobility compared with strained-Si channel MOSFETs, at least, 1.0% biaxial tensile strain to bulk Ge must be induced. Several groups have reported the methods to induce tensile strain into Ge such as the use of the difference of the thermal expansion coefficient between Si and Ge [2], and the use of InxGa1 − xAs buffer layers on a GaAs substrate [3]. By using these methods, the strain value has been reached 0.2% and 1.37%, respectively. In the first case, however, the realizable tensile strain value is limited to 0.25% which is not enough to realize higher effective electron mobility than strainedSi. In the latter case, because Ga and As can diffuse into Ge layer, and they will behave as impurities for Ge as acceptor and donor, respectively; there are some issues for engineering carrier concentration in tensile–strained Ge layer. On the other hand, Ge1 − xSnx is a promising material as a buffer layer for realizing a tensile–strained Ge layer, since Sn is a group IV material and Ge1 − xSnx has a larger lattice
⁎ Corresponding author. Tel.: +81 52 789 3819 ; fax: +81 52 789 2760. E-mail address: [email protected] (Y. Shimura). 0040-6090/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2009.10.044
constant than Ge [4–8]. We have to develop the process technology of a tensile–strained Ge and its buffer layer on the conventional Si platform. Therefore, we have focused strain-relaxed Ge1 − xSnx layers which can be grown on a Si substrate for buffer layers to induce tensile strain to a Ge layer. A lattice constant of the Ge1 − xSnx layer can be controlled by the content of substitutional Sn atoms. Previously, we reported the tensile–strained Ge layer with a strain value of 0.68% on a compositionally step-graded structure consisting of three Ge1 − xSnx buffer layers [5]. Fang et al. also reported the formation of a Ge1 − xSnx layer with a Sn content of 3.5%, and a Ge layer with tensile strain value of 0.43% grown on the Ge1 − xSnx layer by using ultrahigh vacuum CVD (UHV-CVD) [7]. They also reported the formation of a Ge1 − xSnx layer with a Sn content of 5% for application of optical devices [8]. However, the Sn contents of Ge1 − xSnx and the strain values of Ge are insufficient to overcome the effective electron mobility of strainedSi. It is necessary to establish the growth process of Ge1 − xSnx layers with Sn content over 6.8% to induce a tensile strain value of 1.0% to a Ge layer. Ge1 − xSnx layers are compressively strained by a Si substrate just after the growth generally, then, we have to relax the strain of Ge1 − xSnx layers in order to realize a large in-plane lattice constant of Ge1 − xSnx for a large tensile–strain value of a Ge layer. We have reported that the threading dislocations preexisting in a virtual Ge substrate (v-Ge) effectively enhance the strain relaxation by propagation of misfit dislocations [4]. On the other hand, the Sn precipitation from a Ge1 − xSnx, due to the low solid solubility limit of Sn in Ge, reduces the strain of a Ge1 − xSnx layer and leads to decreasing in a lattice constant of a Ge1 − xSnx layer. Therefore, it is very important to enhance the strain relaxation by enhancing lateral propagation of misfit dislocations with suppressing Sn precipitation to realize a Ge1 − xSnx layer with a large lattice constant.
Y. Shimura et al. / Thin Solid Films 518 (2010) S2–S5
S3
Previously, we reported that there is the critical misfit strain for the Sn precipitation [5]. This fact indicates that the Sn precipitation strongly relates with the strain. However, it is difficult to control a behavior of dislocations in case of increasing the number of stacks of Ge1 − xSnx layer [6]. Therefore, we have to increase Sn content in each Ge1 − xSnx layers, rather, we have to fabricate a greatly tensile– strained Ge layer on only one Ge1 − xSnx layer, it is also desired in the viewpoint of simplifying the growth process. We reported that long time annealing at low temperature is effective to realize Ge1 − xSnx layers with high Sn content and degree of strain relaxation (DSR) [6]. In this study, we focus on lowering the growth temperature to control the dislocation behavior and the Sn precipitation during post deposition annealing (PDA). As well known, point defects are induced by lowering the growth temperature. Kasper et al. reported that point defects play a role to form new dislocations which can enhance strain relaxation [9]. However, the effect of the low temperature growth of Ge1 − xSnx layer has not been clear yet. 2. Experimental Ge and Ge1 − xSnx layers were grown by using a molecular beam epitaxy (MBE) system whose base pressure less than 1 × 10− 8 Pa. After a chemical and in situ thermal cleaning performed at 850 °C for a Si(001) substrate, a Ge layer was epitaxially grown on a Si substrate. Then, the Ge layer was, followed by ex situ rapid thermal annealing at 700 °C for 1 min in N2 ambient to fully relax the strain in the Ge layer. We call this structure consisting of a fully-strain-relaxed Ge layer on a Si substrate the “virtual Ge substrate (v-Ge)” [10]. Ge and Sn were deposited on v-Ge by using a Knudsen cell and an arc-plasma gun, respectively, after chemical and in situ thermal cleaning performed at 430 °C for v-Ge. The thickness of the Ge1 − xSnx layer was 100 nm, and the Sn content ranged from 5.3% to 8.1%. Growth temperatures of 100, 150, and 200 °C were chosen to change the density of point defects. The PDA at 500 °C for 60 min in N2 ambient was performed after the Ge1 − xSnx layer growth for the strain relaxation. Four-crystal X-ray diffraction two dimensional reciprocal space mapping (XRD-2DRSM) and cross sectional transmission electron microscopy (XTEM) were used to estimate the content of substitutional Sn atoms in Ge1 − xSnx layers and to characterize the crystallinity and the dislocation structure of epitaxial Ge and Ge1 − xSnx layers. 3. Results and discussion
Fig. 1. The summary of Ge1 − xSnx 224 reciprocal lattice points for Ge1 − xSnx layer grown at 100, 150, and 200 °C. Corresponding symbols to the growth condition are summarized in Table 1.
tionally, these Ge1 − xSnx layers grown at 150 and 100 °C have a potential to induce the tensile strain of 0.86% and 0.82% in a Ge layer, respectively. It is expected that there are two reasons with which Sn precipitation in Ge1 − xSnx can be suppressed by the low temperature growth. One is the enhancement of propagation of dislocation by point defects as mentioned above. As a result, the strain in a Ge1 − xSnx layer is reduced, and then, it is supposed that the Sn precipitation is suppressed. Another is the reduction of the local strain around the Sn atom-vacancy pair generated by a point defect. Sn atoms can be efficient traps for vacancies because it is energetically stable [11]. If there are no point defects, the lattice near a Sn atom is strained compressively from Ge atoms around a Sn atom as shown in Fig. 2(a). In contrast, point defects in a Ge1 − xSnx layer grown at a low temperature effectively contribute to the reduction of the local strain around a Sn atom as shown in Fig. 2(b), and the stabilization of Sn atoms. Therefore, engineering of the local strain in Ge1 − xSnx layers is very important to control the Sn precipitation.
3.1. Dependence of low temperature growth on Sn precipitation 3.2. Crystallinity of Ge1 − xSnx layer As the growth temperature becomes lower, a density of point defect is expected to be higher. The strain relaxation behavior and the Sn content in these samples were compared after PDA. Fig. 1 shows the summary of peak positions of Ge1 − xSnx224 reciprocal lattice points as measured by XRD-2DRSM for Ge1 − xSnx samples grown with various growth temperatures. The symbols and details of the experimental conditions for these samples are summarized in Table 1. In Fig. 1, the vertical and diagonal lines indicate reciprocal lattice peak positions for pseudomorphic and fullystrain-relaxed Ge1 − xSnx layers, respectively, with various Sn contents. Then, we can estimate the Sn content and the DSR of the Ge1 − xSnx layer from the peak position obtained by XRD-2DRSM. For all Ge1 − xSnx layers grown at 200 °C, the substitutional Sn content doesn't exceed 5.5% due to the Sn precipitation after PDA. On the other hand, for Ge1 − xSnx layers grown at 150 and 100 °C, substitutional Sn contents of 6.8% and 7.1%, respectively, can be achieved while the Sn precipitation occurs. This result suggests that there is a limit of the substitutional Sn content in Ge1 − xSnx layers for each growth temperature, and the limit increases with decrease of the growth temperature. In other words, the Sn precipitation is effectively suppressed by lowering the growth temperature. Addi-
We also estimate the full width at half maximum (FWHM) of the Ge1 − xSnx 224 Bragg reflection peak profile along the long axis of the ellipsoidal peak in XRD-2DRSM for the samples listed in Table 1 in order to evaluate the mosaicity in the Ge1 − xSnx layer. Fig. 3 shows the
Table 1 The detail of growth and PDA conditions of prepared Ge1 − xSnx/Ge/Si samples. Sample
Growth temperature (°C)
Sn content (%)
1 2 3
200
5.4 6.3 6.9
4 5 6
150
7.0 7.4 8.1
7 8
100
7.4 8.0
Symbol in Fig. 1
As-grown
After PDA
○ □
♢
● ■ ♦
△ ▽
▲ ▼
Ge1 − xSnx thickness (nm)
Annealing condition
100
500 °C 60 min
S4
Y. Shimura et al. / Thin Solid Films 518 (2010) S2–S5
3.3. Fabrication of tensile–strained Ge layer Finally, we prepared the tensile–strained Ge layer with a thickness of 20 nm on the Ge1 − xSnx layer grown at 150 °C, whose contents of substitutional Sn atoms and the DSR were 6.8% and 87%, respectively (indicated as the rhombus with diagonal stripe in Fig. 1). Fig. 4 shows the XRD-2DRSM result around the Ge 224 reciprocal lattice point for the Ge/Ge1 − xSnx/v-Ge sample. Because of overlapping the diffraction peak attributed to the tensile–strained Ge layer and that of v-Ge, we estimated the in-plane lattice constant of the top Ge layer by the following process: first, we performed the peak deconvolution on the vertical A–A' line drawn to pass the reciprocal lattice peak from the Ge1 − xSnx layer. Then, we defined the B–B' line passing the reciprocal lattice peaks attributed to v-Ge and the tensile–strained Ge layer obtained from the first peak deconvolution, and finally, the peak position of the tensile–strained Ge layer was determined by the peak deconvolution as shown in Fig. 4(b). As a result, the in-plane lattice constant of the tensile– strained Ge layer was estimated to be 0.56978 nm, which is equivalent to an in-plane tensile strain of 0.71% for bulk Ge. This tensile–strain value of Ge layer grown on Si substrate is the largest one in reported values [5].
4. Summary Fig. 2. The schematic diagram of a Sn atom in Ge1 − xSnx; (a) without vacancy, (b) with vacancy.
summary of these FWHMs plotted against the 220 lattice spacing of Ge1 − xSnx layers. The FWHMs for the as-grown samples increase with the lattice spacing. However, they don't depend on the growth temperature. This result suggests that the growth temperature does not influence on the crystallinity of as-grown Ge1 − xSnx layers. On the other hand, the value of the FWHM of all samples after PDA is smaller than that before PDA. This result suggests that propagation of misfit dislocation results in increasing the characteristic domain size in a Ge1 − xSnx layer effectively. Additionally, the FWHM of samples after PDA also depends on the lattice spacing, and it also seems to depend on the growth temperature. Clarifying the detailed dependence of the domain structure related to the strain relaxation on the growth temperature is a future study.
Fig. 3. The summary of the FWHMs of the Ge1 − xSnx 224 Bragg reflection peak profile along the long axis of the ellipsoidal peak estimated from XRD-2DRSM for Ge1 − xSnx layers grown at 100, 150, and 200 °C.
We investigated the dependence of Sn precipitation in Ge1 − xSnx layers on the growth temperature. We found that Sn precipitation was suppressed by lowering the growth temperature. We deduce the reasons why the Sn precipitation was suppressed as follows. One is the reduction of the strain in Ge1 − xSnx layers due to enhancement of lateral propagation of misfit dislocations by the introduction of point defects in the Ge1 − xSnx layers with lowering the growth temperature. Another is the reduction of the local strain around a Sn atom to stabilize with binding a vacancy introduced by lowering the growth temperature. As a result, we realized the Ge1 − xSnx layer with a high substitutional Sn content and the DSR of 7.1% and 79%, respectively.
Fig. 4. XRD-2DRSM result for asymmetric 224 Bragg reflections in the Ge/Ge1 − xSnx/v-Ge sample and diffraction intensity profile; (a) along the A–A' line and (b) B–B' line.
Y. Shimura et al. / Thin Solid Films 518 (2010) S2–S5
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