Growth and characterization of tensile strained Ge on Ge 1-x Sn x buffers for novel channel layers

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1 The 5th International Symposium on Advanced Science and Technology of Silicon Materials (JSPS Si Symposium), Nov , 2008, Kona, Hawaii, USA Growth and characterization of tensile strained Ge on Ge 1-x Sn x buffers for novel channel layers Akira Sakai 1*, Shotaro Takeuchi 2, Osamu Nakatsuka 2, Shigeaki Zaima 2 1 Graduate School of Engineering Science, Osaka University, 1-3 Machikaneyama-cho, Toyonaka-shi, Osaka , Japan 2 Graduate School of Engineering, Nagoya University, Furo-cho, Chikusa-ku, Nagoya , Japan sakai@ee.es.osaka-u.ac.jp Abstract The application of novel materials, alternative to bulk-si(001) substrates, is now under serious consideration for developing next-generation CMOS channel layers. Implementation of the channel fabrication using new materials other than Si offers several practical issues. One of the most crucial ones is a mismatch of lattice parameters between the new materials and Si. Strain and defects are inevitably introduced into the channel but the device performance can be effectively improved if they are carefully controlled in the materials. In this work, we perform growth and characterization of tensile strained Ge which is promising for obtaining higher carrier mobility compared with strained Si and bulk-ge. Ge 1-x Sn x buffers and virtual Ge substrates are effectively used for inducing tensile strain into the Ge layer. Strain and defect engineering for such group IV semiconductor materials is demonstrated. 1. Introduction Limitation of CMOS device scaling requires several new technologies to gain further device performance. In particular, application of new structures and materials to the CMOS channel, alternative to state-of-art high-quality bulk-si substrates, is a promising solution. Several approaches have been reported so far; use of hybrid Si crystal orientation, introduction of uni- or bi-axial strain, and employment of heterostructures consisting of SiGe, Ge and compound semiconductor. Integration of Ge as a channel material is based on its advantageous potential of higher carrier mobility compared with that for Si. In addition, a recent theoretical prediction has made the in-plane tensile-strained Ge(001) more attractive than bulk-ge; both hole and electron mobilities are significantly increased when the Ge layer is bi-axially tensile-strained [1]. It has previously been reported that tensile-strained Ge can be grown on Si(001) substrates by employing the deference of thermal expansion coefficient between Si and Ge [2-4]. However, achievable in-plain tensile strain by this method is still limited whereas larger strain (more than 1%) is required for greater enhancement of carrier mobility in Ge-MOSFETs. Realization of such large strain in Ge relies on heteroepitaxy technology. A buffer layer formed with precise lattice parameter engineering plays a crucial role in accommodating mismatch strain with respect to Si substrates and inducing global strain into the Ge channel. Ge 1-x Sn x alloy having a larger lattice constant than that of bulk-ge is one of the most attractive candidates for the buffer. Several previous studies have been reported for the growth of Ge 1-x Sn x buffer layers on Si(001) substrates and in-plane tensile-strained Ge layers on them [5,6]. In this paper, we review our recent works [7-9] on the growth of Ge 1-x Sn x buffers on Si substrates and tensile-strained Ge on top of them. Detailed structures grown with such group-iv-material-based heteroepitaxy are characterized. Present affiliation: IMEC, Kapeldreef 75, 3001 Leuven, Belgium

2 2. Experimental Film growth was performed using an ultrahigh vacuum solid-source molecular beam epitaxy (MBE) system with a base pressure less than Torr. We used both bulk-ge(001) and virtual Ge substrates (). For growing a, which is a fully strain-relaxed Ge film on Si(001), a 40-nm-thick Ge layer was deposited at 200 C on a clean surface of a Si(001) substrate, followed by ex situ rapid thermal annealing at 700 C for 1 min in N 2 ambient to relax the strain of Ge. Most of strain in was relaxed by the pure-edge dislocation network formed at the Ge/Si interface [10]. Ge 1-x Sn x layers were then grown by MBE at 200 C on the after cleaning the surface of. Ex situ post deposition annealing (PDA) was performed at 600 C for 10 min in N 2 ambient for strain relaxation of the Ge 1-x Sn x layers. A top Ge layer was epitaxially grown on the Ge 1-x Sn x layer at 250 C. Conventional four-crystal X-ray diffraction two-dimensional reciprocal space mapping (XRD-2DRSM) and cross-sectional transmission electron microscopy (XTEM) were used to characterize strain of the epitaxial Ge and Ge 1-x Sn x layers, and dislocation structures, respectively. 3. Results and discussion 3.1. Growth of Ge 1-x Sn x buffers on bulk and virtual Ge substrates First, we compare structures of Ge 1-x Sn x films grown on bulk-ge and. Figure 1 shows a XTEM image of a 30-nm-thick Ge 0.92 Sn 0.08 layer on bulk-ge. A pseudomorphic Ge 0.92 Sn 0.08 layer is found to be grown without any defects in the layer. On the other hand, as shown in Fig. 1, PDA caused many precipitates observed at around the interface while no misfit dislocations were seen. Figure 1(c) shows a high resolution image of one of precipitates. From this image, we confirmed that the observed precipitates have a -Sn crystalline phase. According to XRD-2DRSM results for the PDA-treated sample, reduction of substitutional Sn concentration from 8% to 2.4% was observed. Degree of strain relaxation (DSR) along [110] direction was also estimated to be 6.0%. These results mean that strain relaxation of Ge 1-x Sn x layers with high Sn contents can not readily be achieved on bulk-ge in our process because of the preferential precipitation of -Sn crystallites and few misfit dislocations introduction at the interface. bulk-ge 20 nm bulk-ge 20 nm (c) 5nm Fig. 1. XTEM images of as-grown and PDA-treated Ge 0.92 Sn 0.08 layers on bulk-ge substrates under the two beam condition of g=004. (c) High-resolution TEM image of a spherical dot observed in, showing a -Sn precipitate. Figure 2 shows a dark-field XTEM image of a 210-nm-thick Ge Sn layer grown on. An epitaxial Ge Sn layer was clearly observed with threading dislocations which come mostly from those preexisting in the through the Ge Sn / interface. Such dislocation morphology was found to change drastically after PDA, as shown in Fig. 2. Looping and lateral propagation of the dislocations

3 frequently occurred at the Ge 1-x Sn x / interface and curved dislocations were formed in the layer. These results indicate that the PDA effectively promotes the movement of threading dislocations, resulting in strain relaxation of the Ge 1-x Sn x layer as well as reduction of threading dislocation density. 100 nm Si 100 nm Si Fig. 2. Dark-filed XTEM images of Ge Sn layers on ; as-grown and PDA-treated samples Tensile Ge on Ge 1-x Sn x buffer / Epitaxial growth of Ge was performed on the Ge 1-x Sn x buffer/ structures. In this experiment, a 20-nm-thick pseudomorphic Ge 0.92 Sn 0.08 layer on was firstly prepared, and then followed by PDA treatment and Ge epitaxy on the top. Figure 3 shows a dark-field XTEM image of a sample having a 20-nm-thick Ge epilayer grown on the PDA-treated Ge 0.92 Sn 0.08 /. Misfit dislocations are clearly observed at the Ge 1-x Sn x / interface (see white arrow) but not seen at the Ge/Ge 1-x Sn x interface, indicating that a pseudomorphic Ge layer can be grown on the strain-relaxed Ge 1-x Sn x layer. -Sn crystallites (see white circle) are also observed at the Ge 1-x Sn x / interface but the number density is considerably lower than that for the bulk-ge case (see Fig. 1). A XRD-2DRSM result of the same sample is shown in Fig. 3 for reciprocal lattice points around Ge 224 and Ge 1-x Sn x 224. Reduction of substitutional Sn concentration in the Ge 1-x Sn x layer was measured to be from 8% to 4.5%. DSR along [110] direction was estimated to be 58%. The peak associated with the Ge epilayer can also be observed in Fig. 3, whose Q x value is almost the same as that of the Ge 1-x Sn x 224 peak, meaning pseudomorphic growth of the Ge layer on the Ge Sn layer. Figure 3(c) shows a deconvolution of the diffraction profile along a dashed line shown in Fig. 3. The Ge epilayer has a (110) lattice spacing of nm, which reflects 0.35% tensile strain with respect to that of bulk Ge. Ge 50 nm Si 7.15 (c) Reciprocal lattice Q y [1/nm] virtual-ge Intensity [a.u.] virtual-ge Fig. 3. Dark-field XTEM image taken under the weak beam condition of g 004 /3g 004 of a Ge epilayer on strain-relaxed Ge 0.92 Sn 0.08 /. XRD-2DRSM result around Ge 224 and Ge 1-x Sn x 224 diffraction spots. (c) Diffraction profile along the dashed line shown in.

4 3.3. Tensile Ge on compositionally step-graded Ge 1-x Sn x buffer / As mentioned above, in the case of direct growth of Ge 1-x Sn x layers with high Sn concentration, -Sn crystallites precipitate markedly due to large misfit strain at the Ge 1-x Sn x / interface. In order to suppress such precipitation at the interface and increase the Sn concentration, we performed growth of compositionally step-graded Ge 1-x Sn x buffers on. In the procedure, a stack of Ge 1-x Sn x layers with different Sn concentrations was grown on and PDA was performed after every growth step of Ge 1-x Sn x layers. Figure 4 shows a XTEM image of a Ge epilayer on step-graded Ge 1-x Sn x layers having a structure of Ge Sn /Ge Sn 0.03 /Ge 0.99 Sn 0.01 /. Misfit dislocations originating from pre-existing threading dislocations and lying at each interface were observed except the Ge epilayer/top-ge 1-x Sn x layer interface. This dislocation morphology surely contributes to strain relaxation of each Ge 1-x Sn x layer. Strain value of the Ge epilayer on step-graded Ge 1-x Sn x layers was derived from XRD-2DRSM for asymmetric 224 reflections, as shown in Fig. 4. Deconvolution of diffraction intensity profiles along the line AA and BB shown in the figure clearly revealed the diffraction peak associated with the Ge epilayer. From the profile of BB through the relaxed-ge peak to the Q y of the Ge epilayer estimated by the profile AA, we obtained the reciprocal lattice point Q x : [1/nm] of the epiayer, which corresponds to in-plane (110) lattice spacing a of nm. For the top-ge 1-x Sn x peak, we obtained a of nm, indicating pseudomorphic growth of the Ge epilayer on the top-ge 1-x Sn x layer. A relaxed Ge lattice constant of nm was obtained from the peak of. In-plane strain in the Ge epilayer is, therefore, estimated to be 0.68±0.03% with the standard error of these reciprocal lattice points for the strained Ge peak in the profile BB. This value exceeds the values obtained by the previous methods [2,6]. In strain relaxation process, threading dislocations are often pinned by some reactions between dislocations during PDA. Such pinned dislocations can neither propagate at the interface nor contribute strain relaxation of the layers. However, in the compositionally step-graded case, the upper layer takes over the pinned threading dislocations in the lower one and the dislocations can move again at the interface during PDA due to misfit stress between the upper and lower layers. This process leads to higher DSR of the stacked Ge 1-x Sn x layers, realizing higher tensile strain in the Ge epilayer. virtual-ge 100 nm Ge top- middle- bottom- Si Reciprocal lattice Q y [1/nm] B A A B top Fig. 4. Dark-field XTEM image of the Ge/Ge Sn /Ge Sn /Ge Sn / sample taken by the weak-beam method at g 004 /3g 004. XRD-2DRSM result of the same sample for asymmetric 224 Bragg reflections. Right (Top) of the figure shows a diffraction intensity profile along the line AA (BB ).

5 4. Conclusion Strain and defect structures have been characterized for the growth of tensile-strained Ge on strain-relaxed Ge 1-x Sn x buffers. Employment of and the compositionally step-graded growth method realize high DSR and high Sn concentration in the Ge 1-x Sn x buffers, both of which allow us to obtain highly tensile-strained Ge epilayer. Structural and morphological modification of dislocations is a key to tailoring novel channel materials for mobility enhancement solutions. Acknowledgments This work was partly supported by a Grant-in-Aid. for Scientific Research on Priority Area (No ) from the Ministry of Education, Culture, Sports, Science, and Technology in Japan. References [1] M. V. Fischetti and S. E. Laux, J. Appl. Phys. 80, (1996), [2] Y. Ishikawa, K. Wada, D. D. Cannon, J. Liu, H. -C Luan and L. C. Kimerling, Appl. Phys. Lett. 82, (2003), [3] D. D. Cannon, J. Liu, Y. Ishikawa, K. Wada, D. T. Danielson, S. Jongthammanurak, J. Michel and L. C. Kimerling, Appl. Phys. Lett. 84, (2004), 906. [4] J. Liu, J. Michel, W. Giziewicz, D. Pan, K. Wada, D. D. Cannon, S. Jongthammanurak, D. T. Danielson, L. C. Kimerling, J. Chen, F. Ö. IIday, F. X. Kärtner and J. Yasaitis, Appl. Phys. Lett. 87, (2005), [5] Y. -Y. Fang, J. Tolle, R. Roucka, A. V. G. Chizmeshya, J. Kouvetakis, V. R. D Costa and J. Menéndez, Appl. Phys. Lett. 90, (2007), [6] Y. -Y. Fang, J. Tolle, J. Tice, A. V. G. Chizmeshya, J. Kouvetakis, V. R. D Costa and J. Menéndez, Chem. Mater. 19, (2007), [7] S. Takeuchi, A. Sakai, K. Yamamoto, O. Nakatsuka, M. Ogawa and S. Zaima, Semicond. Sci. Technol. 22, (2007), S231. [8] S. Takeuchi, Y. Shimura, O. Nakatsuka, S. Zaima, M. Ogawa and A. Sakai, Appl. Phys. Lett. 92, (2008), [9] S. Takeuchi, A. Sakai, O. Nakatsuka, M. Ogawa and S. Zaima, Thin Solid Films, to be published. [10] A. Sakai, N. Taoka, O. Nakatsuka, S. Zaima and Y. Yasuda, Appl. Phys. Lett. 86, (2005),