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Germanium growth on Br-terminated Si(1 0 0)
Surface Science 600 (2006) 2907–2912 www.elsevier.com/locate/susc
Germanium growth on Br-terminated Si(1 0 0) B.R. Trenhaile, G.J. Xu 1, J.H. Weaver
*
Department of Physics, Department of Materials Science and Engineering, and Frederick Seitz Materials Research Laboratory, University of Illinois at Urbana-Champaign, Urbana, IL 61801, USA Received 15 December 2005; accepted for publication 16 May 2006 Available online 13 June 2006
Abstract The consequences of Ge deposition on Br-terminated Si(1 0 0) were studied with scanning tunneling microscopy at ambient temperature after annealing at 650 K. One monolayer of Br was sufficient to prevent the formation of Ge huts beyond the critical thickness of 3 ML. This is possible because Br acts as a surfactant whose presence lowered the diffusivity of Ge adatoms. Hindered mobility was manifest at low coverage through the formation of short Ge chains. Further deposition resulted in the extension and connection of the Ge chains and gave rise to the buildup of incomplete layers. The deposition of 7 ML of Ge resulted in a rough surface characterized by irregularly shaped clusters. A short 800 K anneal desorbed the Br and allowed Ge atoms to reorganize into the more energetically favorable ‘‘hut’’ structures produced by conventional Ge overlayer growth on Si(1 0 0). 2006 Elsevier B.V. All rights reserved. Keywords: Scanning tunneling microscopy; Germanium; Si(1 0 0); Epitaxy; Surfactant
1. Introduction The general picture of Ge growth on Si(1 0 0)-(2 · 1) at temperatures above 575 K and deposition rates less than 5 ML/min is that it proceeds via the Stranski–Krastanov mode with a wetting layer of 3 ML followed by threedimensional (3D) islanding [1]. Typically, the wetting layer undergoes a (2 · N) reconstruction where every Nth dimer is missing, and this periodic array of vacancy line defects relieves the misfit strain by allowing the Ge dimers near the trenches to relax [1,2]. The first 3D structures to form on the wetting layer are {5 0 1} faceted ‘‘hut’’ clusters, which have either rectangular or square bases [1,3,4]. Depending upon the growth temperature, the huts either remain to cover the surface or they give way to multifaceted domes or macroscopic clusters after 6–7 ML of Ge have been deposited [4–6]. *
Corresponding author. Tel.: +1 217 333 1440; fax: +1 217 333 2737. E-mail address: [email protected] (J.H. Weaver). 1 Present Address: GE Advanced Materials, Quartz, 24400 Highland Road, Richmond Heights, OH 44143, USA. 0039-6028/$ - see front matter 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.susc.2006.05.035
The growth of thin Ge films on Si(1 0 0) provides a model system for studying strained semiconductor heteroepitaxy as well as offering important device applications for the microelectronics industry. It is well known that growth modes can be manipulated by changing the surface energetics and lowering adatom diffusivity through the use of surfactants. Early studies used As, Sb, and Bi surfactants for Ge on Si(1 0 0), and they revealed that layer-bylayer growth continued well beyond the 3 ML observed on clean Si [7–11]. Surfactants have also been shown to generate more abrupt Ge–Si interfaces by suppressing atomic intermixing of Ge and Si [11–13]. Unfortunately, these surfactants also dope the growing layer. This side effect led to studies of Ge on hydrogen terminated surfaces where the transition to 3D growth can be delayed by a hydrogen flux applied during Ge deposition [14–16]. In this paper, we consider how an adsorbed monolayer of Br affects growth of Ge on Si(1 0 0). Using scanning tunneling microscopy (STM), we show that Br blocks the diffusion of the Ge adatoms during room-temperature deposition. A 650 K anneal of overlayers of nominal thicknesses through 3 ML allowed the random formation
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of short dimer chains. These chains expanded upon further deposition to produce a rugged Ge surface after the deposition of 7 ML. The removal of Br via GeBr2 desorption at 800 K allowed the formation of Ge hut clusters, and annealing at 850 K for 12 h resulted in the coarsening of the hut clusters. The results demonstrate that an adsorbed monolayer of Br acts as a surfactant during Ge growth on Si(1 0 0) by suppressing adatom diffusion and changing surface energetics, without introducing unwanted doping. 2. Experimental considerations The experiments were carried out in a vacuum system with operating pressure <4 · 1011 Torr. Imaging of clean and Ge- and Br-covered surfaces was done at ambient temperature using an Omicron variable temperature STM. The Si wafers were p-type, B-doped to 0.01–0.012 Xcm, and oriented within 0.5 of (1 0 0). Clean surfaces were prepared by degassing at 875 K for 12 h and then heating to 1475 K for 90 s at <5·1010 Torr. An optical pyrometer was used to monitor the sample temperature. For the clean surface, the defect area was 0.01–0.02 ML, primarily in the form of C-type defects and isolated dimer vacancies. A solidstate electrochemical cell derived from AgBr doped with 5 wt.% CdBr2 was used as a Br2 source. The clean Si surface was exposed to a flux of Br2 at room temperature to saturate the dangling bonds. Ge was deposited using an Omicron evaporator that contained a tungsten crucible heated by electron bombardment. All depositions occurred at room temperature with a rate of 9 · 104 ML/s, and the Ge overlayer was built up incrementally. After each deposition, the sample was annealed to 650 K for 30 min and then imaged at room temperature. At the beginning of each cycle, the surface was exposed to an additional Br2 fluence to replenish any Br lost during the short anneal. Cycles of Br passivation, Ge deposition, mild annealing, and room temperature imaging were repeated until a total of 7 ML of Ge had been deposited. A deposition of 0.12 ML, followed by one of 0.5 ML and one of 0.4 ML were performed to reach 1 ML. Thereafter, the remaining overlayer was built up by six depositions of 1 ML each. The sample was then heated to 800 K for 30 min to remove the Br and allow Ge redistribution to form 3D islands. Finally, the sample was annealed at 850 K for 12 h. 3. Results and Discussion STM images revealed that the deposition of 0.12 ML of Ge on Br saturated Si(1 0 0) at room temperature resulted in a poorly ordered surface consisting of mostly isolated adatoms. We speculate that arriving Ge atoms exchanged with Br at the site of arrival. This site exchange has been calculated to occur spontaneously, and it has been observed in studies of Ge and Si deposition on H–Si(1 0 0) [16–19]. The Ge atoms were immobile during subsequent room-temperature imaging because of Br site blocking.
Therefore, the presence of Br lowered the diffusivity of Ge adatoms with respect to the clean surface. Fig. 1(a) shows an STM image of the same surface after annealing to 650 K for 30 min. The dimer rows of the Br terminated Si (layer 0) run diagonally from the upper right to lower left. The post deposition anneal led to a more ordered surface as the Ge adatoms in layer 1 formed dimer chains that are perpendicular to the Si dimer rows. These chains show that the mild anneal has promoted adatom diffusion. For Ge deposition on clean Si(1 0 0), entropy considerations drive Ge atoms to interchange with Si atoms [20–23]. While adsorbed Br should increase the effective exchange barrier, it is possible that the dimer chains observed here include Si atoms. Since the mild anneal promoted Ge diffusion, it is important to discuss atomic scale mechanisms for adatom migration in the presence of surfactants. Kandel and Kaxiras used first-principles calculations to study surfactant mediated growth of semiconductor films, and they developed a model called diffusion-de-exchange-passivation (DDP) [24,25]. Rather than explaining the change in growth mode through reduced diffusion length for the adatoms, the DDP model considers the competition between the diffusion of adatoms on top of the surfactant layer, exchange/de-exchange of adatoms with the surfactant layer, and the attachment of adatoms to steps that are passivated with the surfactant. The DDP model has been successfully applied to semiconductor growth in the presence of groupIII and group-V surfactants. The adsorption and diffusion of Si adatoms on H-terminated Si(1 0 0) surfaces has been studied using first-principles calculations [18,26,19,27]. Those studies found that Si adatoms migrate on the hydrogenated surface through a H atom capture and release mechanism that is assisted by H atom mobility at elevated temperatures. This process gives a much higher barrier for Si adatom diffusion compared to clean Si(1 0 0). Such a capture and release diffusion process should also be applicable to Ge adatoms. The reduced diffusion length in the presence of H has been used to explain the experimental observation that H prevents 3D growth of Ge on Si(1 0 0) [14,16]. Bromine is chemically more similar to H than to the group-III and group-V atoms since it can only form one strong bond. Hence, the capture and release model of diffusion seems more relevant than the DDP model for Br-terminated Si(1 0 0). We also note that Br desorption from Si(1 0 0) is negligible at 650 K (0.005 ML of Br desorb for a 30 min anneal) and it is unlikely that desorption plays a significant role in the Ge mobility [28]. While the attachment of adatoms to steps in the presence of Br is certainly a critical part of the growth process, we take the restricted mobility of Ge adatoms to be the main effect of the Br layer. However, further investigations are needed to fully determine the microscopic mechanisms that occur during Br mediated growth. Fig. 1(b) shows that layer 1 was highly irregular after an additional 0.5 ML deposition and anneal. A type A step of
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Fig. 1. STM micrographs of Ge growth on Br–Si(1 0 0) (2.0 V sample bias). Ge was deposited at room temperature: (a) 0.12 ML of Ge imaged after 600 K anneal for 30 min, 15 · 15 nm2; (b) 0.62 ML of Ge imaged after an additional 0.5 ML Ge deposition followed by a 650 K anneal for 30 min, 25 · 25 nm2. A type-A step of the Si(1 0 0) surface is indicated by the dashed line; (c) 1 ML of Ge imaged after an additional 0.4 ML Ge was deposited followed by a 650 K anneal for 30 min, 30 · 30 nm2. A type-A and a type-B step of the Si(1 0 0) surface are indicated by the dashed lines and (d) 2 ML of Ge imaged after an additional 1 ML Ge deposition followed by an anneal to 650 K for 30 min, 30 · 30 nm2. The surface was passivated with Br before each Ge deposition.
the Si surface runs from the lower left to the upper right corner, as indicated by the dotted line. At this coverage, the dimer chains have connected to form features several dimer rows in width. Since features of this sort did not form at lower coverage, their development is most probably due to the connection of elongating dimer chains. The brightest features in the image reflect Ge adatoms that are beginning the formation of layer 2, and only about 0.03 ML of the Ge adatoms occupy layer 2. These chains appear to have formed at anti-phase boundaries in layer 1. Those boundaries form at the intersection of dimer chains that are out of phase by one atomic unit, and previously studies of Si growth on Br–Si(1 0 0) showed them to be chain nucleation sites [29]. Fig. 1(c) shows the morphology after a total of 1 ML Ge deposition (0.12 + 0.5 + 0.4 ML). Three layers are still
seen, but the dimer chains of layer 2 have increased in number and length to account for about 0.14 ML and layer 1 is composed of 0.87 ML. These chains make it more difficult to distinguish the original Si(1 0 0) steps, but the locations of two steps are indicated by the dotted lines. Fig. 1(d) shows the morphology after a total of 2 ML deposition. There are no substrate steps in this image, and four Ge layers can be seen. Layer 2 covers 0.85 ML, layer 3 accounts for 0.23 ML, and layer 4 is less than 0.01 ML Ge so that 0.09 ML of layer 1 remains incomplete. The limited diffusivity of the Ge adatoms on the Br terminated surface prevented the formation of a 2 · N reconstruction or large islands on the surfaces of Fig. 1 but analysis shows that there is interlayer diffusion. Assuming random adatom arrival, deposition onto the surface represented by Fig. 1(b) would result in a 59% probability of
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arriving on layer 1. Therefore, if diffusing atoms could not overcome the Ehrlich–Schwoebel barrier, the deposition of
an additional 0.4 ML of Ge would have resulted in a minimum of 0.27 ML of Ge in layer 2. Fig. 1(c) reveals that
Fig. 2. STM micrographs of 7 ML Ge grown on Br–Si(1 0 0) (2.0 V sample bias): (a) a rough film was produced after an anneal to 650 K for 30 min, but no Ge hut clusters were observed, 200 · 200 nm2; (b) higher resolution image showing the dimer rows on top of the irregular-shaped clusters, 35 · 35 nm2; (c) annealing to 800 K for 30 min desorbed the Br as GeBr2 and allowed square and rectangular {5 0 1} faceted huts to form, 200 · 200 nm2; (d) higher resolution image of a hut cluster on the 2 · N wetting layer, 35 · 35 nm2 and (e) shows the hut cluster with 3D rendering to emphasize the {5 0 1} facets. Line profiles across the top of such huts shows them to be fairly flat for 5 nm.
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only 0.14 ML occupancy of layer 2 and more than 0.13 ML moved from layer 2 to layer 1 during the anneal. Similar analysis of the surfaces represented by Fig. 1(b) and 1(c) show atom transfer to lower layers. We assume that the atoms staying on the layer of arrival encountered an antiphase boundary or another adatom to nucleate chains. The composition of the layers as they are built up results from a complex interplay between adatom diffusivity in the presence of Br and the number of atoms deposited between each annealing cycle. Fig. 2(a) shows the morphology after a total deposition of 7 ML of Ge. The substrate steps can no longer be seen and the surface is very irregular because of incomplete layer filling. Instead of huts bounded by {5 0 1} facets representative of Ge growth on clean Si(1 0 0), we observe clusters with shallower sides that do not appear to be bounded by a particular family of planes. Fig. 2(b) shows that the clusters are not amorphous because the top layer is composed of dimer-like rows that run in the [0 1 1] and ½011 directions, as on a (1 0 0) surface. Most of the dimer rows on top of the clusters are separated by a missing dimer row, and the lower layers on many clusters can be observed sticking out with dimer rows oriented perpendicular or parallel to the dimer rows of the top layer. The average angle of incline for the sides of the clusters is 8.0 with a range of 6.7–10.0, which is a slightly gentler slope then the 11 observed for {5 0 1} facets. The clusters range in height up to ˚ with an average of 7.3 A ˚. about 11 A We attribute these clusters to decreased adatom diffusivities and altered surface energetics associated with the high concentration of Br. The combination of limited mobility and decreased surface energies prevented the adatoms from reaching the conventional configuration of hut clusters on a 2 · N wetting layer. An increase in surface roughness has also been observed for Si growth on Br terminated Si(1 0 0) [29], and it was manifested here through the irregularly shaped clusters of varying heights. To determine if these clusters would transform into huts if the Br was removed, we annealed the sample at 800 K for 30 min. This should allow ample time for the removal of Br through GeBr2 desorption [30]. Fig. 2(c) reveals a morphology, that is, characteristic of Ge growth on clean Si(1 0 0) at high temperature, with huts on a Ge wetting layer. For comparison with Fig. 2(b), Fig. 2(d) shows a high resolution image of hut cluster on the 2 · N wetting layer. Unlike the clusters of Fig. 2(c), the huts are bounded by {5 0 1} facets as can be clearly seen in the 3D image of Fig. 2(e). As expected, extended annealing at 850 K allowed the larger huts to grow at the expense of smaller ones due to Ostwald ripening as has been observed in real time by several investigators [5,6,31]. 4. Conclusions We have studied the epitaxial growth of Ge on Br terminated Si(1 0 0). Br frustrates Ge adatom diffusion and isolated adatoms are observed after room temperature
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deposition. A mild anneal allows dimer chains to form at low coverage and leads to a rough Ge film after the deposition of seven monolayers. Thus, Br prevents Ge hut formation. Removal of Br allows the system to reach a more energetically favored configuration of large Ge huts setting on wetting layers with a (2 · N) reconstruction. In this way, the onset of huts can be delayed until the desired thickness of Ge film is reached. By more precisely controlling the Br coverage, the interplay between kinetic and thermodynamic effects might be explored further. Acknowledgements This research was supported by the National Science Foundation. The experiments were performed in the Center for Microanalysis of Materials which is supported in part by the US Department of Energy, Division of Materials Science under Award No. DEFG02-91ER45439. We acknowledge the expert assistance of V. Petrova, S. Burdin, and E. Sammann. References [1] B. Voigtlander, Surf. Sci. Rep. 43 (2001) 127, and references cited therein. [2] F. Liu, F. Wu, M.G. Lagally, Chem. Rev. 97 (1997) 1045. [3] Y.-W. Mo, D.E. Savage, B.S. Swartzentruber, M.G. Lagally, Phys. Rev. Lett. 65 (1990) 1020. [4] M. Tomitori, K. Watanabe, M. Kobayashi, O. Nishikawa, Appl. Surf. Sci. 76/77 (1994) 322. [5] G. Medeiros-Ribeiro, A.M. Bratkovski, T.I. Kamins, D.A.A. Ohlberg, R.S. Williams, Science 279 (1998) 353. [6] F.M. Ross, R.M. Tromp, M.C. Reuter, Science 286 (1999) 1931. [7] M. Copel, M.C. Reuter, E. Kaxiras, R.M. Tromp, Phys. Rev. Lett. 63 (1989) 632. [8] M. Copel, M.C. Reuter, M. Horn von Hoegen, R.M. Tromp, Phys. Rev. B 42 (1990) 11682. [9] A. Kawano, I. Konomi, H. Azuma, T. Hioki, S. Noda, J. Appl. Phys. 74 (1993) 4265. [10] B. Voigtlander, A. Zinner, T. Weber, H.P. Bonzel, Phys. Rev. B 51 (1995) 7583. [11] M. Katayama, T. Nakayama, M. Aono, Phys. Rev. B 54 (1996) 8600. [12] M. Copel, R.M. Tromp, Appl. Phys. Lett. 58 (1991) 2648. [13] K. Oura, V.G. Lifshits, A.A. Saranin, A.V. Zotov, M. Katayama, Surf. Sci. Rep 35 (1999) 1. [14] A. Sakai, T. Tatsumi, Appl. Phys. Lett. 64 (1994) 52. [15] S.-J. Kahng, J.-Y. Park, K.H. Booth, J. Lee, Y. Kang, Y. Kuk, J. Vac. Sci. Technol. A 15 (1997) 927. [16] S.-J. Kahng, Y.H. Ha, J.-Y. Park, S. Kim, D.W. Moon, Y. Kuk, Phys. Rev. Lett. 80 (1998) 4931. [17] S.-J. Kahng, J.-Y. Park, Y. Kuk, Phys. Rev. B 60 (1999) 16558. [18] S. Jeong, A. Oshiyama, Phys. Rev. Lett. 79 (1997) 4425. [19] J. Nara, T. Sasaki, T. Ohno, Phys. Rev. Lett. 79 (1997) 4421. [20] K. Nakajima, A. Konishi, K. Kimura, Phys. Rev. Lett. 83 (1999) 1802. [21] X.R. Qin, B.S. Swartzentruber, M.G. Lagally, Phys. Rev. Lett. 84 (2000) 4645. [22] D.-S. Lin, J.-L. Wu, S.-Y. Pan, T.-C. Chiang, Phys. Rev. Lett. 90 (2003) 046102. [23] R.J. Wagner, E. Gulari, Phys. Rev. B 69 (2004) 195312. [24] D. Kandel, E. Kaxiras, Phys. Rev. Lett. 75 (1995) 2742. [25] D. Kandel, E. Kaxiras, Sol. State Phys. 54 (2000) 219. [26] S. Jeong, A. Oshiyama, Phys. Rev. B 58 (1998) 122958.
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[27] T. Hashizume, H. Kajiyama, Y. Suwu, S. Heike, S. Matsuura, J. Nara, T. Ohno, Appl. Surf. Sci. 216 (1998) 15. [28] B.R. Trenhaile, V.N. Antonov, G.J. Xu, Koji S. Nakayama, J.H. Weaver, Surf. Sci. Lett. 583 (2005) L135. [29] G.J. Xu, J.H. Weaver, Phys. Rev. B 70 (2004) 165321.
[30] S.M. Cohen, T.I. Hukka, Y.L. Yang, M.P. D’Evelyn, Thin Solid Films 225 (1993) 155. [31] G. Medeiros-Ribeiro, T.I. Kamins, D.A.A. Ohlberg, R.S. Williams, Phys. Rev. B 58 (1998) 3533.
Germanium growth on Br-terminated Si(1 0 0) B.R. Trenhaile, G.J. Xu 1, J.H. Weaver
*
Department of Physics, Department of Materials Science and Engineering, and Frederick Seitz Materials Research Laboratory, University of Illinois at Urbana-Champaign, Urbana, IL 61801, USA Received 15 December 2005; accepted for publication 16 May 2006 Available online 13 June 2006
Abstract The consequences of Ge deposition on Br-terminated Si(1 0 0) were studied with scanning tunneling microscopy at ambient temperature after annealing at 650 K. One monolayer of Br was sufficient to prevent the formation of Ge huts beyond the critical thickness of 3 ML. This is possible because Br acts as a surfactant whose presence lowered the diffusivity of Ge adatoms. Hindered mobility was manifest at low coverage through the formation of short Ge chains. Further deposition resulted in the extension and connection of the Ge chains and gave rise to the buildup of incomplete layers. The deposition of 7 ML of Ge resulted in a rough surface characterized by irregularly shaped clusters. A short 800 K anneal desorbed the Br and allowed Ge atoms to reorganize into the more energetically favorable ‘‘hut’’ structures produced by conventional Ge overlayer growth on Si(1 0 0). 2006 Elsevier B.V. All rights reserved. Keywords: Scanning tunneling microscopy; Germanium; Si(1 0 0); Epitaxy; Surfactant
1. Introduction The general picture of Ge growth on Si(1 0 0)-(2 · 1) at temperatures above 575 K and deposition rates less than 5 ML/min is that it proceeds via the Stranski–Krastanov mode with a wetting layer of 3 ML followed by threedimensional (3D) islanding [1]. Typically, the wetting layer undergoes a (2 · N) reconstruction where every Nth dimer is missing, and this periodic array of vacancy line defects relieves the misfit strain by allowing the Ge dimers near the trenches to relax [1,2]. The first 3D structures to form on the wetting layer are {5 0 1} faceted ‘‘hut’’ clusters, which have either rectangular or square bases [1,3,4]. Depending upon the growth temperature, the huts either remain to cover the surface or they give way to multifaceted domes or macroscopic clusters after 6–7 ML of Ge have been deposited [4–6]. *
Corresponding author. Tel.: +1 217 333 1440; fax: +1 217 333 2737. E-mail address: [email protected] (J.H. Weaver). 1 Present Address: GE Advanced Materials, Quartz, 24400 Highland Road, Richmond Heights, OH 44143, USA. 0039-6028/$ - see front matter 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.susc.2006.05.035
The growth of thin Ge films on Si(1 0 0) provides a model system for studying strained semiconductor heteroepitaxy as well as offering important device applications for the microelectronics industry. It is well known that growth modes can be manipulated by changing the surface energetics and lowering adatom diffusivity through the use of surfactants. Early studies used As, Sb, and Bi surfactants for Ge on Si(1 0 0), and they revealed that layer-bylayer growth continued well beyond the 3 ML observed on clean Si [7–11]. Surfactants have also been shown to generate more abrupt Ge–Si interfaces by suppressing atomic intermixing of Ge and Si [11–13]. Unfortunately, these surfactants also dope the growing layer. This side effect led to studies of Ge on hydrogen terminated surfaces where the transition to 3D growth can be delayed by a hydrogen flux applied during Ge deposition [14–16]. In this paper, we consider how an adsorbed monolayer of Br affects growth of Ge on Si(1 0 0). Using scanning tunneling microscopy (STM), we show that Br blocks the diffusion of the Ge adatoms during room-temperature deposition. A 650 K anneal of overlayers of nominal thicknesses through 3 ML allowed the random formation
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of short dimer chains. These chains expanded upon further deposition to produce a rugged Ge surface after the deposition of 7 ML. The removal of Br via GeBr2 desorption at 800 K allowed the formation of Ge hut clusters, and annealing at 850 K for 12 h resulted in the coarsening of the hut clusters. The results demonstrate that an adsorbed monolayer of Br acts as a surfactant during Ge growth on Si(1 0 0) by suppressing adatom diffusion and changing surface energetics, without introducing unwanted doping. 2. Experimental considerations The experiments were carried out in a vacuum system with operating pressure <4 · 1011 Torr. Imaging of clean and Ge- and Br-covered surfaces was done at ambient temperature using an Omicron variable temperature STM. The Si wafers were p-type, B-doped to 0.01–0.012 Xcm, and oriented within 0.5 of (1 0 0). Clean surfaces were prepared by degassing at 875 K for 12 h and then heating to 1475 K for 90 s at <5·1010 Torr. An optical pyrometer was used to monitor the sample temperature. For the clean surface, the defect area was 0.01–0.02 ML, primarily in the form of C-type defects and isolated dimer vacancies. A solidstate electrochemical cell derived from AgBr doped with 5 wt.% CdBr2 was used as a Br2 source. The clean Si surface was exposed to a flux of Br2 at room temperature to saturate the dangling bonds. Ge was deposited using an Omicron evaporator that contained a tungsten crucible heated by electron bombardment. All depositions occurred at room temperature with a rate of 9 · 104 ML/s, and the Ge overlayer was built up incrementally. After each deposition, the sample was annealed to 650 K for 30 min and then imaged at room temperature. At the beginning of each cycle, the surface was exposed to an additional Br2 fluence to replenish any Br lost during the short anneal. Cycles of Br passivation, Ge deposition, mild annealing, and room temperature imaging were repeated until a total of 7 ML of Ge had been deposited. A deposition of 0.12 ML, followed by one of 0.5 ML and one of 0.4 ML were performed to reach 1 ML. Thereafter, the remaining overlayer was built up by six depositions of 1 ML each. The sample was then heated to 800 K for 30 min to remove the Br and allow Ge redistribution to form 3D islands. Finally, the sample was annealed at 850 K for 12 h. 3. Results and Discussion STM images revealed that the deposition of 0.12 ML of Ge on Br saturated Si(1 0 0) at room temperature resulted in a poorly ordered surface consisting of mostly isolated adatoms. We speculate that arriving Ge atoms exchanged with Br at the site of arrival. This site exchange has been calculated to occur spontaneously, and it has been observed in studies of Ge and Si deposition on H–Si(1 0 0) [16–19]. The Ge atoms were immobile during subsequent room-temperature imaging because of Br site blocking.
Therefore, the presence of Br lowered the diffusivity of Ge adatoms with respect to the clean surface. Fig. 1(a) shows an STM image of the same surface after annealing to 650 K for 30 min. The dimer rows of the Br terminated Si (layer 0) run diagonally from the upper right to lower left. The post deposition anneal led to a more ordered surface as the Ge adatoms in layer 1 formed dimer chains that are perpendicular to the Si dimer rows. These chains show that the mild anneal has promoted adatom diffusion. For Ge deposition on clean Si(1 0 0), entropy considerations drive Ge atoms to interchange with Si atoms [20–23]. While adsorbed Br should increase the effective exchange barrier, it is possible that the dimer chains observed here include Si atoms. Since the mild anneal promoted Ge diffusion, it is important to discuss atomic scale mechanisms for adatom migration in the presence of surfactants. Kandel and Kaxiras used first-principles calculations to study surfactant mediated growth of semiconductor films, and they developed a model called diffusion-de-exchange-passivation (DDP) [24,25]. Rather than explaining the change in growth mode through reduced diffusion length for the adatoms, the DDP model considers the competition between the diffusion of adatoms on top of the surfactant layer, exchange/de-exchange of adatoms with the surfactant layer, and the attachment of adatoms to steps that are passivated with the surfactant. The DDP model has been successfully applied to semiconductor growth in the presence of groupIII and group-V surfactants. The adsorption and diffusion of Si adatoms on H-terminated Si(1 0 0) surfaces has been studied using first-principles calculations [18,26,19,27]. Those studies found that Si adatoms migrate on the hydrogenated surface through a H atom capture and release mechanism that is assisted by H atom mobility at elevated temperatures. This process gives a much higher barrier for Si adatom diffusion compared to clean Si(1 0 0). Such a capture and release diffusion process should also be applicable to Ge adatoms. The reduced diffusion length in the presence of H has been used to explain the experimental observation that H prevents 3D growth of Ge on Si(1 0 0) [14,16]. Bromine is chemically more similar to H than to the group-III and group-V atoms since it can only form one strong bond. Hence, the capture and release model of diffusion seems more relevant than the DDP model for Br-terminated Si(1 0 0). We also note that Br desorption from Si(1 0 0) is negligible at 650 K (0.005 ML of Br desorb for a 30 min anneal) and it is unlikely that desorption plays a significant role in the Ge mobility [28]. While the attachment of adatoms to steps in the presence of Br is certainly a critical part of the growth process, we take the restricted mobility of Ge adatoms to be the main effect of the Br layer. However, further investigations are needed to fully determine the microscopic mechanisms that occur during Br mediated growth. Fig. 1(b) shows that layer 1 was highly irregular after an additional 0.5 ML deposition and anneal. A type A step of
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Fig. 1. STM micrographs of Ge growth on Br–Si(1 0 0) (2.0 V sample bias). Ge was deposited at room temperature: (a) 0.12 ML of Ge imaged after 600 K anneal for 30 min, 15 · 15 nm2; (b) 0.62 ML of Ge imaged after an additional 0.5 ML Ge deposition followed by a 650 K anneal for 30 min, 25 · 25 nm2. A type-A step of the Si(1 0 0) surface is indicated by the dashed line; (c) 1 ML of Ge imaged after an additional 0.4 ML Ge was deposited followed by a 650 K anneal for 30 min, 30 · 30 nm2. A type-A and a type-B step of the Si(1 0 0) surface are indicated by the dashed lines and (d) 2 ML of Ge imaged after an additional 1 ML Ge deposition followed by an anneal to 650 K for 30 min, 30 · 30 nm2. The surface was passivated with Br before each Ge deposition.
the Si surface runs from the lower left to the upper right corner, as indicated by the dotted line. At this coverage, the dimer chains have connected to form features several dimer rows in width. Since features of this sort did not form at lower coverage, their development is most probably due to the connection of elongating dimer chains. The brightest features in the image reflect Ge adatoms that are beginning the formation of layer 2, and only about 0.03 ML of the Ge adatoms occupy layer 2. These chains appear to have formed at anti-phase boundaries in layer 1. Those boundaries form at the intersection of dimer chains that are out of phase by one atomic unit, and previously studies of Si growth on Br–Si(1 0 0) showed them to be chain nucleation sites [29]. Fig. 1(c) shows the morphology after a total of 1 ML Ge deposition (0.12 + 0.5 + 0.4 ML). Three layers are still
seen, but the dimer chains of layer 2 have increased in number and length to account for about 0.14 ML and layer 1 is composed of 0.87 ML. These chains make it more difficult to distinguish the original Si(1 0 0) steps, but the locations of two steps are indicated by the dotted lines. Fig. 1(d) shows the morphology after a total of 2 ML deposition. There are no substrate steps in this image, and four Ge layers can be seen. Layer 2 covers 0.85 ML, layer 3 accounts for 0.23 ML, and layer 4 is less than 0.01 ML Ge so that 0.09 ML of layer 1 remains incomplete. The limited diffusivity of the Ge adatoms on the Br terminated surface prevented the formation of a 2 · N reconstruction or large islands on the surfaces of Fig. 1 but analysis shows that there is interlayer diffusion. Assuming random adatom arrival, deposition onto the surface represented by Fig. 1(b) would result in a 59% probability of
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arriving on layer 1. Therefore, if diffusing atoms could not overcome the Ehrlich–Schwoebel barrier, the deposition of
an additional 0.4 ML of Ge would have resulted in a minimum of 0.27 ML of Ge in layer 2. Fig. 1(c) reveals that
Fig. 2. STM micrographs of 7 ML Ge grown on Br–Si(1 0 0) (2.0 V sample bias): (a) a rough film was produced after an anneal to 650 K for 30 min, but no Ge hut clusters were observed, 200 · 200 nm2; (b) higher resolution image showing the dimer rows on top of the irregular-shaped clusters, 35 · 35 nm2; (c) annealing to 800 K for 30 min desorbed the Br as GeBr2 and allowed square and rectangular {5 0 1} faceted huts to form, 200 · 200 nm2; (d) higher resolution image of a hut cluster on the 2 · N wetting layer, 35 · 35 nm2 and (e) shows the hut cluster with 3D rendering to emphasize the {5 0 1} facets. Line profiles across the top of such huts shows them to be fairly flat for 5 nm.
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only 0.14 ML occupancy of layer 2 and more than 0.13 ML moved from layer 2 to layer 1 during the anneal. Similar analysis of the surfaces represented by Fig. 1(b) and 1(c) show atom transfer to lower layers. We assume that the atoms staying on the layer of arrival encountered an antiphase boundary or another adatom to nucleate chains. The composition of the layers as they are built up results from a complex interplay between adatom diffusivity in the presence of Br and the number of atoms deposited between each annealing cycle. Fig. 2(a) shows the morphology after a total deposition of 7 ML of Ge. The substrate steps can no longer be seen and the surface is very irregular because of incomplete layer filling. Instead of huts bounded by {5 0 1} facets representative of Ge growth on clean Si(1 0 0), we observe clusters with shallower sides that do not appear to be bounded by a particular family of planes. Fig. 2(b) shows that the clusters are not amorphous because the top layer is composed of dimer-like rows that run in the [0 1 1] and ½011 directions, as on a (1 0 0) surface. Most of the dimer rows on top of the clusters are separated by a missing dimer row, and the lower layers on many clusters can be observed sticking out with dimer rows oriented perpendicular or parallel to the dimer rows of the top layer. The average angle of incline for the sides of the clusters is 8.0 with a range of 6.7–10.0, which is a slightly gentler slope then the 11 observed for {5 0 1} facets. The clusters range in height up to ˚ with an average of 7.3 A ˚. about 11 A We attribute these clusters to decreased adatom diffusivities and altered surface energetics associated with the high concentration of Br. The combination of limited mobility and decreased surface energies prevented the adatoms from reaching the conventional configuration of hut clusters on a 2 · N wetting layer. An increase in surface roughness has also been observed for Si growth on Br terminated Si(1 0 0) [29], and it was manifested here through the irregularly shaped clusters of varying heights. To determine if these clusters would transform into huts if the Br was removed, we annealed the sample at 800 K for 30 min. This should allow ample time for the removal of Br through GeBr2 desorption [30]. Fig. 2(c) reveals a morphology, that is, characteristic of Ge growth on clean Si(1 0 0) at high temperature, with huts on a Ge wetting layer. For comparison with Fig. 2(b), Fig. 2(d) shows a high resolution image of hut cluster on the 2 · N wetting layer. Unlike the clusters of Fig. 2(c), the huts are bounded by {5 0 1} facets as can be clearly seen in the 3D image of Fig. 2(e). As expected, extended annealing at 850 K allowed the larger huts to grow at the expense of smaller ones due to Ostwald ripening as has been observed in real time by several investigators [5,6,31]. 4. Conclusions We have studied the epitaxial growth of Ge on Br terminated Si(1 0 0). Br frustrates Ge adatom diffusion and isolated adatoms are observed after room temperature
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deposition. A mild anneal allows dimer chains to form at low coverage and leads to a rough Ge film after the deposition of seven monolayers. Thus, Br prevents Ge hut formation. Removal of Br allows the system to reach a more energetically favored configuration of large Ge huts setting on wetting layers with a (2 · N) reconstruction. In this way, the onset of huts can be delayed until the desired thickness of Ge film is reached. By more precisely controlling the Br coverage, the interplay between kinetic and thermodynamic effects might be explored further. Acknowledgements This research was supported by the National Science Foundation. The experiments were performed in the Center for Microanalysis of Materials which is supported in part by the US Department of Energy, Division of Materials Science under Award No. DEFG02-91ER45439. We acknowledge the expert assistance of V. Petrova, S. Burdin, and E. Sammann. References [1] B. Voigtlander, Surf. Sci. Rep. 43 (2001) 127, and references cited therein. [2] F. Liu, F. Wu, M.G. Lagally, Chem. Rev. 97 (1997) 1045. [3] Y.-W. Mo, D.E. Savage, B.S. Swartzentruber, M.G. Lagally, Phys. Rev. Lett. 65 (1990) 1020. [4] M. Tomitori, K. Watanabe, M. Kobayashi, O. Nishikawa, Appl. Surf. Sci. 76/77 (1994) 322. [5] G. Medeiros-Ribeiro, A.M. Bratkovski, T.I. Kamins, D.A.A. Ohlberg, R.S. Williams, Science 279 (1998) 353. [6] F.M. Ross, R.M. Tromp, M.C. Reuter, Science 286 (1999) 1931. [7] M. Copel, M.C. Reuter, E. Kaxiras, R.M. Tromp, Phys. Rev. Lett. 63 (1989) 632. [8] M. Copel, M.C. Reuter, M. Horn von Hoegen, R.M. Tromp, Phys. Rev. B 42 (1990) 11682. [9] A. Kawano, I. Konomi, H. Azuma, T. Hioki, S. Noda, J. Appl. Phys. 74 (1993) 4265. [10] B. Voigtlander, A. Zinner, T. Weber, H.P. Bonzel, Phys. Rev. B 51 (1995) 7583. [11] M. Katayama, T. Nakayama, M. Aono, Phys. Rev. B 54 (1996) 8600. [12] M. Copel, R.M. Tromp, Appl. Phys. Lett. 58 (1991) 2648. [13] K. Oura, V.G. Lifshits, A.A. Saranin, A.V. Zotov, M. Katayama, Surf. Sci. Rep 35 (1999) 1. [14] A. Sakai, T. Tatsumi, Appl. Phys. Lett. 64 (1994) 52. [15] S.-J. Kahng, J.-Y. Park, K.H. Booth, J. Lee, Y. Kang, Y. Kuk, J. Vac. Sci. Technol. A 15 (1997) 927. [16] S.-J. Kahng, Y.H. Ha, J.-Y. Park, S. Kim, D.W. Moon, Y. Kuk, Phys. Rev. Lett. 80 (1998) 4931. [17] S.-J. Kahng, J.-Y. Park, Y. Kuk, Phys. Rev. B 60 (1999) 16558. [18] S. Jeong, A. Oshiyama, Phys. Rev. Lett. 79 (1997) 4425. [19] J. Nara, T. Sasaki, T. Ohno, Phys. Rev. Lett. 79 (1997) 4421. [20] K. Nakajima, A. Konishi, K. Kimura, Phys. Rev. Lett. 83 (1999) 1802. [21] X.R. Qin, B.S. Swartzentruber, M.G. Lagally, Phys. Rev. Lett. 84 (2000) 4645. [22] D.-S. Lin, J.-L. Wu, S.-Y. Pan, T.-C. Chiang, Phys. Rev. Lett. 90 (2003) 046102. [23] R.J. Wagner, E. Gulari, Phys. Rev. B 69 (2004) 195312. [24] D. Kandel, E. Kaxiras, Phys. Rev. Lett. 75 (1995) 2742. [25] D. Kandel, E. Kaxiras, Sol. State Phys. 54 (2000) 219. [26] S. Jeong, A. Oshiyama, Phys. Rev. B 58 (1998) 122958.
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