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Heteroepitaxial growth of InAs on GaAs(0 0 1) by in situ STM located inside MBE growth chamber
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Microelectronics Journal 37 (2006) 1498–1504 www.elsevier.com/locate/mejo
Heteroepitaxial growth of InAs on GaAs(0 0 1) by in situ STM located inside MBE growth chamber S. Tsukamotoa,, G.R. Bellb, Y. Arakawaa a
Nanoelectronics Collaborative Research Center, The Institute of Industrial Science, The University of Tokyo, 4-6-1 Komaba, Tokyo 153-8505, Japan b Department of Physics, University of Warwick, Coventry CV4 7AL, UK Available online 5 July 2006
Abstract The growth of InAs on GaAs(0 0 1) is of great interest primarily due to the self-assembly of arrays of quantum dots (QDs) with excellent opto-electronic properties. However, a basic understanding of their spontaneous formation is lacking. Advanced experimental methods are required to probe these nanostructures dynamically in order to elucidate their growth mechanism. Scanning tunneling microscopy (STM) has been successfully applied to many GaAs-based materials grown by molecular beam epitaxy (MBE). Typical STM–MBE experiments involve quenching the sample and transferring it to a remote STM chamber under arsenic-free ultra-high vacuum. In the case of GaAs-based materials grown at substrate temperatures of 400–600 1C, operating the STM at room temperature ensures that the surface is essentially static on the time scale of STM imaging. To attempt dynamic experiments requires a system in which STM and MBE are incorporated into one unit in order to scan in situ during growth. Here, we discuss in situ STM results from just such a system, covering both QDs and the dynamics of the wetting layer. r 2006 Elsevier Ltd. All rights reserved. Keywords: Scanning tunneling microscopy; Molecular beam epitaxy; Quantum dots; InAs; GaAs
1. Introduction Scanning tunneling microscopy (STM) has been applied to GaAs and related materials since the 1980s. Somewhat more recently it has been applied to GaAs grown in vacuo by molecular beam epitaxy (MBE). The ability to image surface reconstructions and morphology as well as island structures and dopants has added substantially to the understanding of fundamental MBE processes. In a typical STM–MBE experiment, the ultra-high vacuum (UHV) chambers housing the MBE system and STM system are separate. After halting the MBE growth, the sample is moved through UHV to the STM for imaging. An early system of this type is described by Orr et al. [1] while a more recent design is detailed by Geng et al. [2]. These instruments enable one to interrupt growth and obtain ‘snapshots’ of the growing surface. The length scales of interest range from 10 10 m (surface reconstruction, dopant incorporation, alloying) to 10 7 m (islands, doCorresponding author. Tel./fax: +81 3 5452 6553.
E-mail address: [email protected] (S. Tsukamoto). 0026-2692/$ - see front matter r 2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.mejo.2006.05.011
mains, step-terrace structure). Of course, MBE is a dynamic process covering timescales from seconds or tens of seconds, which are typical monolayer (ML) completion times, to the vibrational periods of adatoms (10 13 s) which govern attempt frequencies for migration. Two conditions are placed on the system if one wishes to faithfully image the state of the surface during growth. Firstly, the surface should not change during imaging. Adsorption of contaminants and MBE materials is easily limited by the UHV conditions and the isolation of the MBE chamber from the STM system, respectively. However, the possibility of rearrangement of surface structures during imaging remains, so that the imaged surface may not represent the as-grown surface. MBE growth of III–V materials normally takes place at elevated surface temperatures (4500 1C for GaAs, 4400 1C for InAs). In most GaAs STM–MBE systems, STM imaging occurs at room temperature, and this is sufficient to suppress adatom migration. Furthermore, the stability of surface structures is easily checked by repeated imaging. The second condition relates to the process of quenching the sample from the MBE chamber. In the case of GaAs
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MBE, there is normally an excess flux of As2 or As4 molecules which stabilizes the surface at high temperatures. In some STM–MBE experiments, the substrate temperature is lowered at a fixed rate while the As flux is also lowered. The two rates must be balanced to avoid the surface becoming As-deficient or As-rich. Surface reconstructions have been imaged successfully at atomic resolution in this type of experiment: see, for example, Avery et al. [3] and LaBella et al. [4]. These surfaces are fairly uniform and can be stabilized without disrupting the surface reconstruction, which can be monitored throughout using reflection high energy electron diffraction (RHEED). Moreover, discrete structures such as islands can be affected by the continued annealing inherent in such a cooling protocol. A detailed cooling/growth scheme for the STM study of GaAs islands on GaAs(0 0 1) has been presented by Yang et al. [5]. An alternative to controlled cooling is to quench the surfaces to clean UHV as rapidly as possible. This technique has been applied to image discrete growth structures such as co-existing reconstruction domains [6], GaAs islands [7], and InAs quantum dots (QDs) [8]. In these studies, the small size of the samples required for STM combined with the relatively large thermal mass of the transfer mechanism produced an estimated cooling rate of around 50 1C s 1. The samples could be quickly rotated away from the impinging As flux and the total time to remove samples from the MBE chamber was around 5 s. An example of annealing discrete structures before such rapid-quenching is given in Fig. 1. Approximately, 0.1 ML GaAs was deposited on to GaAs(0 0 1)-(2 4) to produce an array of ML-height islands. Both the islands and substrate are (2 4) reconstructed. The sample was either rapid-quenched immediately (Fig. 1a) or allowed to anneal for 10 s at the growth temperature of 550 1C under the As2 flux prior to rapid quenching (Fig. 1b). In the latter case, the number density has dropped by nearly a factor of three with a corresponding increase in the mean island area. This is due to adatom migration during the short anneal and demonstrates that metrics such as the island size distribution can evolve markedly during the quench process. Note that all of the in vacuo images shown in this paper (Figs. 1, 3b, and 4b) were produced using the STM–MBE system at Imperial College London, by rapid-quench methods [6–8]. It is not possible to perform true dynamic imaging by using in vacuo remote STM–MBE. The development of a particular surface feature during growth cannot be followed since the tip cannot normally be returned to the same location after additional MBE growth. Of course, observations of ‘typical’ features and metrics such as island size distributions are still extremely useful. In order to overcome these limitations and any uncertainties about the effects of sample quenching, it is desirable to place the STM inside the MBE chamber (STMBE) and perform true in situ imaging during MBE growth [9]. This technique presents significant experimental challenges and shows the possibility of STM observations on GaAs surfaces even
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under As4 irradiation at 440 1C [10]. We discuss the use of an in situ STM–MBE (STMBE) system for III–V experiments, comparing them to conventional quenched in vacuo STM–MBE. We also show key results for InAs– GaAs(0 0 1) heteroepitaxial growth [11] which relate to the fundamental processes governing QD formation. 2. Experimental details The in situ STM–MBE (STMBE) system is based on an Omicron Micro-STM supported on a modified spring-anddamper vibration isolation system [9]. The tip points vertically upwards in this configuration for compatibility with the MBE chamber, in which the sample faces vertically downwards towards the cell ports. Effusion cells for In, Ga, As4 and other materials are included, along with a RHEED system. An efficient cryo-shroud covers both the upper chamber and a bay below the level of the cell ports for the STM when it is not in use. This prevents excess build-up of As on the STM mechanism, which is also protected by some additional shielding. To engage the STM with the sample, a motorized linear drive raises the whole STM stage until it docks with the MBE sample holder. The substrate holder and STM are designed in such a way that the sample is lifted from its mounting in the heater and becomes vibrationally isolated while maintained at high temperature by the radiative substrate heater. The STM is modified to allow molecular beams to impinge on the sample surface without the microscope shadowing the surface—in the present experiments, these were In and As4. The STM and MBE capabilities are therefore fully integrated. Running the STM within the MBE chamber produces several additional experimental challenges beyond those faced by conventional two-chamber STM–MBE systems. First is contamination of the microscope by the molecular beams: metallic deposits could easily short electrical wiring in the system, e.g. to the piezoelectric scanner tube. However, the modified STM was sufficiently well shielded to avoid any such problems even over imaging runs of many hours duration with the As4 flux engaged. Second is contamination of the STM tip itself by the beams. This did occur and will be discussed below. Vibrational isolation of the sample and microscope was good, although careful adjustment of the docking mechanism is necessary. Thermal drift was more of a problem, principally occurring due to changes in the substrate temperature or due to opening or closing of an effusion cell shutter. The radiant heat from the In cell in particular was sufficient to affect the STM when the shutter was closed to terminate growth or opened to initiate it. It was therefore difficult to image immediately after starting or stopping growth. Lastly, STM tip exchange was not possible with the past system unless the whole chamber was vented. However, it is possible with a new system in University of Tokyo, which aids experiments considerably given the increased probability of irreversible tip contamination when imaging under molecular beam fluxes.
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Fig. 1. Two 100 nm in vacuo STM images of 0.1 ML GaAs grown on GaAs(0 0 1)–(2 4) at 550 1C. The bright features are 1 ML high GaAs islands. In (a) the sample was rapid-quenched immediately after the termination of growth, while in (b) it was annealed for 10 s under the As flux before rapid-quenching.
Experiments were performed on 1–2 ML thick layers of InAs on GaAs(0 0 1) substrates 11 miscut towards (1 1 1)A grown by standard MBE. InAs was deposited at rates o0.01 ML s 1 and at a substrate temperature of 400 1C. Low growth rates are useful for producing long wavelength emission from InAs QDs [12] but in the present context also assist thermal stabilization of the STM after growth is initiated. STM images were obtained in constant current mode using tunneling currents of around 0.2–0.8 nA and a sample bias voltage of 3 V. Etched tungsten tips were employed. The substrate temperature was 400 1C and the typical As4 flux was (1.570.5) 1014 molecules cm 2 s 1 estimated from the beam equivalent pressure. The growth was monitored using RHEED and the normal sequence of RHEED patterns was observed [13]. In particular, a symmetric (1 3) pattern was observed for the InAs wetting layer (WL) and a ‘chevron’ pattern for the QDs. The images presented in this paper were obtained with the sample held at 400 1C under As4 flux without co-deposition of In (Figs. 2, 3a, and 4a) and at 300 1C without As4 and In fluxes (Figs. 5 and 6). 3. Results Fig. 2 displays two 300 nm STM images obtained sequentially after the growth of 2 ML InAs, acquired over a time of 100 s per image. The overall drift was estimated at 0.04 nm s 1 and the tip is rastered left-to-right moving down the images as shown. The first (upper) part of image (a) is affected by strong ‘streaking’ in the direction of tip motion. This effect suddenly disappears, resulting in clear resolution of several surface structures. The stepterrace structure of the substrate is revealed, with mostly single ML-height steps. Islands with heights of 1 ML and quite isotropic shape are also observed on the terraces. This morphology is typical of the InAs WL observed in rapidquench in vacuo STM studies [13]. The bright features are InAs QDs whose average dimensions are 1.7 nm height and 5 nm width (although with a broad size fluctuation). The
density of QDs is 6 1011 cm 2 which is in line with the value expected for the growth conditions [12]. The QDs are preferentially sited at the bases of ML steps, a feature also often observed in quenched STM studies. A careful comparison of images (a) and (b) (including direct image subtraction) shows that there is essentially no change in the shape or position of the QDs between the two images, nor in the shape of the step edges or ML-height islands. Fig. 3 shows two 200 nm STM images of 1 ML InAs deposited on GaAs(0 0 1). Image (a) was obtained using in situ STM at 400 1C under the As4 flux, while image (b) was obtained using rapid-quench in vacuo STM. The drift during imaging is 0.16 nm s 1 for image (a), while the bright irregular feature is a tip deposit rather than a QD. The horizontal streaks are probably due to tip instability caused by As molecules—they are absent in the in vacuo image (b). As for the larger images in Fig. 2, the stepterrace structure of the surface is easily resolved in (a). Rather than ML-height islands, however, the surface is covered by ML-depth pits. This ‘pit’ morphology is typical of InAs WLs at overages close to 1 ML [14]. For the sample (a) grown and held at 400 1C the density is around 1 1011 cm 2 and the pits have an average lateral size of 30 nm. For the quenched sample (b) grown at 480 1C, the pits are rather larger and less numerous (about 5 1010 cm 2 and 4 nm 10 nm typical dimensions). This reflects the greater adatom mobility at higher growth temperatures. Sequential imagings of fixed areas of the WL show no evolution in the shapes of islands or step edges. As in Fig. 2, the WL appears to be stable over the imaging timescale of hundreds of seconds. Two higher magnification images of the WL are shown in Fig. 4. Again, image (a) was obtained by in situ STM, holding the sample at 400 1C under the As4 flux after depositing 1.5 ML InAs, while image (b) was obtained by in vacuo STM after quenching a 1.2 ML InAs layer from a growth temperature of 350 1C. Both images are 25 nm 25 nm. Image (a) is not atomically resolved, although clear anisotropic structure is visible with short,
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Fig. 2. Two 300 nm sequential in situ STM images of InAs QDs on vicinal GaAs(0 0 1) [11]. The upper half of image (a) is partly obscured by streaking (see text). The white dots are InAs QDs with average height 1.7 nm and width 5 nm. The step array and ML-height islands in the WL are also visible.
Fig. 3. Two 200 nm STM images of the InAs WL on GaAs(0 0 1). Image (a): in situ STM at 400 1C under As4 flux, InAs coverage 1 ML [11]. Image (b): quenched in vacuo STM, 0.9 ML InAs deposited at 480 1C on a singular substrate.
bright segments elongated along the [1¯ 1 0] direction. Atomic resolution is achieved in image (b) although considerable structural disorder is evident. The phase diagram for surface reconstructions of the WL has been mapped by Belk et al. [13] and for both the samples imaged in Fig. 3 the symmetric (1 3) RHEED pattern is observed (the fractional order rods are spaced equally between the integer orders). In image (b) one observes dark rows elongated along the [1¯ 1 0] azimuth with brighter rows in between, on top of which are superimposed individual bright features. The spacing between these rows along the [1 1 0] azimuth is typically 1.2 nm, indicating threefold periodicity, although there is considerable ‘kinking’ in the rows. The in situ STM images are less clear, but [1 1 0] periodicities are shown by the cross-sections below the images (their location on image (a) is indicated by white
lines). As well as threefold spacings of 1.2 nm (crosssection 2), both two-fold and four-fold spacings can be seen (cross-section 1). The individual bright features evident on the quenched STM image (b) are absent in the in situ image (a). The effects of imaging conditions on the atomic-scale structure of the STM results are further highlighted by Fig. 5, which shows an in situ image obtained at 300 1C but without any incident As flux. The images in Figs. 4 and 5 are on the same scale (25 nm), and although the coverage in Fig. 5 is slightly higher at 1.6 ML, the expected reconstructions are comparable [13]. At elevated temperature but in the absence of an As flux, the 3-fold periodicity is clearer than in the case of Fig. 4(a) but without the additional discrete features present in Fig. 4(b) after quenching. Some streaking remains in the image.
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Fig. 4. Two 25 nm filled-states STM images of the InAs WL on GaAs(0 0 1). In (a) the InAs coverage is 1.5 ML and the growth temperature is 400 1C; the image was obtained using in situ STM [11]. In (b), the InAs coverage is 1.2 ML and the growth temperature was 350 1C; the image was obtained by quenched in vacuo STM. The cross-sections below are taken from image (a) as indicated by the white lines and show 2 , 3 and 4 spacings [11].
4. Discussion
Fig. 5. A 25 nm 25 nm STM image of 1.6 ML InAs on GaAs(0 0 1), obtained in situ at 300 1C but without an incident As flux. Clear 3-fold periodicity is visible.
The streaking in the upper half of image (a) in Fig. 2 illustrates one of the additional difficulties with imaging under the As4 flux. During typical MBE growth with excess As, a surface population of mobile As molecules is present [7]. Contamination of the tip, due to As molecules becoming unstably adsorbed near the tip apex, is the most likely cause of the streaking. These effects were rarely observed without the incident As4 flux. It was found that higher scan speeds (tip motion rates X2000 nm s 1 compared to more typical values of hundreds of nm s 1) helped to prevent the onset of streaking and instability [10]. However, such high scan speeds make it more difficult to obtain atomic resolution. After lengthy scanning under As4, the tip would normally become very contaminated and no resolution of the substrate could be obtained. This could be alleviated by scanning at relatively high tunneling current and bias voltage well away from the area of interest. The resulting combination of tip heating and field desorption was sufficient to remove the contamination and restore resolution. Subsequent large-scale scans of the areas used for such procedures showed the presence of significant deposits, sometimes up to tens of nm in height, confirming
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that material had been desorbed from the tip. A final point regards shadowing of As4 beam by the tip itself. Since the final radius of curvature of the STM tips may be some tens of nm, one would expect the tip itself to block the incoming fluxes over some lateral distance relating to this radius and the polar incidence angles of the incoming beams. In the case of InAs–GaAs(0 0 1), the incoming As flux stabilizes the (1 3) reconstruction. In As-deficient conditions, this reverts to a (4 2)/c(8 2) In-stable phase which is very easy to observe with STM due to its characteristic, highly anisotropic domain structure [14]. No such As-deficient phases were observed in the present experiments, indicating that the migration lengths of surface As molecules are sufficient to maintain As-rich reconstructions even in areas where the incident beam might be blocked by the tip. Despite the additional tip cleaning required while imaging under As4, the in situ STM images shown in Figs. 2 and 3 are comparable to in vacuo results, i.e. the nanometer-scale resolution is essentially unimpaired. This confirms the ability of in situ STM to follow the development of epitaxial structures on this scale. It is useful to assess the possible effects of annealing under the As4 flux on the morphology of the QD-covered surface (note that this in situ STM–MBE (STMBE) experiment is distinct from in vacuo high-temperature STM because of the group V overpressure). Subject to kinetic hindrance through slow adatom exchange between QDs, one expects to observe ripening of the QD array simply due to the energetic balance between surface area and strain relaxation. In practical terms, surfaces are often prepared ex situ by As-decapping [15] which may involve extended annealing at 300 1C. In the case of InAs–GaAs QDs, ex situ AFM is also widely used on large MBE platens which may cool rather slowly. The present results indicate that the WL and QD morphology is quite stable over timescales of at least hundreds of seconds at 400 1C. Quenched mode microscopy involving a similar or smaller post-growth thermal load should therefore faithfully reflect the as-grown morphology. We now turn to high resolution imaging of the WL (Figs. 4 and 5). The absence of As flux allows stable in situ imaging at much lower scan speeds, which in turn tends to produce improved resolution (the scan speed for the images in Figs. 5 and 6 is approximately 70 nm s 1). The overall structure of the WL surface observed by the in situ STM is in agreement with previous RHEED and STM work [13,16]. In particular, the 3-fold periodicity is dominant but there is considerable structural disorder. Atomic resolution is regularly achieved in quenched-mode STM, but images of the InAs WL normally show the surface covered in small discrete features which have been assigned to In adatoms [13,14]. This is the case even for ‘static’ surfaces annealed under As flux for a long time before quenching [13], implying that there is a substantial equilibrium population of In adatoms which are frozen in by the quench. Mobile surface species are expected to cause some streaking in the in situ images, and indeed this is normally observed.
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Despite the additional difficulty imposed the high scan speeds, high resolution images under As flux typically show more severe streaking (Fig. 4(a)) than in vacuum-anneal conditions (Fig. 5). This is probably due to a combination of mobile As molecular surface species [7] and In adatoms. In the case of the annealed surface without As flux, streaking is still present but is less pronounced, indicating the presence of mobile In adatoms only. Experimental measurements of total surface stress [17] indicate that a significant fraction of deposited In in the WL does not exist in a surface phase supporting the epitaxial stress. Recent STMBE results also indicate that not all deposited In is incorporated into the lattice [18]. Furthermore, theoretical work by Ishii et al. [19] and Kratzer et al. [20] indicates that In adatoms should be highly mobile on InAs/GaAs(0 0 1) thin film surfaces. In the latter study, the alloyed In0.66Ga0.34As (2 3) structure was shown to be stable, while the activation energies derived for In migration on this surface were derived as just 0.13 eV in the [1¯ 1 0] direction and 0.29 eV along the [1 1 0] direction. The present observations (Fig. 5) are consistent with the existence of a population of In adatoms which are mobile even at 300 1C and migrate on a rather disordered (1 3) reconstructed WL surface. Although the morphology of steps, pits and QDs appears quite stable over several hundred seconds when annealed at 400 1C under As flux (Fig. 2), it is possible to observe a small degree of step motion when the WL is annealed at 300 1C without As flux. In Fig. 6, several step edges are visible. The upper step edges lie mainly parallel to the rows of the reconstruction, while the lower left step edges cut across the reconstruction rows. Careful comparison of the images shows that some step erosion has taken place going from image (a) to image (b). This is highlighted by the white arrows. However, the step edges parallel to the reconstruction rows near the top of the images are hardly affected. This implies that As desorption may take place preferentially at steps running along the [1 1 0] azimuth (first-principles calculations would be useful in elucidating this point). Further comparison of the images shows that the detailed pattern of the segments and trenches of the (1 3) reconstruction is most strongly altered in the vicinity of the step erosion between images (a) and (b). This is highlighted by a white circle on the images. This is consistent with the release of additional In adatoms by the As desorption process which migrate and locally alter the reconstruction domains. It is probable that an incident As flux suppresses this process and leads to a more stable step morphology (Fig. 2). 5. Conclusions A comparison between in situ and in vacuo STM–MBE studies has been made, concentrating on the InAs– GaAs(0 0 1) system. Despite the necessity of extra STM tip treatment when scanning under the As4 flux, it has been shown that nanometer-scale imaging is comparable to the
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Fig. 6. Two 50 nm 50 nm STM image of 1.6 ML InAs on GaAs(0 0 1), obtained in situ under continuous annealing at 300 1C without an As flux. The location of each image is the same (subject to some drift) and the image acquisition time was 400 s for each image. Image (b) was obtained directly after image (a). The 3-fold periodicity of the surface reconstruction is visible as approximately horizontal striped segments and several step edges are also present. The white arrows and circle highlight changes between the two images.
conventional in vacuo methods. Furthermore, comparisons of atomic scale imaging with or without As flux strongly support the presence of a substantial population of mobile In adatoms at temperatures as low as 300 1C, and the presence of additional mobile As-related species under As flux. The disordered (1 3) reconstruction of the WL can be clearly observed using in situ imaging. Both QDs and the terrace/island/pit structures of the WL appear to be stable when annealed under As4 flux at 400 1C over timescales of several hundred seconds. In the absence of As flux, some step erosion has been observed at 300 1C on the InAs WL. The preferential site for As re-evaporation appears to be at step edges running parallel to the [1 1 0] direction and this leads to the release of In adatoms which locally modify the (1 3) reconstruction. Acknowledgments The authors thank Dr. Nobuyuki Koguchi at the National Institute for Materials Science in Japan (NIMS) for sincere encouragement and support. The STMBE work was performed at NIMS (Figs. 2, 3a, and 4a) and continues with a new modified STMBE system at University of Tokyo (Figs. 5 and 6) assisted by IT program of the Ministry of Education, Culture, Sports, Science and Technology of the Japanese Government. References [1] B.G. Orr, C.W. Snyder, M. Johnson, Rev. Sci. Instrum. 62 (1991) 1400–1403. [2] P. Geng, J. Marquez, L. Geelhar, J. Platen, C. Stezer, K. Jacobi, Rev. Sci. Instrum. 71 (2000) 504–508.
[3] A.R. Avery, D.M. Holmes, J. Sudijono, T.S. Jones, B.A. Joyce, Surf. Sci. 323 (1995) 91–101. [4] V.P. LaBella, H. Yang, D.W. Bullock, P.M. Thibado, P. Kratzer, M. Scheffler, Phys. Rev. Lett. 83 (1999) 2989–2992. [5] H. Yang, V.P. LaBella, D.W. Bullock, P.M. Thibado, J. Vac. Sci. Technol. B 17 (1999) 1778–1780. [6] G.R. Bell, J.G. Belk, C.F. McConville, T.S. Jones, Phys. Rev. B 59 (1999) 2947–2955. [7] M. Itoh, G.R. Bell, A.R. Avery, T.S. Jones, B.A. Joyce, D.D. Vvedensky, Phys. Rev. Lett. 81 (1998) 633–636. [8] P.B. Joyce, T.J. Krzyzewski, G.R. Bell, T.S. Jones, B.A. Joyce, Phys. Rev. B 58 (1998) R15981–R15984. [9] S. Tsukamoto, N. Koguchi, J. Crystal Growth 201/202 (1999) 118–123; S. Tsukamoto, N. Koguchi, J. Crystal Growth 209 (2000) 258–262. [10] S. Tsukamoto, N. Koguchi, In: MRS Symposium Proceedings 648 (2001) P11.20.1–P.11.20.6. [11] G.R. Bell, M. Pristovsek, S. Tsukamoto, B.G. Orr, Y. Arakawa, N. Koguchi, Surf. Sci. 544 (2003) 234–240. [12] P.B. Joyce, T.J. Krzyzewski, G.R. Bell, T.S. Jones, S. Malik, D. Childs, R. Murray, Phys. Rev. B 62 (2000) 10891–10895. [13] J.G. Belk, C.F. McConville, J. Sudijono, T.S. Jones, B.A. Joyce, Surf. Sci. 387 (1997) 213–226. [14] T.J. Krzyzewski, P.B. Joyce, G.R. Bell, T.S. Jones, Surf. Sci. 532–535 (2003) 822–827. [15] F. Patella, S. Nufris, F. Arciprete, M. Fanfoni, E. Placidi, A. Sgarlata, A. Balzarotti, Phys. Rev. B 67 (2003) 205308. [16] T.J. Krzyzewski, P.B. Joyce, G.R. Bell, T.S. Jones, Surf. Sci. 517 (2002) 8–16. [17] J.M. Garcı´ a, J.P. Silveira, F. Briones, Appl. Phys. Lett. 77 (2000) 409–411. [18] S. Tsukamoto, T. Honma, G.R. Bell, A. Ishii, Y. Arakawa, Small 2 (2006) 386–389. [19] A. Ishii, K. Fujiwara, T. Aisaka, Appl. Surf. Sci. 216 (2003) 478–482. [20] P. Kratzer, E. Penev, M. Scheffler, Appl. Surf. Sci. 216 (2003) 436–446.
Microelectronics Journal 37 (2006) 1498–1504 www.elsevier.com/locate/mejo
Heteroepitaxial growth of InAs on GaAs(0 0 1) by in situ STM located inside MBE growth chamber S. Tsukamotoa,, G.R. Bellb, Y. Arakawaa a
Nanoelectronics Collaborative Research Center, The Institute of Industrial Science, The University of Tokyo, 4-6-1 Komaba, Tokyo 153-8505, Japan b Department of Physics, University of Warwick, Coventry CV4 7AL, UK Available online 5 July 2006
Abstract The growth of InAs on GaAs(0 0 1) is of great interest primarily due to the self-assembly of arrays of quantum dots (QDs) with excellent opto-electronic properties. However, a basic understanding of their spontaneous formation is lacking. Advanced experimental methods are required to probe these nanostructures dynamically in order to elucidate their growth mechanism. Scanning tunneling microscopy (STM) has been successfully applied to many GaAs-based materials grown by molecular beam epitaxy (MBE). Typical STM–MBE experiments involve quenching the sample and transferring it to a remote STM chamber under arsenic-free ultra-high vacuum. In the case of GaAs-based materials grown at substrate temperatures of 400–600 1C, operating the STM at room temperature ensures that the surface is essentially static on the time scale of STM imaging. To attempt dynamic experiments requires a system in which STM and MBE are incorporated into one unit in order to scan in situ during growth. Here, we discuss in situ STM results from just such a system, covering both QDs and the dynamics of the wetting layer. r 2006 Elsevier Ltd. All rights reserved. Keywords: Scanning tunneling microscopy; Molecular beam epitaxy; Quantum dots; InAs; GaAs
1. Introduction Scanning tunneling microscopy (STM) has been applied to GaAs and related materials since the 1980s. Somewhat more recently it has been applied to GaAs grown in vacuo by molecular beam epitaxy (MBE). The ability to image surface reconstructions and morphology as well as island structures and dopants has added substantially to the understanding of fundamental MBE processes. In a typical STM–MBE experiment, the ultra-high vacuum (UHV) chambers housing the MBE system and STM system are separate. After halting the MBE growth, the sample is moved through UHV to the STM for imaging. An early system of this type is described by Orr et al. [1] while a more recent design is detailed by Geng et al. [2]. These instruments enable one to interrupt growth and obtain ‘snapshots’ of the growing surface. The length scales of interest range from 10 10 m (surface reconstruction, dopant incorporation, alloying) to 10 7 m (islands, doCorresponding author. Tel./fax: +81 3 5452 6553.
E-mail address: [email protected] (S. Tsukamoto). 0026-2692/$ - see front matter r 2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.mejo.2006.05.011
mains, step-terrace structure). Of course, MBE is a dynamic process covering timescales from seconds or tens of seconds, which are typical monolayer (ML) completion times, to the vibrational periods of adatoms (10 13 s) which govern attempt frequencies for migration. Two conditions are placed on the system if one wishes to faithfully image the state of the surface during growth. Firstly, the surface should not change during imaging. Adsorption of contaminants and MBE materials is easily limited by the UHV conditions and the isolation of the MBE chamber from the STM system, respectively. However, the possibility of rearrangement of surface structures during imaging remains, so that the imaged surface may not represent the as-grown surface. MBE growth of III–V materials normally takes place at elevated surface temperatures (4500 1C for GaAs, 4400 1C for InAs). In most GaAs STM–MBE systems, STM imaging occurs at room temperature, and this is sufficient to suppress adatom migration. Furthermore, the stability of surface structures is easily checked by repeated imaging. The second condition relates to the process of quenching the sample from the MBE chamber. In the case of GaAs
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MBE, there is normally an excess flux of As2 or As4 molecules which stabilizes the surface at high temperatures. In some STM–MBE experiments, the substrate temperature is lowered at a fixed rate while the As flux is also lowered. The two rates must be balanced to avoid the surface becoming As-deficient or As-rich. Surface reconstructions have been imaged successfully at atomic resolution in this type of experiment: see, for example, Avery et al. [3] and LaBella et al. [4]. These surfaces are fairly uniform and can be stabilized without disrupting the surface reconstruction, which can be monitored throughout using reflection high energy electron diffraction (RHEED). Moreover, discrete structures such as islands can be affected by the continued annealing inherent in such a cooling protocol. A detailed cooling/growth scheme for the STM study of GaAs islands on GaAs(0 0 1) has been presented by Yang et al. [5]. An alternative to controlled cooling is to quench the surfaces to clean UHV as rapidly as possible. This technique has been applied to image discrete growth structures such as co-existing reconstruction domains [6], GaAs islands [7], and InAs quantum dots (QDs) [8]. In these studies, the small size of the samples required for STM combined with the relatively large thermal mass of the transfer mechanism produced an estimated cooling rate of around 50 1C s 1. The samples could be quickly rotated away from the impinging As flux and the total time to remove samples from the MBE chamber was around 5 s. An example of annealing discrete structures before such rapid-quenching is given in Fig. 1. Approximately, 0.1 ML GaAs was deposited on to GaAs(0 0 1)-(2 4) to produce an array of ML-height islands. Both the islands and substrate are (2 4) reconstructed. The sample was either rapid-quenched immediately (Fig. 1a) or allowed to anneal for 10 s at the growth temperature of 550 1C under the As2 flux prior to rapid quenching (Fig. 1b). In the latter case, the number density has dropped by nearly a factor of three with a corresponding increase in the mean island area. This is due to adatom migration during the short anneal and demonstrates that metrics such as the island size distribution can evolve markedly during the quench process. Note that all of the in vacuo images shown in this paper (Figs. 1, 3b, and 4b) were produced using the STM–MBE system at Imperial College London, by rapid-quench methods [6–8]. It is not possible to perform true dynamic imaging by using in vacuo remote STM–MBE. The development of a particular surface feature during growth cannot be followed since the tip cannot normally be returned to the same location after additional MBE growth. Of course, observations of ‘typical’ features and metrics such as island size distributions are still extremely useful. In order to overcome these limitations and any uncertainties about the effects of sample quenching, it is desirable to place the STM inside the MBE chamber (STMBE) and perform true in situ imaging during MBE growth [9]. This technique presents significant experimental challenges and shows the possibility of STM observations on GaAs surfaces even
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under As4 irradiation at 440 1C [10]. We discuss the use of an in situ STM–MBE (STMBE) system for III–V experiments, comparing them to conventional quenched in vacuo STM–MBE. We also show key results for InAs– GaAs(0 0 1) heteroepitaxial growth [11] which relate to the fundamental processes governing QD formation. 2. Experimental details The in situ STM–MBE (STMBE) system is based on an Omicron Micro-STM supported on a modified spring-anddamper vibration isolation system [9]. The tip points vertically upwards in this configuration for compatibility with the MBE chamber, in which the sample faces vertically downwards towards the cell ports. Effusion cells for In, Ga, As4 and other materials are included, along with a RHEED system. An efficient cryo-shroud covers both the upper chamber and a bay below the level of the cell ports for the STM when it is not in use. This prevents excess build-up of As on the STM mechanism, which is also protected by some additional shielding. To engage the STM with the sample, a motorized linear drive raises the whole STM stage until it docks with the MBE sample holder. The substrate holder and STM are designed in such a way that the sample is lifted from its mounting in the heater and becomes vibrationally isolated while maintained at high temperature by the radiative substrate heater. The STM is modified to allow molecular beams to impinge on the sample surface without the microscope shadowing the surface—in the present experiments, these were In and As4. The STM and MBE capabilities are therefore fully integrated. Running the STM within the MBE chamber produces several additional experimental challenges beyond those faced by conventional two-chamber STM–MBE systems. First is contamination of the microscope by the molecular beams: metallic deposits could easily short electrical wiring in the system, e.g. to the piezoelectric scanner tube. However, the modified STM was sufficiently well shielded to avoid any such problems even over imaging runs of many hours duration with the As4 flux engaged. Second is contamination of the STM tip itself by the beams. This did occur and will be discussed below. Vibrational isolation of the sample and microscope was good, although careful adjustment of the docking mechanism is necessary. Thermal drift was more of a problem, principally occurring due to changes in the substrate temperature or due to opening or closing of an effusion cell shutter. The radiant heat from the In cell in particular was sufficient to affect the STM when the shutter was closed to terminate growth or opened to initiate it. It was therefore difficult to image immediately after starting or stopping growth. Lastly, STM tip exchange was not possible with the past system unless the whole chamber was vented. However, it is possible with a new system in University of Tokyo, which aids experiments considerably given the increased probability of irreversible tip contamination when imaging under molecular beam fluxes.
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Fig. 1. Two 100 nm in vacuo STM images of 0.1 ML GaAs grown on GaAs(0 0 1)–(2 4) at 550 1C. The bright features are 1 ML high GaAs islands. In (a) the sample was rapid-quenched immediately after the termination of growth, while in (b) it was annealed for 10 s under the As flux before rapid-quenching.
Experiments were performed on 1–2 ML thick layers of InAs on GaAs(0 0 1) substrates 11 miscut towards (1 1 1)A grown by standard MBE. InAs was deposited at rates o0.01 ML s 1 and at a substrate temperature of 400 1C. Low growth rates are useful for producing long wavelength emission from InAs QDs [12] but in the present context also assist thermal stabilization of the STM after growth is initiated. STM images were obtained in constant current mode using tunneling currents of around 0.2–0.8 nA and a sample bias voltage of 3 V. Etched tungsten tips were employed. The substrate temperature was 400 1C and the typical As4 flux was (1.570.5) 1014 molecules cm 2 s 1 estimated from the beam equivalent pressure. The growth was monitored using RHEED and the normal sequence of RHEED patterns was observed [13]. In particular, a symmetric (1 3) pattern was observed for the InAs wetting layer (WL) and a ‘chevron’ pattern for the QDs. The images presented in this paper were obtained with the sample held at 400 1C under As4 flux without co-deposition of In (Figs. 2, 3a, and 4a) and at 300 1C without As4 and In fluxes (Figs. 5 and 6). 3. Results Fig. 2 displays two 300 nm STM images obtained sequentially after the growth of 2 ML InAs, acquired over a time of 100 s per image. The overall drift was estimated at 0.04 nm s 1 and the tip is rastered left-to-right moving down the images as shown. The first (upper) part of image (a) is affected by strong ‘streaking’ in the direction of tip motion. This effect suddenly disappears, resulting in clear resolution of several surface structures. The stepterrace structure of the substrate is revealed, with mostly single ML-height steps. Islands with heights of 1 ML and quite isotropic shape are also observed on the terraces. This morphology is typical of the InAs WL observed in rapidquench in vacuo STM studies [13]. The bright features are InAs QDs whose average dimensions are 1.7 nm height and 5 nm width (although with a broad size fluctuation). The
density of QDs is 6 1011 cm 2 which is in line with the value expected for the growth conditions [12]. The QDs are preferentially sited at the bases of ML steps, a feature also often observed in quenched STM studies. A careful comparison of images (a) and (b) (including direct image subtraction) shows that there is essentially no change in the shape or position of the QDs between the two images, nor in the shape of the step edges or ML-height islands. Fig. 3 shows two 200 nm STM images of 1 ML InAs deposited on GaAs(0 0 1). Image (a) was obtained using in situ STM at 400 1C under the As4 flux, while image (b) was obtained using rapid-quench in vacuo STM. The drift during imaging is 0.16 nm s 1 for image (a), while the bright irregular feature is a tip deposit rather than a QD. The horizontal streaks are probably due to tip instability caused by As molecules—they are absent in the in vacuo image (b). As for the larger images in Fig. 2, the stepterrace structure of the surface is easily resolved in (a). Rather than ML-height islands, however, the surface is covered by ML-depth pits. This ‘pit’ morphology is typical of InAs WLs at overages close to 1 ML [14]. For the sample (a) grown and held at 400 1C the density is around 1 1011 cm 2 and the pits have an average lateral size of 30 nm. For the quenched sample (b) grown at 480 1C, the pits are rather larger and less numerous (about 5 1010 cm 2 and 4 nm 10 nm typical dimensions). This reflects the greater adatom mobility at higher growth temperatures. Sequential imagings of fixed areas of the WL show no evolution in the shapes of islands or step edges. As in Fig. 2, the WL appears to be stable over the imaging timescale of hundreds of seconds. Two higher magnification images of the WL are shown in Fig. 4. Again, image (a) was obtained by in situ STM, holding the sample at 400 1C under the As4 flux after depositing 1.5 ML InAs, while image (b) was obtained by in vacuo STM after quenching a 1.2 ML InAs layer from a growth temperature of 350 1C. Both images are 25 nm 25 nm. Image (a) is not atomically resolved, although clear anisotropic structure is visible with short,
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Fig. 2. Two 300 nm sequential in situ STM images of InAs QDs on vicinal GaAs(0 0 1) [11]. The upper half of image (a) is partly obscured by streaking (see text). The white dots are InAs QDs with average height 1.7 nm and width 5 nm. The step array and ML-height islands in the WL are also visible.
Fig. 3. Two 200 nm STM images of the InAs WL on GaAs(0 0 1). Image (a): in situ STM at 400 1C under As4 flux, InAs coverage 1 ML [11]. Image (b): quenched in vacuo STM, 0.9 ML InAs deposited at 480 1C on a singular substrate.
bright segments elongated along the [1¯ 1 0] direction. Atomic resolution is achieved in image (b) although considerable structural disorder is evident. The phase diagram for surface reconstructions of the WL has been mapped by Belk et al. [13] and for both the samples imaged in Fig. 3 the symmetric (1 3) RHEED pattern is observed (the fractional order rods are spaced equally between the integer orders). In image (b) one observes dark rows elongated along the [1¯ 1 0] azimuth with brighter rows in between, on top of which are superimposed individual bright features. The spacing between these rows along the [1 1 0] azimuth is typically 1.2 nm, indicating threefold periodicity, although there is considerable ‘kinking’ in the rows. The in situ STM images are less clear, but [1 1 0] periodicities are shown by the cross-sections below the images (their location on image (a) is indicated by white
lines). As well as threefold spacings of 1.2 nm (crosssection 2), both two-fold and four-fold spacings can be seen (cross-section 1). The individual bright features evident on the quenched STM image (b) are absent in the in situ image (a). The effects of imaging conditions on the atomic-scale structure of the STM results are further highlighted by Fig. 5, which shows an in situ image obtained at 300 1C but without any incident As flux. The images in Figs. 4 and 5 are on the same scale (25 nm), and although the coverage in Fig. 5 is slightly higher at 1.6 ML, the expected reconstructions are comparable [13]. At elevated temperature but in the absence of an As flux, the 3-fold periodicity is clearer than in the case of Fig. 4(a) but without the additional discrete features present in Fig. 4(b) after quenching. Some streaking remains in the image.
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Fig. 4. Two 25 nm filled-states STM images of the InAs WL on GaAs(0 0 1). In (a) the InAs coverage is 1.5 ML and the growth temperature is 400 1C; the image was obtained using in situ STM [11]. In (b), the InAs coverage is 1.2 ML and the growth temperature was 350 1C; the image was obtained by quenched in vacuo STM. The cross-sections below are taken from image (a) as indicated by the white lines and show 2 , 3 and 4 spacings [11].
4. Discussion
Fig. 5. A 25 nm 25 nm STM image of 1.6 ML InAs on GaAs(0 0 1), obtained in situ at 300 1C but without an incident As flux. Clear 3-fold periodicity is visible.
The streaking in the upper half of image (a) in Fig. 2 illustrates one of the additional difficulties with imaging under the As4 flux. During typical MBE growth with excess As, a surface population of mobile As molecules is present [7]. Contamination of the tip, due to As molecules becoming unstably adsorbed near the tip apex, is the most likely cause of the streaking. These effects were rarely observed without the incident As4 flux. It was found that higher scan speeds (tip motion rates X2000 nm s 1 compared to more typical values of hundreds of nm s 1) helped to prevent the onset of streaking and instability [10]. However, such high scan speeds make it more difficult to obtain atomic resolution. After lengthy scanning under As4, the tip would normally become very contaminated and no resolution of the substrate could be obtained. This could be alleviated by scanning at relatively high tunneling current and bias voltage well away from the area of interest. The resulting combination of tip heating and field desorption was sufficient to remove the contamination and restore resolution. Subsequent large-scale scans of the areas used for such procedures showed the presence of significant deposits, sometimes up to tens of nm in height, confirming
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that material had been desorbed from the tip. A final point regards shadowing of As4 beam by the tip itself. Since the final radius of curvature of the STM tips may be some tens of nm, one would expect the tip itself to block the incoming fluxes over some lateral distance relating to this radius and the polar incidence angles of the incoming beams. In the case of InAs–GaAs(0 0 1), the incoming As flux stabilizes the (1 3) reconstruction. In As-deficient conditions, this reverts to a (4 2)/c(8 2) In-stable phase which is very easy to observe with STM due to its characteristic, highly anisotropic domain structure [14]. No such As-deficient phases were observed in the present experiments, indicating that the migration lengths of surface As molecules are sufficient to maintain As-rich reconstructions even in areas where the incident beam might be blocked by the tip. Despite the additional tip cleaning required while imaging under As4, the in situ STM images shown in Figs. 2 and 3 are comparable to in vacuo results, i.e. the nanometer-scale resolution is essentially unimpaired. This confirms the ability of in situ STM to follow the development of epitaxial structures on this scale. It is useful to assess the possible effects of annealing under the As4 flux on the morphology of the QD-covered surface (note that this in situ STM–MBE (STMBE) experiment is distinct from in vacuo high-temperature STM because of the group V overpressure). Subject to kinetic hindrance through slow adatom exchange between QDs, one expects to observe ripening of the QD array simply due to the energetic balance between surface area and strain relaxation. In practical terms, surfaces are often prepared ex situ by As-decapping [15] which may involve extended annealing at 300 1C. In the case of InAs–GaAs QDs, ex situ AFM is also widely used on large MBE platens which may cool rather slowly. The present results indicate that the WL and QD morphology is quite stable over timescales of at least hundreds of seconds at 400 1C. Quenched mode microscopy involving a similar or smaller post-growth thermal load should therefore faithfully reflect the as-grown morphology. We now turn to high resolution imaging of the WL (Figs. 4 and 5). The absence of As flux allows stable in situ imaging at much lower scan speeds, which in turn tends to produce improved resolution (the scan speed for the images in Figs. 5 and 6 is approximately 70 nm s 1). The overall structure of the WL surface observed by the in situ STM is in agreement with previous RHEED and STM work [13,16]. In particular, the 3-fold periodicity is dominant but there is considerable structural disorder. Atomic resolution is regularly achieved in quenched-mode STM, but images of the InAs WL normally show the surface covered in small discrete features which have been assigned to In adatoms [13,14]. This is the case even for ‘static’ surfaces annealed under As flux for a long time before quenching [13], implying that there is a substantial equilibrium population of In adatoms which are frozen in by the quench. Mobile surface species are expected to cause some streaking in the in situ images, and indeed this is normally observed.
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Despite the additional difficulty imposed the high scan speeds, high resolution images under As flux typically show more severe streaking (Fig. 4(a)) than in vacuum-anneal conditions (Fig. 5). This is probably due to a combination of mobile As molecular surface species [7] and In adatoms. In the case of the annealed surface without As flux, streaking is still present but is less pronounced, indicating the presence of mobile In adatoms only. Experimental measurements of total surface stress [17] indicate that a significant fraction of deposited In in the WL does not exist in a surface phase supporting the epitaxial stress. Recent STMBE results also indicate that not all deposited In is incorporated into the lattice [18]. Furthermore, theoretical work by Ishii et al. [19] and Kratzer et al. [20] indicates that In adatoms should be highly mobile on InAs/GaAs(0 0 1) thin film surfaces. In the latter study, the alloyed In0.66Ga0.34As (2 3) structure was shown to be stable, while the activation energies derived for In migration on this surface were derived as just 0.13 eV in the [1¯ 1 0] direction and 0.29 eV along the [1 1 0] direction. The present observations (Fig. 5) are consistent with the existence of a population of In adatoms which are mobile even at 300 1C and migrate on a rather disordered (1 3) reconstructed WL surface. Although the morphology of steps, pits and QDs appears quite stable over several hundred seconds when annealed at 400 1C under As flux (Fig. 2), it is possible to observe a small degree of step motion when the WL is annealed at 300 1C without As flux. In Fig. 6, several step edges are visible. The upper step edges lie mainly parallel to the rows of the reconstruction, while the lower left step edges cut across the reconstruction rows. Careful comparison of the images shows that some step erosion has taken place going from image (a) to image (b). This is highlighted by the white arrows. However, the step edges parallel to the reconstruction rows near the top of the images are hardly affected. This implies that As desorption may take place preferentially at steps running along the [1 1 0] azimuth (first-principles calculations would be useful in elucidating this point). Further comparison of the images shows that the detailed pattern of the segments and trenches of the (1 3) reconstruction is most strongly altered in the vicinity of the step erosion between images (a) and (b). This is highlighted by a white circle on the images. This is consistent with the release of additional In adatoms by the As desorption process which migrate and locally alter the reconstruction domains. It is probable that an incident As flux suppresses this process and leads to a more stable step morphology (Fig. 2). 5. Conclusions A comparison between in situ and in vacuo STM–MBE studies has been made, concentrating on the InAs– GaAs(0 0 1) system. Despite the necessity of extra STM tip treatment when scanning under the As4 flux, it has been shown that nanometer-scale imaging is comparable to the
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Fig. 6. Two 50 nm 50 nm STM image of 1.6 ML InAs on GaAs(0 0 1), obtained in situ under continuous annealing at 300 1C without an As flux. The location of each image is the same (subject to some drift) and the image acquisition time was 400 s for each image. Image (b) was obtained directly after image (a). The 3-fold periodicity of the surface reconstruction is visible as approximately horizontal striped segments and several step edges are also present. The white arrows and circle highlight changes between the two images.
conventional in vacuo methods. Furthermore, comparisons of atomic scale imaging with or without As flux strongly support the presence of a substantial population of mobile In adatoms at temperatures as low as 300 1C, and the presence of additional mobile As-related species under As flux. The disordered (1 3) reconstruction of the WL can be clearly observed using in situ imaging. Both QDs and the terrace/island/pit structures of the WL appear to be stable when annealed under As4 flux at 400 1C over timescales of several hundred seconds. In the absence of As flux, some step erosion has been observed at 300 1C on the InAs WL. The preferential site for As re-evaporation appears to be at step edges running parallel to the [1 1 0] direction and this leads to the release of In adatoms which locally modify the (1 3) reconstruction. Acknowledgments The authors thank Dr. Nobuyuki Koguchi at the National Institute for Materials Science in Japan (NIMS) for sincere encouragement and support. The STMBE work was performed at NIMS (Figs. 2, 3a, and 4a) and continues with a new modified STMBE system at University of Tokyo (Figs. 5 and 6) assisted by IT program of the Ministry of Education, Culture, Sports, Science and Technology of the Japanese Government. References [1] B.G. Orr, C.W. Snyder, M. Johnson, Rev. Sci. Instrum. 62 (1991) 1400–1403. [2] P. Geng, J. Marquez, L. Geelhar, J. Platen, C. Stezer, K. Jacobi, Rev. Sci. Instrum. 71 (2000) 504–508.
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