Ripening of self-organized InAs quantum dots

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1 Available online at Physica E 21 (2004) Ripening of self-organized InAs quantum dots K. Potschke a;,l.muller-kirsch a, R. Heitz a;1, R.L. Sellin a, U.W. Pohl a, D. Bimberg a, N. Zakharov b, P. Werner b a Institut fuer Ferstkorperphysik, Technische Universitat Berlin, Hardenbergstr. 36, Berlin 10623, Germany b Max-Planck-Institut fur Mikrostrukturphysik, Weinberg 2, Halle 06120, Germany Abstract The temporal evolution of the size and the shape of self-organized InAs/GaAs quantum dots (QDs) grown using MOCVD is investigated. During a growth interruption after the deposition of the QD material a ripening process is observed, where some QDs grow at the expense of other QDs. A multimodal distribution of the QD ground-state transition energies is observed and attributed to QDs diering in height by entire numbers of atomic monolayers. This distribution is used to track the evolution of the QD ensemble during the growth interruption more detailed. A shape transition from very at, truncated-pyramid-like QDs to higher, more pyramidal QDs is suggested. An additional antimony ux at the end of the growth interruption leads to an accelerated ripening resulting in a signicant red shift of the QD luminescence, which is explained by the surfactant properties of antimony on InAs.? 2003 Elsevier B.V. All rights reserved. PACS: Ta; Hc; Hb Keywords: Quantum dot; Monolayer splitting; Ripening; MOCVD 1. Introduction Self-organized quantum dots (QDs) are promising for applications as active medium in optoelectronic devices [1]. The evolution of the QD ensemble towards an equilibrium state [2] featuring a narrow QD size distribution is important for most QD-based optoelectronic devices. Such an equilibrium state implies the existence of a limited island size which can be explained either by thermodynamic arguments [2,3] or by strain-related adatom diusion barriers around the Corresponding author. Tel.: ; fax: address: konst@physik.tu-berlin.de (K. Potschke). 1 Deceased. QD bases, leading to kinetic self-limitation of the QD size [4,5] or both. Growth interruptions of varying duration previous to the deposition of the GaAs cap layer on top of InAs/GaAs QDs were performed, which allow the QD ensemble to develop to a potentially stable conguration via ripening. Ripening has previously been reported for QDs in dierent material systems [6 8]. For dierent durations of the growth interruption, the QD ensemble shows a spectral modulation of the photoluminescence intensity, which is attributed to QD subsets, which dier in height by integer numbers of atomic monolayers (ML). This ML-splitting allows us to monitor the ripening process of the QD ensemble in detail. An additional antimony ux added at the end of the growth interruption (GRI) leads to an /$ - see front matter? 2003 Elsevier B.V. All rights reserved. doi: /j.physe

2 acceleration of the QD ripening, due to surfactant properties of antimony. K. Potschke et al. / Physica E 21(2004) Experimental The samples were grown in an Aixtron 200/4 horizontal quartz glass reactor at 100 mbar on GaAs(0 0 1) substrates. TMIn, TMGa, TMAl, and the alternative arsenic precursor tbuas were used as precursors. The carrier gas was H 2. The deposition of a 300 nm thick GaAs buer layer was followed by 60 nm Al 0:6 Ga 0:4 As, and 90 nm GaAs at 625 C. Subsequently, the growth temperature was reduced to 485 C for the deposition of the InAs layer. About 1:8 ML InAs were deposited at a nominal V/III ratio of 1.5 at a growth rate of about 0:4 ML=s. After the deposition of InAs a GRI of variable duration was introduced to allow QD formation and ripening. All precursor supply was switched o during the GRI, except for experiments with antimony at the end of the GRI. The QDs were then capped at 485 C with 5 nm GaAs. During the deposition of the following 50 nm GaAs the temperature was ramped up to 600 C. Finally, 20 nm Al 0:33 Ga 0:67 As were deposited as diusion barrier for photoexcited charge carriers and 10 nm GaAs was followed to prevent oxidation. The samples were structurally characterized by dark-eld transmission electron microscopy (TEM), using the strain-sensitive (2 2 0) reection. The morphology of the cap layer was checked by atomic force microscopy (AFM). Photoluminescence measurements (PL) were performed using the 514:5 nm line of an Ar + laser for excitation, a 0:3 m double-grating monochromator and a LN 2 -cooled Ge pin-diode for detection using lock-in technique. 3. Experimental results We investigated the temporal evolution of QD formation, using a series of samples for which the duration of the GRI was varied. Without GRI only wetting-layer (WL) related luminescence is observed. For a GRI of 2 s and up to the longest GRI of 270 s, the WL luminescence disappears at low excitation density and the QD-related luminescence dominates the spectrum (Fig. 1). Fig. 1. PL spectra of InAs QD samples for which the duration of the GRI was varied recorded at room temperature with low excitation density. For a GRI of 270 s a signicant decrease of PL intensity can be observed at room temperature. The cap layer of this sample shows hillocks in AFM, while for all other samples the cap layer is smooth on a monolayer scale. This dierence can be explained by the larger volume of the QDs after a GRI of 270 s. Obviously some islands exceed the critical size for dislocation formation. Due to the faster growth of dislocated islands as compared to coherent QDs [9] these dislocated islands become large, plastically relaxed islands (clusters). Note that the presence of dislocated islands strongly aects the evolution of the strained QDs due to an increased material transfer to the dislocated islands. For the intended study of Stranski Krastanow growth mode, only GRIs up to 90 s are therefore useful. With increasing duration of the GRI, the QD luminescence exhibits a red shift, as previously observed [10]. The red shift slows down for longer GRI but can be observed during the whole GRI (see inset of Fig. 1). The reason for the red shift of the QD luminescence is an increase of the average QD volume. The transition energy of the WL, taken from PLE measurements (not shown), stays constant, independent of the applied GRI. This means that the thickness of the WL does not change during the applied GRI and therefore the WL does not supply the material for the QDs. The volume increase of the QDs can thus only be ascribed to material exchange between dierent QDs. This implies that during the GRI, a

3 608 K. Potschke et al. / Physica E 21(2004) Fig. 2. Normalized integrated ground-state luminescence of QD samples with dierent durations of GRI in the low and high excitation density regime recorded at 10 K. Fig. 4. PL spectra of InAs QD samples for which the duration of the GRI was varied recorded at 10 K with low excitation density. Fig. 3. Plan view TEM images of InAs QD samples with 5 s GRI (left) and 90 s GRI (right). certain fraction of QDs dissolve to allow other QDs to grow. At low excitation density ( 5W=cm 2 ), the integrated QD ground-state luminescence at 10 K (Figs. 2 and 4) hardly changes during 270 s GRI whereas at higher excitation density ( 500 W=cm 2 ), the integrated ground-state luminescence decreases monotonically (Fig. 2). Due to saturation of the QD luminescence at higher excitation densities, the integrated QD luminescence of the ground state reects the QD densities. Therefore, the data suggests a decrease of the QD density during the GRI. This is exemplarily shown by plan-view TEM images for two samples (Fig. 3). The QD density drops by a factor of about 2 from 5 to 90 s GRI. No clusters could be found in both samples. Fig. 4 shows the PL spectra of the QD samples for which the duration of the GRI was varied at 10 K. Several maxima can be distinguished in the quantum-dot-related luminescence bands. The PL spectra were recorded with a low excitation density of 5W=cm 2, suggesting that no excited states are responsible for the features in the PL spectrum. The PL lines must stem from structurally dierent subsets of the QD ensemble. The maxima originate from different subsets of QDs with dierent heights exceeding the WL thickness with sharp upper and lower interfaces. The explicit contrast between subsets of QDs, which are dierenced in height by 1 ML, is related to the large dierence in z-connement for these very at QDs. For a detailed discussion of the monolayer splitting see Heitz et al. [11]. The dierent peaks are broadened by a distribution of the island widths. A height of 2 ML including the WL for the QDs belonging to the highest energy maximum at 1:35 ev is estimated from the transition energy of a 2 ML InAs quantum well in a GaAs matrix [12]. The height of the other maxima is subsequently increased by integer numbers of monolayers. The observed ML splitting allows a more detailed discussion of the evolution of the QDs. It shows clearly that parts of the QD ensemble increase in height during the GRI. For this increase in QD height, an evolution of QD shape from very at, truncated pyramids to higher, more pyramidal like QDs is proposed. Such a shape transition connected to an increase in QD volume should lead to the observed red shift of the QD luminescence. So we expect that the transferred material leads to an increase in QD volume with a

4 K. Potschke et al. / Physica E 21(2004) simultaneous evolution in QD shape. This is similar to theoretical predictions from Moll et al. [3]. Up to a GRI of 270 s the evolution of QDs towards larger volume and the decrease of the QD density is comparable with an Ostwald ripening-like behavior of the QD ensemble. Ostwald ripening denotes an increasing particle volume when a critical particle size is exceeded, featuring growth of larger particles at the expense of smaller ones. For a detailed discussion see Kamins et al. [6] and references therein. On the other hand, strain relaxation is also a driving force towards larger QDs due to favorable strain relaxation in larger QDs. Furthermore, one has also to take into account that the strain eld surrounding the QDs is proposed to hinder further growth of larger QDs [4,5] and a shape transition also inuences the growth process. Therefore the presented kind of ripening should not be confused with Ostwald ripening and termed dierently. It is not excluded that QDs can develop to an equilibrium state, which however could not be shown due to the formation of clusters in MOCVD. On the other hand the shown material transfer between different QDs is a precondition for a thermodynamic equilibrium. Antimony is known to decrease the surface energy and therefore impacts the formation of QDs and the evolution towards thermodynamic equilibrium. Adding an antimony ux after 90 s GRI without any metalorganic ux at this point to the rather stable QD ensemble leads to a signicant red shift of about 100 mev (for the larger Sb ux) of the QD luminescence (Fig. 5). A signicant incorporation of Sb through an As Sb exchange is excluded from other investigations [13] and from the low probability according to Yano et al., [14]. Furthermore, with the smaller Sb ux new maxima on the high-energy side of QD luminescence appear. These maxima are assumed to come from dissolving QDs. It is assumed that Sb inuences the growth in two ways. Firstly, the decrease of the surface energy changes the equilibrium conditions. The critical point for QDs to dissolve or to grow is shifted by the presence of antimony, which leads to larger QDs. The dissolution of QDs seems to be comparable to the presented formation of QDs in this report. Secondly, Sb may alter the surface kinetics in a way that accelerates the evolution towards equilibrium. Fig. 5. InAs QD samples with 90 s GRI and additional 5 s GRI with dierent antimony uxes. In conclusion, we have investigated the temporal evolution of QD ensembles during a GRI applied after deposition of the QD material. Some QDs grow at the expense of other QDs, which dissolve. Furthermore, we have shown that the evolution of the QDs is connected to a shape transition towards more pyramidal like QDs. Adding antimony aects the equilibrium conditions and impacts the temporal evolution of the QDs accordingly. The authors acknowledge support by the Deutsche Forschungsgemeinschaft in the framework of SFB 296 and by the E.U. project DOTCOM. References [1] V.A. Shchukin, N.N. Ledentsov, D. Bimberg, Epitaxy of Nanostructures, Springer, Berlin, [2] V.A. Shchukin, D. Bimberg, Rev. Mod. Phys. 71 (1999) [3] N. Moll, M. Scheer, E. Pehlke, Phys. Rev. B 58 (1998) [4] A.-L. Barabasi, Appl. Phys. Lett. 70 (1997) [5] D.E. Jesson, G. Chen, K.M. Chen, S.J. Pennycook, Phys. Rev. Lett. 80 (1998) [6] T.I. Kamins, G. Medeiros-Ribeiro, D.A.A. Ohlberg, R.S. Williams, J. Appl. Phys. 85 (1999) [7] A. Raab, G. Springholz, Appl. Phys. Lett. 77 (2000) [8] F.M. Ross, J. Terso, R.M. Tromp, Phys. Rev. Lett. 80 (1998) 984. [9] J. Drucker, Phys. Rev. B 48 (1993) [10] F. Heinrichsdor, A. Krost, M. Grundmann, D. Bimberg, F. Bertram, J. Christen, A. Kosogov, P. Werner, J. Cryst. Growth 170 (1997) 568.

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