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GaAs quantum wires incorporated in planar Bragg microcavities
Journal of Crystal Growth 207 (1999) 161}173
Organometallic chemical vapor deposition of V-groove InGaAs/GaAs quantum wires incorporated in planar Bragg microcavities C. Constantin*, E. Martinet, A. Rudra, K. Leifer, F. Lelarge, G. Biasiol, E. Kapon Department of Physics, Swiss Federal Institute of Technology - EPFL, 1015 Lausanne, Switzerland Received 7 April 1999; accepted 14 June 1999 Communicated by A. Zangwill
Abstract We report the fabrication and the optical properties of dense arrays of strained InGaAs/GaAs V-groove quantum wires (QWRs) embedded in wavelength size planar Bragg microcavities, made using a two-step organometallic chemical vapor deposition (OMCVD). Growth front evolution and top surface morphology of GaAs on the corrugated substrate were investigated as a function of the growth temperature (550}6253C) and the pitch of the V-groove grating (3}0.25 lm). Based on these studies, a growth temperature of 5503C and a grating pitch of 0.25 lm were selected to achieve sizes compatible with simultaneous quantum con"nement for electron (+10 nm wide wires) and photon (+0.25 lm thick cavities) states. Reference arrays of uniform, nanometer-size, crescent-shaped, InGaAs QWRs with densities up to 4QWRs/lm were realized, exhibiting a narrow (8 meV) and intense emission from one-dimensional excitonic states at low temperature. Similar QWRs were then successfully embedded in planar microcavities, showing brighter (]50) emission with a strongly reduced linewidth (1 meV) due to resonant coupling between the wire emission and the microcavity modes. ( 1999 Elsevier Science B.V. All rights reserved. PACS: 81.15.Gh; 68.65.#g; 42.60.Da Keywords: Self-ordered OMCVD growth; Strain; Quantum wires; Microcavities
1. Introduction Low-dimensional (LD) semiconductor nanostructures in which the charge carrier motion is restricted in two- or three-dimensions, such as
* Corresponding author. Tel.: #41-21-693-3387; fax: #4121-693-5480. E-mail address: christophe.constantin@ep#.ch (C. Constantin)
quantum wires (QWRs) or quantum dots (QDs), show interesting physical properties and are potentially useful for applications in novel electronic and optoelectronic devices [1,2]. In particular, the reduced dimensionality in one-dimensional (1D) QWRs leads to unique features such as a peaked free-particle density of states, intrinsic polarization anisotropy [3], increased optical nonlinearities and enhanced excitonic e!ects [4,5]. The incorporation of strain in such structures, by combining lattice mismatched materials, brings another degree of
0022-0248/99/$ - see front matter ( 1999 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 0 2 4 8 ( 9 9 ) 0 0 3 6 1 - 9
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freedom in tailoring the material's properties and in further improving device performances (e.g., reduction of threshold current and higher optical gain for lasers) [6]. In particular, high strain levels can be accommodated without defect formation by these structures, due to their limited lateral dimension [7]. Moreover, combining electronic and photonic con"nement, by incorporating LD nanostructures in optical resonators with dimensions of the order of the emission wavelength (microcavities), allows to tailor the light}matter interaction (spontaneous emission pattern and rate, coupling strength, etc.) by controlling simultaneously electron and photon states [8]. Novel phenomena in quantum optics and important technological applications have been predicted and demonstrated for quantum wells (QWs) embedded in microcavities, for example, strong light}matter coupling and high e$ciency LEDs [9]. Microcavities incorporating quantum structures of lower dimensionality will open new directions in this area. In particular, coupling between 1D electron states and 2D photon states can be obtained by inserting an array of QWRs into a planar microcavity. Theoretical predictions have already been made for such structures, showing dramatic modi"cations in the exciton}polariton dispersion [10], polarization-dependent absorption, emission and exciton radiative lifetime [11], as well as cooperative e!ects between the wires (mode coupling) [12] and reduced lasing thresholds due to larger spontaneous emission coupling factors [13]. Unfortunately, few experimental data on these structures are available due to the extreme di$culty involved in their fabrication [14,15]. On one hand, the observation of 1D quantum e!ects requires QWRs with small sizes, high uniformity in terms of size and composition, and defect free interfaces. On the other hand, the incorporation of the QWRs in a sub-micrometer planar cavity, with an accurate layer thickness and composition control in order to yield a resonant coupling between electron and photon states, is another major obstacle for realizing such structures. Various schemes have been developed to fabricate QWRs. Recently, unstrained T- and V-shaped GaAs/AlGaAs QWRs, obtained, respectively, by the cleaved-edge overgrowth method and growth
on pre-patterned substrates, have exhibited unambiguous 1D features [16}18]. The latter technique relies on in situ seeded self-ordered growth at sites pre-de"ned by V-grooves etched in the substrate. The resulting wires are defect-free and uniform thanks to a strong driving force towards the selfordered phase [19]. The recently reported similar self-ordering behavior of GaAs grown on V-groove substrates by low-pressure OMCVD [20] makes it possible to prepare GaAs templates for producing high-quality strained InGaAs/GaAs V-groove QWRs. Moreover, this approach should allow the quick planarization of the V-grooves, necessary for embedding strained wires in a planar microcavity, if su$ciently shallow corrugations are employed. We report here on growth studies of InGaAs/GaAs V-groove QWRs aimed at identifying the growth parameters needed for realizing onewavelength (j) planar Bragg microcavities incorporating high-quality strained wires. A microcavity QWR structure is then successfully demonstrated. The paper is organized as follows: in Section 2, we describe the experimental approach chosen and the subsequent stringent speci"cations faced. In Section 3, we establish the experimental feasibility of the structure by investigating and optimizing the growth of GaAs on V-grooved-substrates. In Sections 4 and 5, we present evidences for the high structural and optical properties achieved for arrays of strained QWRs and for microcavity-embedded arrays of strained QWRs, respectively.
2. Experimental approach and structural requirements 2.1. General approach and requirements A sketch of the structure we aim at is shown in Fig. 1. It consists of V-groove InGaAs/GaAs QWRs located within a j-thick GaAs spacer (j+0.25 lm is the low-temperature QWR emission wavelength in the material), sandwiched between two planar AlAs/AlGaAs highly re#ective distributed Bragg re#ectors (DBRs). Note that a growth interruption is necessary for corrugating the GaAs spacer prior to the deposition of the wires.
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Fig. 1. Sketch of a j planar Bragg microcavity incorporating an array of V-groove QWRs. The structure is obtained in a threestage process: (1) growth of an AlAs/AlGaAs bottom DBR mirror (2) etching of the V-groove array onto the mirror (3) regrowth of the InGaAs/GaAs QWRs and of an AlAs/AlGaAs top DBR mirror.
For the 1D character of the V-groove QWRs to be clearly revealed, the growth of GaAs on the corrugated surface needs to produce a uniform template of su$ciently small size (+10 nm), before depositing the InGaAs QWRs. The interfaces should then be defect free, thus making the surface preparation before the regrowth process crucial. Moreover, the wires should exhibit good uniformity in terms of size variations (less than a few %) and composition variation (less than 1%). Other structural requirements are speci"c to the microcavity structure. Care should be taken not to damage the AlAs-containing bottom mirror while etching the V-grooves onto the thin (+130 nm) GaAs spacer. Planarization of the grating after the deposition of the QWRs must occur within the allowed 250 nm thickness of the j GaAs spacer, and the planarized surface should exhibit a smooth morphology in order to obtain a top mirror with high quality, #at interfaces. Finally, resonant coupling between the wires and the microcavity requires suitable control of the spectral and geometric features of the structure. The QWRs should be located at an antinode of the optical "eld inside the cavity and their emission wavelength should match that of the cavity mode. The last two conditions have been quantitatively evaluated, using a standard transfer matrix method [21] to calculate the optical "eld
Fig. 2. Standard transfer matrix calculation of (a) the optical "eld distribution in the cavity, normalized to the vacuum "eld intensity and of (b) the re#ectivity spectrum for exact thicknesses (solid line), !4% in spacer thickness (short dashed line), and #3% in DBRs thickness (long dashed line); a schematic Gaussian PL emission is also represented.
distribution in the structure and its re#ectivity spectrum. The calculations, depicted in Fig. 2, demonstrate the tight speci"cations on the accuracy of the thickness of the di!erent layers: (i) if located within $20 nm away from the antinode of the optical "eld the QWRs couple to 80% of the "eld maximum; (ii) a $1% thickness error in the spacer or in the DBR layers causes a $+3 nm shift of the
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cavity mode wavelength. As a result, a 3% growth rate control on the nonplanar surface is necessary to reach resonant coupling, using the spatial growth rate nonuniformities in the OMCVD reactor and the temperature as post growth degrees of freedom to spectrally tune the microcavity QWR structure. 2.2. Technical details To investigate the growth of GaAs on the corrugated substrate, we have monitored its growth front evolution by inserting a series of thin InGaAs marker layers in GaAs. MIn Ga As (1.2 nm)/ 0.15 0.85 GaAs (8.4 nm)N]8-MIn Ga As (1.2 nm)/GaAs 0.15 0.85 (28 nm)N]7 layer sequences were grown on a semi-insulating (1 0 0)GaAs substrate patterned with an array of V-grooves oriented along the [0 1 11 ] direction. Note that the thicknesses mentioned correspond to nominal thicknesses obtained on a planar substrate. After optimization of the growth parameters, strained In Ga As/GaAs 0.15 0.85 QWRs in arrays of di!erent densities (0.3/lm}4/lm) were realized with nominal thicknesses between 3 and 25 nm. Finally, complete microcavity QWR structures were realized in a three-stage process: (i) a 25 pairs AlAs/Al Ga As DBR bottom 0.07 0.93 mirror with a 125 nm-thick GaAs cap layer was "rst grown; (ii) the sample was then removed from the growth reactor and the cap layer patterned with a V-groove array; (iii) a GaAs spacer with embedded In Ga As/GaAs QWRs and a 0.15 0.85 20 pairs AlAs/Al Ga As DBR top mirror 0.07 0.93 were regrown. The InGaAs/GaAs rather than the InGaAs/AlGaAs material system was preferred for the spacer in order to avoid contamination by oxidation before the regrowth step, as well as to obtain a fast planarization of the corrugated surface. The AlAs/AlGaAs material system was chosen to yield mirrors that are transparent at the emission wavelength of the QWRs. All samples were grown at 20 mbar in a horizontal OMCVD reactor using trimethylgallium (TMGa), trimethylindium (TNIn), trimethylaluminium (TMAl), arsine and Pd-puri"ed H as 2 a carrier gas. The susceptor was rotated at about 60 rpm under a total reactor #ow of 6 l/mn. The AlAs/Al Ga As DBRs were grown at 7503C 0.07 0.93
with growth rates of 0.3 and 0.7 nm/s under a V/III ratio of 600 and 145, respectively. The partial pressures of TMAl were 83 and 3.5 mPa, while that of TMGa was 345 mPa. During the growth of the In Ga As/GaAs spacer material, the partial 0.15 0.85 pressure of TMIn was 26 mPa and that of TMGa was 152 mPa while arsine was kept at 33 Pa. Moreover, the substrate temperature during the growth of the spacer was found to be one of the critical parameters to optimize in order to meet the microcavity requirements described earlier. A detailed growth study for temperatures ranging from 550 to 6253C is carried out in Section 2. We fabricated gratings of V-grooves with periodicities ranging from 3 to 0.25 lm using conventional or holographic photolithography. First, a two-level mask, consisting of a negative photoresist (AZ 5200) on top of a Si N dielectric antiref3 4 lection coating (ARC) serving also as an adhesion promoter, was patterned with a rectangular grating. The grating was de"ned in the photoresist by means of holographic photolithography in the case of submicrometer pitch corrugations, using a UV (364 nm) argon laser and a two beams interferometer system, and then transferred into the ARC by CF reactive ion etching (RIE). There4 after, anisotropic wet chemical etching in an H SO : H O : H O solution (1 : 8 : 80 by volume) 2 4 2 2 2 yielded sharp V-shaped grooves delimited by quasi-M1 1 1NA sidewalls in GaAs. Finally, the mask was removed and the substrate cleaned in a buffered HF solution.
3. Feasibility studies and growth parameters optimization 3.1. Cross-sectional growth front evolution To address the issue of the growth of a j-thick GaAs spacer including InGaAs QWRs in view of microcavity applications, we have monitored the evolution of the GaAs growth front on a V-grooved surface using thin InGaAs markers. Growth runs at di!erent temperatures, ¹ "550, 575, 600 and 4 6253C, were implemented on substrates patterned with a 0.25 lm pitch V-groove array. Fig. 3 shows a dark-"eld transmission electron microscope
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Fig. 3. Dark-"eld TEM cross-section of GaAs grown at 6253C and 5503C on a 0.25 lm pitch V-groove grating. Thin InGaAs markers, appearing with a darker contrast, were inserted to monitor the evolution of the GaAs growth front, composed of di!erent facets delimited by white dots. Circles "tting the self-limiting GaAs growth front in the groove and planarization of the grating predicted by the geometrical model (dashed lines) are also shown.
(TEM) cross-section of the samples grown at the extreme temperatures used (625 and 5503C). InGaAs markers appear with a darker contrast. Di!erent facets (delimited by white dots on the micrograph) are formed during the "rst stage of the GaAs growth: a central (1 0 0) facet in the groove, surrounded by two inner M3 1 1NA facets, and, further out, by high index sidewalls; the grooves are then connected by a (1 0 0) ridge surrounded by two outer M3 1 1NA facets. The length of the facets evolves, and after su$cient material is deposited, only the top (1 0 0) facet remains and the groove is planarized. We hereafter quantitatively describe and discuss this evolution of the growth front at di!erent growth temperatures. In the groove, the center (1 0 0) and M3 1 1NA facets extension varies in the "rst stage of GaAs growth before reaching a self-limiting value; moreover, once perturbed by an InGaAs marker, the
same self-limiting pro"le is quickly recovered. Those features are evidenced in Fig. 4, where the radius of curvature o of the GaAs growth front G!A4 before the deposition of the di!erent markers are shown for the samples grown at di!erent temperatures. The measured o corresponds to the G!A4 radius of the circle de"ned by the center of the bottom (1 0 0) facet and the intersection between this facet and the two surrounding M3 1 1NA facets, as shown in Fig. 3 at the "fth marker. o inG!A4 creases slightly at 6253C and decreases slowly at 6003C, in contrast to the fast decrease observed in the case of ¹ "575 and 5503C. A self-limiting 4 is reached after depositing about value o4G!A4 40}50 nm GaAs nominal thickness (i.e. "ve decreases for demarkers) for all samples. o4G!A4 creasing temperatures: o4- +50, 42, 16 and G!A4 10 nm at ¹ "625, 600, 575 and 5503C, respective4 ly. Note that the initial pro"le of the groove is
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Fig. 4. Evolution of the GaAs growth front radius of curvature o as a function of the nominal GaAs thickness (markers), for G!A4 di!erent growth temperatures (550, 575, 600 and 6253C). The solid lines are the calculated evolutions using the analytic model developed in Ref. [19].
slightly di!erent for the di!erent samples investigated (di!erent initial o in Fig. 4); nevertheless, G!A4 a stable pro"le is obtained in all cases due to the strong driving force towards the self-ordered phase of the system. Moreover, when the self-limiting pro"le is perturbed by an InGaAs marker, o4- is G!A4 re-established within at most 10 nm (i.e. one marker period) for all temperatures. A similar behavior has been reported for Al Ga As grown on V0.47 0.53 grooved substrates; however, a growth temperature of 7003C was necessary in that case to obtain sub10 nm self-limiting radius of curvature [19]. In Fig. 5a, we show the GaAs growth front height from the initial groove surface, at the bottom (H , 7 "lled markers) and at the ridge (H , empty markers), 3 as a function of the nominal thickness, for ¹ "6253C (dashed line}circles) and ¹ "5503C 4 4 (solid line}squares). As a reference, we also present the usual linear dependence of the growth front height evolution on a planar surface (dashed}dotted line). The slope of the curve H versus nominal 7 thickness (full markers) and the one of H versus 3 nominal thickness (empty markers) are, respective-
Fig. 5. (a) GaAs height from the bottom of the V-groove above the groove (H , "lled markers) and above the ridge (H , empty 7 3 markers) at 5503C (solid line, squares) and 6253C (dashed line, circles) as a function of the nominal GaAs thickness; planarization points are marked by an arrow. GaAs height on a planar surface is shown as reference (dashed}dotted line). (b) Right scale: measured growth rate anisotropy R between the sidewalls and the ridge (squares) and exponential law of the temperature dependence of R (solid line). Left scale: normalized planarization height H /K as a function of the temperature (dashed lines) for 71 two groove geometries, obtained by a geometrical model taking into account the measured growth rate anisotropy R only.
ly, larger and smaller than the slope of the growth on a planar surface (dashed}dotted line), which indicates that more material is deposited in the groove than on the ridge, leading to the planarization of the corrugated surface. Notice that planarization occurs faster for the sample grown at low temperature than for the one grown at high temperature: a thickness from the bottom of the groove H of 215 nm (about 182 nm in nominal 71
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thickness) was necessary to reach planarization at ¹ "6253C, in contrast to only 148 nm (90 nm 4 nominal) at ¹ "5503C (arrows in Fig. 5a). 4 These observations above can be quantitatively understood when describing nonplanar growth in terms of both facet-dependent incorporation rate and subsequent surface di!usion of adatoms [22]. Incorporation rate anisotropy on the di!erent facets forming the V-groove originates mainly from local surface chemistry (bondings and steps density) which determines a facet-dependent activation energy for the decomposition of the precursors. In OMCVD with trimethyl precursors, the incorporation rate is typically maximum on the high-index sidewalls of the grooves and lower on the ridge (1 0 0) facets in this geometry [23]. Indeed, the measured ratio R between the growth rate on the high index sidewalls and the (1 0 0) ridge at ¹ "625, 600, 575 and 5503C, depicted as full 4 squares in Fig. 5b, amounts to 1.75, 1.85, 1.86 and 2.5, respectively.1 The solid line in Fig. 5b represents an exponential law of the temperature dependence of R based on values measured on many samples grown at di!erent temperatures, discussed in Ref. [19]. As the facet-dependant decomposition of the precursors is likely to be thermally activated, a di!erence in the activation energy of the decomposition between the high index sidewalls and the (1 0 0) ridge facet of 0.2 eV is observed. Thus, a decrease in temperature leads to an increase in incorporation rate anisotropy. Note that an incomplete decomposition of the precursors on the (1 0 0) ridge at 5503C may contribute to the large anisotropy observed (2.5) [24]. In addition to the growth rate anisotropy, surface di!usion stems from the lateral gradient in the surface chemical potential, which is lower in the groove that a!ects the surface evolution [25]. Thus, the migration of adatoms results in `capillaritya #uxes towards the bottom of the groove, which increase with increasing surface curvature [25]. A decrease in temperature results in smaller #uxes, as the di!usion length of GaAs
1 The values of R given were measured for facet lengths much larger than the self-limiting radius of curvature at the corresponding temperature, in order to neglect capillarity-related growth rate variations.
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species ¸ is thermally activated. Indeed, a di!uG!A4 sion barrier for GaAs on such high index sidewalls of 1.9$0.3 eV has been experimentally observed, from which we can deduce values of 58, 41, 28 and 19 nm for ¸ at ¹ "625, 600, 575 and 5503C G!A4 4 [19]. The observed evolution of the growth front pro"le towards self-ordering in the groove can be reproduced using the analytic model developed in Ref. [19], based on anisotropic incorporation rate and surface di!usion as the only ingredients. The calculated evolution of the radius of curvature o at the di!erent temperatures is depicted by G!A4 solid lines in Fig. 4, taking into account the temperature dependence of the incorporation rate anisotropy and of the surface di!usion; the calculated values are in quantitative agreement with the measured ones for all temperatures. At 5503C, during the "rst stage of growth, the fast linear decrease of o between 0 and 30 nm in nominal thickness is G!A4 mainly governed by the incorporation rate anisotropy; as the pro"le sharpens up, capillarity #uxes increase and o decreases more slowly between G!A4 30 and 40 nm; "nally, an equilibrium between incorporation rate anisotropy that tends to sharpen the pro"le and surface migration that tends to broaden it is reached after 40 nm, yielding a selflimiting value o4- "10 nm. When this equilibG!A4 rium is perturbed after the deposition of an InGaAs marker, capillarity #uxes will adjust in order to stabilize the pro"le back to its self-limiting value o4- . When the growth temperature is increased, G!A4 the same behavior holds. However, due to the decrease in incorporation rate anisotropy and the increase in di!usion length, the initial decrease of o is less pronounced and the equilibrium point G!A4 is shifted towards a larger value of o4- . Note that G!A4 the analytic model used is valid when the di!usion length of adatoms is much smaller than the size of the sidewalls, which is not the case at 6253C (58 and about 100 nm, respectively); however, the calculated evolution still agrees with the measured one. The planarization of the groove after thick enough material is deposited originates both from the incorporation rate anisotropy, as the smaller rate on the (1 0 0) ridge facets results in a gradual expansion of these facets at the expense of the
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other ones, and to surface di!usion leading to mass transport towards the bottom of the groove. However, the observed faster planarization at low temperature (¹ "5503C) than at high temperature 4 (¹ "6253C), indicates that the intuitive mecha4 nism of thermal smoothening of the nonplanar surfaces due to mass transport [25] plays a minor role in the planarization process. Furthermore, the dominant role played by the anisotropic incorporation rate has been clari"ed by a geometrical model predicting the evolution of the extension of the top (1 0 0) facet neglecting surface di!usion (see dashed lines in Fig. 3). The planarization thickness from the bottom of the groove H normalized by the 71 grating pitch K, is then related to the growth rate anisotropy R and to the length of the groove ¸ by: 7 H /K"[R/(R!1)][¸ /(2K)]. The modeled tem71 7 perature dependence of H /K is plotted in Fig. 5b 71 (dashed lines) for two di!erent groove geometries (¸ /K"0.5 and 0.8). The observed reduction of the 7 planarization thickness upon reducing the temperature is well reproduced with this model. Note, however, that the absolute height is overestimated by about 10% in all cases, as it does not take into account the acceleration of the planarization process visible in Fig. 3 between the two last markers before planarization. This acceleration is probably due to the predominant role played by surface di!usion during the very last stage before planarization. 3.2. Top surface morphology In order to address the issue of the regrowth of a DBR mirror on a V-grooved substrate, we have studied the GaAs top surface morphology after planarization for samples grown at di!erent temperatures (¹ "550, 575, 600 and 6253C) on sub4 strates patterned with a 0.25 lm pitch grating, and on planar substrates for reference. Fig. 6 displays atomic force microscopy (AFM) images of the top &&&&&&&&&&&&&&&&&&&&&c Fig. 6. (4]4 lm) AFM images of the GaAs top surface morphology after planarization for a reference sample grown at 5503C on a planar substrate, and for samples grown at 5503C and 6253C on a substrate patterned with a 0.25 lm pitch V-groove array.
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surface after planarization for the reference sample grown at ¹ "5503C, for the patterned sub4 strate grown at ¹ "5503C, and for the patterned 4 substrate grown at ¹ "6253C, respectively. All 4 samples show atomic terraces of di!erent size and shape. The planar reference sample exhibits a step #ow mode of growth, with lm-long steps oriented along the substrate unintentional miss-orientation; note that islands have nucleated on the terraces, indicating that the di!usion length of GaAs at 5503C on (1 0 0) facets is on the order of 500 nm, a value which is consistent with previously reported ones [26]. Both patterned substrates show a stable quasi-two-dimensional mode of growth after smoothening of the initial grooves, with a temperature dependent residual roughness. The terraces have no preferential orientation in this case, as the substrate miss-orientation has been lost during the V-groove etching process, resulting in a hills-andvalleys like morphology. At ¹ "5503C, a typical 4 height variation of about 20 monolayers on a 4 lm length scale is found on the surface. At ¹ "6253C, 4 a height variation of about 10 monolayers is found on the same length scale. The reduced roughness at 6253C can be explained by a better healing of the substrate after lithography, as the surface di!usion length for GaAs increases with temperature [25]. Note that in all cases, the observed roughness (a few nm) amounts to a small fraction of the wavelength of the light in the material (+250 nm) only, and should therefore not a!ect the optical properties of a DBR mirror grown on top of such a surface. This has been con"rmed by re#ectivity measurements on two nominally identical DBR mirrors, one grown on the planar reference surface and the other grown on the planarized surface of the GaAs
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patterned substrate, which yielded the same peak re#ectivity value. In summary, the growth conditions have been optimized by reducing the grating pitch and the growth temperature to, respectively, K"0.25 lm and ¹ "5503C, in order to yield stable GaAs 4 radius of curvatures compatible with electron quantum con"nement (+10 nm) and planarization thicknesses compatible with photon quantum con"nement (+250 nm). Moreover, these growth conditions lead to DBR mirrors on planarized V-grooved substrates with smooth interfaces.
4. InGaAs QWRs Periodic arrays of strained In Ga As/GaAs 0.15 0.85 QWRs with a nominal thickness of 5.8 nm and with di!erent grating pitches K"3, 0.5 and 0.25 lm were grown at 5503C, in order to study the e!ect of lateral packing density on the optical properties of the wires. Another similar array of QWRs with a nominal thickness of 4.7 nm was grown and used as a reference for the microcavity embedded QWRs discussed in the next section. 4.1. Structural properties A typical cross-sectional view of a dense (K"0.25 lm) array of QWRs, 4.7 nm in nominal thickness, is depicted in the TEM micrograph of Fig. 7. The QWRs in the bottom of the V-grooves, connected by side and top QWs, are highly uniform thanks to the seeded self-organization involved in the fabrication scheme used. The radius at the QWR interfaces expands from o "10 nm G!A4
Fig. 7. Dark-"eld TEM cross-sectional view of an array of strained In Ga As/GaAs QWRs connected by side and top QWs, 0.15 0.85 grown by low pressure OMCVD at 5503C on a 0.25 lm V-grooved substrate.
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(lower GaAs interface) to o "23 nm (upper I/G!A4 InGaAs interface), due to the larger di!usion length of In species towards the bottom of the groove as compared with Ga species [27], resulting in the typical crescent shape of the QWRs. Those sharp pro"les yield QWRs with vertical and lateral thicknesses at the crescent center of 10.6 and 19 nm, respectively. The faceting of the upper QWR interface is less pronounced than in similar GaAs/ Al Ga As V-groove QWRs grown at 6503C 0.42 0.58 [17], possibly due to previously reported straininduced height modulation along the wire axis on the M3 1 1NA facets [28]. The darker contrast at the crescent center in Fig. 7 suggests a non-uniform indium distribution in the InGaAs layer, owing to the di!erent di!usion length of Ga and In adatoms. An In content of 0.2 at the center of the wires and in the top QWs as compared with a value of 0.1 in the side QWs was measured by electron energy loss spectroscopy on a similar wire of 25 nm in nominal thickness [29]. This signi"cant segregation is expected to further enhance the lateral quantum con"nement in these structures. Note that QWRs with a nominal In content of 0.2 and a vertical thickness at the crescent center of 38 nm did not show any dislocation in high-resolution TEM images, suggesting a large critical thickness for systems of "nite lateral width, as already reported [7]. 4.2. Optical properties The samples were characterized by photoluminescence (PL) and photoluminescence excitation (PLE) spectroscopy with polarization analysis of the emission. The measurements were performed at low temperature (15 K), in a back scattering con"guration, using linearly polarized light from a tunable titanium}sapphire laser. The excitation energy was chosen so as to excite carriers in the GaAs barrier. Typical excitation power densities of 5 W/cm2 were used unless speci"ed otherwise. The emission was dispersed by a 0.85 m double spectrometer and detected by a germanium detector using standard lock-in techniques. The spectral resolution was "xed at 0.6 meV. Fig. 8 shows the normalized PL spectra of an array of nominally 5.8 nm thick QWRs with K"3,
Fig. 8. Normalized 15K PL spectra of strained In Ga As/ 0.15 0.85 GaAs QWR structures with di!erent grating pitches (K"3, 0.5 and 0.25 lm).
0.5 and 0.25 lm, respectively. For all samples, the lowest energy peak around 1.31 eV originates from the QWRs; the 20 meV energy distribution between the samples is attributed to inadvertent variations in growth parameters from run to run. The second peak at higher energy, around 1.37 eV, is assigned to emission from the top QWs at the ridge region. Note that the energy separation between the top QWs and the QWRs emission increases when decreasing the grating pitch (31 meV for K"3 lm and 51 meV for K"0.25 lm). A third peak still at higher energy around 1.39 eV, only visible in the spectrum of the K"3 lm sample, originates from the side QWs; no such emission was found for K"0.5 and 0.25 lm. These peak assignments were con"rmed by spatially resolved cathodoluminescence, PL of selectively etched samples where the top and side QWs were removed, and polarization analysis of the emission. The observed luminescence is attributed to excitonic transitions, as they dominate the low-temperature recombination process in LD nanostructures [1]. The high quality of the strained QWRs is attested to by a full-width at half-maximum (FWHM) as low as 8 meV. Moreover, the QWR line exhibits a symmetrical shape, as the low-energy tail present for similar GaAs/Al Ga As V-groove QWRs is absent 0.3 0.7 [18]. This indicates a di!erent localization dynamics, due to changes in interface and alloy disorders
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in the case of InGaAs/GaAs wires. The QWR emission peak intensity relative to that of the top QW, QWR/QW , increases from 0.025 for 501 K"3 lm up to 2.4 for K"0.25 lm. This increase in relative intensity when the grating pitch is reduced cannot be accounted for by the geometry of the samples solely (length of the QWs). In the case of sub-micrometer pitch gratings, more e$cient carrier capture into the wires plays a dominant role as con"rmed by measuring the e!ects of temperature and power excitation on the PL spectra. When increasing the pump power, the QWR/QW emis501 sion intensity ratio decreases as the QWR emission saturates before the QW emission due to the smaller joint density of states; this decrease is more pronounced for K"0.25 lm. When increasing the temperature, the QWR/QW emission intensity 501 ratio increases, due to more e$cient carrier transport into the wires as a consequence of the increased carrier mobility at higher temperature; this increase is less pronounced for K"0.25 lm. Both trends point towards a more e!ective carrier transfer into the QWRs in the case of sub-micrometer pitch grating, as this value is on the order of the exciton di!usion length at low temperature. For comparison purpose, we show in Fig. 10 the polarization analysis of the emission (dashed lines) and the PLE spectrum (dashed}dotted line) of the dense (K"0.25 lmN array of QWRs with a 4.7 nm nominal thickness, which will serve as a reference sample for microcavity QWRs. The dependence of the PL spectrum on the orientation of the linear polarization (parallel or perpendicular to the wire axis) reveals a striking anisotropy: the QWR emission is preferentially polarized along the wire with a degree of linear polarization P"(I1!3!I1%31)/ (I1!3#I1%31)"33%. This anisotropy is related to the details of the valence subband mixing [18] and to the anisotropic three-dimensional strain distribution [30] in the wires. Note that the top QW emission does not show any polarization anisotropy. The PLE spectrum was recorded under a pump power of about 100 W/cm2. It exhibits three heavy hole-like excitonic resonances associated with laterally con"ned subbands. The subband separation (about 40 meV) is larger than that observed in similar unstrained V-groove QWRs [18]. The observed small Stokes shift
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(9 meV) con"rms the high uniformity of the wires.
5. Microcavity QWRs Complete j planar Bragg microcavities incorporating similar arrays of In Ga As QWRs 0.15 0.85 with K"0.25 lm could thus be fabricated. As already mentioned, spatial and spectral tuning between the QWR emission and the cavity mode require an accurate and reproducible control of the location and of the emission energy of the QWRs on one hand, and of the cavity mode energy on the other hand. For that purpose, compositions and growth rates were carefully calibrated. The growth rate on V-grooved surfaces was found to be rather sensitive to the exact pro"le of the grating, and to the geometry of the deposition (substrate size and position). As a result, the tuning of the cavity turned out to be delicate. The microcavity structure was designed to be resonant at low temperature. 5.1. Structural properties Special care was taken during the patterning process so as to keep the AlAs-containing bottom mirror unaltered, as well as to obtain defect free interfaces. Fig. 9a shows a cross-sectional scanning electron microscope image of the microcavity structure prior to the second growth. A typical TEM cross-sectional view of the complete structure is depicted in Fig. 9b. The array of QWRs, similar to the reference one, is located in between the bottom and the top mirrors. Planarization of the GaAs growth front after the wires occurred before the deposition of the top DBR mirror, yielding #at interfaces. In spite of the regrowth process, no degradation of the structural properties has been introduced; a high structural quality is achieved as neither the regrown GaAs interface nor defects are visible in TEM cross-sections of the structure. The high structural properties of the microcavity were con"rmed by room-temperature re#ectivity measurements: a cavity mode with a linewidth of about 1 meV was observed, yielding a cavity quality factor Q of 1300.
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Fig. 9. (a) Typical SEM cross-sectional view of the bottom DBR mirror patterned with a grating of V-grooves. (b) Typical dark "eld TEM cross-section of a l planar AlAs/Al Ga As Bragg microcavity structure incorporating an array of In Ga As/GaAs 0.07 0.93 0.15 0.85 QWRs with a 0.25 lm pitch.
Fig. 10. 15 K PLE spectrum of a reference QWR sample without mirrors (dashed}dotted line); 15 K PL spectra of the reference QWR sample (dashed line) and of a microcavity QWR sample (solid line) of slightly di!erent thickness, for emission linearly polarized parallel or perpendicular to the wire axis.
5.2. Optical properties Fig. 10 shows the PL spectra of the microcavity (solid lines) for light linearly polarized parallel or
perpendicular to the wire axis. The corresponding spectra of the reference sample, already discussed, are also shown for comparison (dashed lines). Two photon energy scales are used in the plot as the two samples have a slightly di!erent nominal thickness. The microcavity spectra exhibit an intense and narrow peak at the `barea QWR emission energy (1.37 eV), and a much weaker and broader one at the `barea QW energy (1.41 eV). Note that the energy separation between these two peaks is exactly the same as the one measured on the reference sample. The peak at 1.37 eV exhibits the narrow (1 meV) characteristic linewidth of the cavity mode, instead of the broad (8 meV) emission linewidth of the reference QWRs. Furthermore, the on-axis emission intensity is enhanced by a factor of about 50 with respect to the reference QWR sample. Moreover, both structures show exactly the same emission polarization anisotropy, the emission polarized along the wires being favored. We interpret those results as a clear redistribution of the broad QWR spontaneous emission in free space into the narrow mode of the planar cavity, resulting from the resonant spatial and spectral coupling between the QWRs 1D carriers and the microcavity 2D photon states [8]. The optical properties of the system as a function of the detuning between the wire emission and the
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cavity mode are investigated in greater details in Ref. [31]. 6. Conclusions We have fabricated periodic arrays of strained In Ga As/GaAs V-groove QWRs with dens0.15 0.85 ities up to 4 QWRs/lm, and incorporated such dense arrays in a j thick (j+0.25 lm) planar Bragg microcavity. Prior to the fabrication of such demanding structures in terms of size, composition and reproducibility, we have studied quantitatively the GaAs growth processes on [0 1 11 ]-corrugated (1 0 0) substrates. Cross-sectional TEM analysis of the growth front evolution and AFM top surface morphology characterization established the key role played by the growth temperature, which was optimized to 5503C in order to meet the severe requirements of a microcavity QWR structure. All samples displayed high structural properties in spite of the two growth steps involved in the microcavity fabrication. The reference QWRs exhibited 1D narrow (8 meV) and intense emission due to e$cient carrier capture. The microcavity QWRs showed brighter (]50) emission with strongly reduced linewidth (1 meV) resulting from the resonant coupling between the wires and the microcavity. This e$cient light emission control demonstrates the possibility for novel light sources, combining the attributes of both QWRs (1D-DOS, exciton stability and preferential polarization) and microcavities (control of emission spectrum, pattern and rate). Future work on these structures include further investigations of the optical properties speci"c to this system, such as the emission rate and pattern, the polariton dispersion or cooperative e!ects between the wires [12]. Furthermore, this structure is a good candidate for investigating weak and strong coupling between 1D electrons and 1D or 0D photon states, realized by etching narrow ridge- or pillarlike structures on top of the microcavity [32,33]. Acknowledgements We thank A. Condo for help with the TEM work. This work was supported by the Swiss Priority Program Optics II.
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References [1] C. Weisbuch, B. Vinter, Quantum Semiconductor Structures, Academic Press, Boston, 1991. [2] R. Cingolani, R. Rinaldi, Nuovo Cimento D 16 (9) (1993) 1. [3] U. Bockelmann, G. Bastard, Europhys. Lett. 15 (2) (1991) 215. [4] F. Rossi, E. Molinari, Phys. Rev. B 53 (24) (1996) 16462. [5] T. Ogawa, T. Takagahara, Phys. Rev. B 44 (15) (1991) 8138. [6] S.L. Chuang, Physics of Optoelectronic Devices, Wiley, New York, 1995. [7] S.C. Jain et al., Phys. Scripta T 69 (1997) 65. [8] E. Burstein, C. Weisbuch (Eds.), Con"ned Electrons and Photons: New Physics and Applications, Vol. 340, NATO ASI Series B, Plenum Press, New York, 1995. [9] J. Rarity, C. Weisbuch (Eds.), Microcavities and Photonic bandgaps: Physics and Applications, Vol. 324, NATO ASI Series E, Kluwer Academic Publishers, Dordrecht, 1996. [10] A.V. Kavokin, M.A. Kaliteevski, M.R. Vladimirova, Phys. Rev. B 54 (3) (1996) 1490. [11] A.V. Kavokin et al., Superlattices Microstruct. 23 (2) (1998) 389. [12] J.Y. Duboz, J. Phys. 7 (1997) 1693. [13] T. Baba et al., IEEE J. Quantum Electron 28 (5) (1992) 1310. [14] A. Chavez Pirson et al., Appl. Phys. Lett. 64 (14) (1994) 1759. [15] T. Arakawa et al., Appl. Phys. Lett. 64 (17) (1994) 2200. [16] T. Someya, H. Akiyama, H. Sakaki, Phys. Rev. Lett. 76 (16) (1996) 2965. [17] A. Gustafsson et al., Appl. Phys. Lett. 67 (25) (1995) 3673. [18] F. Vouilloz et al., Phys. Rev. Lett. 78 (8) (1997) 1580. [19] G. Biasiol, E. Kapon, Phys. Rev. Lett. 81 (14) (1998) 2962. [20] G. Biasiol et al., Appl. Phys. Lett. 71 (13) (1997) 1831. [21] P. Yeh (Ed.), Optical Waves in Layered Media, Wiley, New York, 1988. [22] D.G. Coronell, K.F. Jensen, J. Crystal Growth 114 (1991) 581. [23] S.H. Jones, L.S. Salinas, J. Crystal Growth 154 (1995) 163. [24] G.B. Stringfellow, Organometallic Vapor Phase Epitaxy: Theory and Practice, Vol. XVIII, 1989, Academic Press, Boston, 1989, p. 398. [25] W.W. Mullins, J. Appl. Phys. 28 (3) (1957) 333. [26] M. Shinohara et al., Appl. Phys. Lett. 65 (11) (1994) 1418. [27] M. Walther et al., Appl. Phys. Lett. 62 (18) (1993) 2170. [28] E. Martinet et al., Appl. Phys. Lett. 72 (6) (1998) 701. [29] K. Leifer et al., in: Institute of Physics Conference Series, Oxford University Press, Oxford, 1999. [30] M. Notomi et al., Phys. Rev. Lett. 80 (14) (1998) 3125. [31] C. Constantin et al., Phys. Rev. B 59 (12) (1999) 7809R. [32] A.I. Tartakovskii et al., Phys. Rev. B 57 (12) (1998) R6807. [33] J.M. GeH rard et al., Phys. Rev. Lett. 81 (5) (1998) 1110.
Organometallic chemical vapor deposition of V-groove InGaAs/GaAs quantum wires incorporated in planar Bragg microcavities C. Constantin*, E. Martinet, A. Rudra, K. Leifer, F. Lelarge, G. Biasiol, E. Kapon Department of Physics, Swiss Federal Institute of Technology - EPFL, 1015 Lausanne, Switzerland Received 7 April 1999; accepted 14 June 1999 Communicated by A. Zangwill
Abstract We report the fabrication and the optical properties of dense arrays of strained InGaAs/GaAs V-groove quantum wires (QWRs) embedded in wavelength size planar Bragg microcavities, made using a two-step organometallic chemical vapor deposition (OMCVD). Growth front evolution and top surface morphology of GaAs on the corrugated substrate were investigated as a function of the growth temperature (550}6253C) and the pitch of the V-groove grating (3}0.25 lm). Based on these studies, a growth temperature of 5503C and a grating pitch of 0.25 lm were selected to achieve sizes compatible with simultaneous quantum con"nement for electron (+10 nm wide wires) and photon (+0.25 lm thick cavities) states. Reference arrays of uniform, nanometer-size, crescent-shaped, InGaAs QWRs with densities up to 4QWRs/lm were realized, exhibiting a narrow (8 meV) and intense emission from one-dimensional excitonic states at low temperature. Similar QWRs were then successfully embedded in planar microcavities, showing brighter (]50) emission with a strongly reduced linewidth (1 meV) due to resonant coupling between the wire emission and the microcavity modes. ( 1999 Elsevier Science B.V. All rights reserved. PACS: 81.15.Gh; 68.65.#g; 42.60.Da Keywords: Self-ordered OMCVD growth; Strain; Quantum wires; Microcavities
1. Introduction Low-dimensional (LD) semiconductor nanostructures in which the charge carrier motion is restricted in two- or three-dimensions, such as
* Corresponding author. Tel.: #41-21-693-3387; fax: #4121-693-5480. E-mail address: christophe.constantin@ep#.ch (C. Constantin)
quantum wires (QWRs) or quantum dots (QDs), show interesting physical properties and are potentially useful for applications in novel electronic and optoelectronic devices [1,2]. In particular, the reduced dimensionality in one-dimensional (1D) QWRs leads to unique features such as a peaked free-particle density of states, intrinsic polarization anisotropy [3], increased optical nonlinearities and enhanced excitonic e!ects [4,5]. The incorporation of strain in such structures, by combining lattice mismatched materials, brings another degree of
0022-0248/99/$ - see front matter ( 1999 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 0 2 4 8 ( 9 9 ) 0 0 3 6 1 - 9
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freedom in tailoring the material's properties and in further improving device performances (e.g., reduction of threshold current and higher optical gain for lasers) [6]. In particular, high strain levels can be accommodated without defect formation by these structures, due to their limited lateral dimension [7]. Moreover, combining electronic and photonic con"nement, by incorporating LD nanostructures in optical resonators with dimensions of the order of the emission wavelength (microcavities), allows to tailor the light}matter interaction (spontaneous emission pattern and rate, coupling strength, etc.) by controlling simultaneously electron and photon states [8]. Novel phenomena in quantum optics and important technological applications have been predicted and demonstrated for quantum wells (QWs) embedded in microcavities, for example, strong light}matter coupling and high e$ciency LEDs [9]. Microcavities incorporating quantum structures of lower dimensionality will open new directions in this area. In particular, coupling between 1D electron states and 2D photon states can be obtained by inserting an array of QWRs into a planar microcavity. Theoretical predictions have already been made for such structures, showing dramatic modi"cations in the exciton}polariton dispersion [10], polarization-dependent absorption, emission and exciton radiative lifetime [11], as well as cooperative e!ects between the wires (mode coupling) [12] and reduced lasing thresholds due to larger spontaneous emission coupling factors [13]. Unfortunately, few experimental data on these structures are available due to the extreme di$culty involved in their fabrication [14,15]. On one hand, the observation of 1D quantum e!ects requires QWRs with small sizes, high uniformity in terms of size and composition, and defect free interfaces. On the other hand, the incorporation of the QWRs in a sub-micrometer planar cavity, with an accurate layer thickness and composition control in order to yield a resonant coupling between electron and photon states, is another major obstacle for realizing such structures. Various schemes have been developed to fabricate QWRs. Recently, unstrained T- and V-shaped GaAs/AlGaAs QWRs, obtained, respectively, by the cleaved-edge overgrowth method and growth
on pre-patterned substrates, have exhibited unambiguous 1D features [16}18]. The latter technique relies on in situ seeded self-ordered growth at sites pre-de"ned by V-grooves etched in the substrate. The resulting wires are defect-free and uniform thanks to a strong driving force towards the selfordered phase [19]. The recently reported similar self-ordering behavior of GaAs grown on V-groove substrates by low-pressure OMCVD [20] makes it possible to prepare GaAs templates for producing high-quality strained InGaAs/GaAs V-groove QWRs. Moreover, this approach should allow the quick planarization of the V-grooves, necessary for embedding strained wires in a planar microcavity, if su$ciently shallow corrugations are employed. We report here on growth studies of InGaAs/GaAs V-groove QWRs aimed at identifying the growth parameters needed for realizing onewavelength (j) planar Bragg microcavities incorporating high-quality strained wires. A microcavity QWR structure is then successfully demonstrated. The paper is organized as follows: in Section 2, we describe the experimental approach chosen and the subsequent stringent speci"cations faced. In Section 3, we establish the experimental feasibility of the structure by investigating and optimizing the growth of GaAs on V-grooved-substrates. In Sections 4 and 5, we present evidences for the high structural and optical properties achieved for arrays of strained QWRs and for microcavity-embedded arrays of strained QWRs, respectively.
2. Experimental approach and structural requirements 2.1. General approach and requirements A sketch of the structure we aim at is shown in Fig. 1. It consists of V-groove InGaAs/GaAs QWRs located within a j-thick GaAs spacer (j+0.25 lm is the low-temperature QWR emission wavelength in the material), sandwiched between two planar AlAs/AlGaAs highly re#ective distributed Bragg re#ectors (DBRs). Note that a growth interruption is necessary for corrugating the GaAs spacer prior to the deposition of the wires.
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Fig. 1. Sketch of a j planar Bragg microcavity incorporating an array of V-groove QWRs. The structure is obtained in a threestage process: (1) growth of an AlAs/AlGaAs bottom DBR mirror (2) etching of the V-groove array onto the mirror (3) regrowth of the InGaAs/GaAs QWRs and of an AlAs/AlGaAs top DBR mirror.
For the 1D character of the V-groove QWRs to be clearly revealed, the growth of GaAs on the corrugated surface needs to produce a uniform template of su$ciently small size (+10 nm), before depositing the InGaAs QWRs. The interfaces should then be defect free, thus making the surface preparation before the regrowth process crucial. Moreover, the wires should exhibit good uniformity in terms of size variations (less than a few %) and composition variation (less than 1%). Other structural requirements are speci"c to the microcavity structure. Care should be taken not to damage the AlAs-containing bottom mirror while etching the V-grooves onto the thin (+130 nm) GaAs spacer. Planarization of the grating after the deposition of the QWRs must occur within the allowed 250 nm thickness of the j GaAs spacer, and the planarized surface should exhibit a smooth morphology in order to obtain a top mirror with high quality, #at interfaces. Finally, resonant coupling between the wires and the microcavity requires suitable control of the spectral and geometric features of the structure. The QWRs should be located at an antinode of the optical "eld inside the cavity and their emission wavelength should match that of the cavity mode. The last two conditions have been quantitatively evaluated, using a standard transfer matrix method [21] to calculate the optical "eld
Fig. 2. Standard transfer matrix calculation of (a) the optical "eld distribution in the cavity, normalized to the vacuum "eld intensity and of (b) the re#ectivity spectrum for exact thicknesses (solid line), !4% in spacer thickness (short dashed line), and #3% in DBRs thickness (long dashed line); a schematic Gaussian PL emission is also represented.
distribution in the structure and its re#ectivity spectrum. The calculations, depicted in Fig. 2, demonstrate the tight speci"cations on the accuracy of the thickness of the di!erent layers: (i) if located within $20 nm away from the antinode of the optical "eld the QWRs couple to 80% of the "eld maximum; (ii) a $1% thickness error in the spacer or in the DBR layers causes a $+3 nm shift of the
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cavity mode wavelength. As a result, a 3% growth rate control on the nonplanar surface is necessary to reach resonant coupling, using the spatial growth rate nonuniformities in the OMCVD reactor and the temperature as post growth degrees of freedom to spectrally tune the microcavity QWR structure. 2.2. Technical details To investigate the growth of GaAs on the corrugated substrate, we have monitored its growth front evolution by inserting a series of thin InGaAs marker layers in GaAs. MIn Ga As (1.2 nm)/ 0.15 0.85 GaAs (8.4 nm)N]8-MIn Ga As (1.2 nm)/GaAs 0.15 0.85 (28 nm)N]7 layer sequences were grown on a semi-insulating (1 0 0)GaAs substrate patterned with an array of V-grooves oriented along the [0 1 11 ] direction. Note that the thicknesses mentioned correspond to nominal thicknesses obtained on a planar substrate. After optimization of the growth parameters, strained In Ga As/GaAs 0.15 0.85 QWRs in arrays of di!erent densities (0.3/lm}4/lm) were realized with nominal thicknesses between 3 and 25 nm. Finally, complete microcavity QWR structures were realized in a three-stage process: (i) a 25 pairs AlAs/Al Ga As DBR bottom 0.07 0.93 mirror with a 125 nm-thick GaAs cap layer was "rst grown; (ii) the sample was then removed from the growth reactor and the cap layer patterned with a V-groove array; (iii) a GaAs spacer with embedded In Ga As/GaAs QWRs and a 0.15 0.85 20 pairs AlAs/Al Ga As DBR top mirror 0.07 0.93 were regrown. The InGaAs/GaAs rather than the InGaAs/AlGaAs material system was preferred for the spacer in order to avoid contamination by oxidation before the regrowth step, as well as to obtain a fast planarization of the corrugated surface. The AlAs/AlGaAs material system was chosen to yield mirrors that are transparent at the emission wavelength of the QWRs. All samples were grown at 20 mbar in a horizontal OMCVD reactor using trimethylgallium (TMGa), trimethylindium (TNIn), trimethylaluminium (TMAl), arsine and Pd-puri"ed H as 2 a carrier gas. The susceptor was rotated at about 60 rpm under a total reactor #ow of 6 l/mn. The AlAs/Al Ga As DBRs were grown at 7503C 0.07 0.93
with growth rates of 0.3 and 0.7 nm/s under a V/III ratio of 600 and 145, respectively. The partial pressures of TMAl were 83 and 3.5 mPa, while that of TMGa was 345 mPa. During the growth of the In Ga As/GaAs spacer material, the partial 0.15 0.85 pressure of TMIn was 26 mPa and that of TMGa was 152 mPa while arsine was kept at 33 Pa. Moreover, the substrate temperature during the growth of the spacer was found to be one of the critical parameters to optimize in order to meet the microcavity requirements described earlier. A detailed growth study for temperatures ranging from 550 to 6253C is carried out in Section 2. We fabricated gratings of V-grooves with periodicities ranging from 3 to 0.25 lm using conventional or holographic photolithography. First, a two-level mask, consisting of a negative photoresist (AZ 5200) on top of a Si N dielectric antiref3 4 lection coating (ARC) serving also as an adhesion promoter, was patterned with a rectangular grating. The grating was de"ned in the photoresist by means of holographic photolithography in the case of submicrometer pitch corrugations, using a UV (364 nm) argon laser and a two beams interferometer system, and then transferred into the ARC by CF reactive ion etching (RIE). There4 after, anisotropic wet chemical etching in an H SO : H O : H O solution (1 : 8 : 80 by volume) 2 4 2 2 2 yielded sharp V-shaped grooves delimited by quasi-M1 1 1NA sidewalls in GaAs. Finally, the mask was removed and the substrate cleaned in a buffered HF solution.
3. Feasibility studies and growth parameters optimization 3.1. Cross-sectional growth front evolution To address the issue of the growth of a j-thick GaAs spacer including InGaAs QWRs in view of microcavity applications, we have monitored the evolution of the GaAs growth front on a V-grooved surface using thin InGaAs markers. Growth runs at di!erent temperatures, ¹ "550, 575, 600 and 4 6253C, were implemented on substrates patterned with a 0.25 lm pitch V-groove array. Fig. 3 shows a dark-"eld transmission electron microscope
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Fig. 3. Dark-"eld TEM cross-section of GaAs grown at 6253C and 5503C on a 0.25 lm pitch V-groove grating. Thin InGaAs markers, appearing with a darker contrast, were inserted to monitor the evolution of the GaAs growth front, composed of di!erent facets delimited by white dots. Circles "tting the self-limiting GaAs growth front in the groove and planarization of the grating predicted by the geometrical model (dashed lines) are also shown.
(TEM) cross-section of the samples grown at the extreme temperatures used (625 and 5503C). InGaAs markers appear with a darker contrast. Di!erent facets (delimited by white dots on the micrograph) are formed during the "rst stage of the GaAs growth: a central (1 0 0) facet in the groove, surrounded by two inner M3 1 1NA facets, and, further out, by high index sidewalls; the grooves are then connected by a (1 0 0) ridge surrounded by two outer M3 1 1NA facets. The length of the facets evolves, and after su$cient material is deposited, only the top (1 0 0) facet remains and the groove is planarized. We hereafter quantitatively describe and discuss this evolution of the growth front at di!erent growth temperatures. In the groove, the center (1 0 0) and M3 1 1NA facets extension varies in the "rst stage of GaAs growth before reaching a self-limiting value; moreover, once perturbed by an InGaAs marker, the
same self-limiting pro"le is quickly recovered. Those features are evidenced in Fig. 4, where the radius of curvature o of the GaAs growth front G!A4 before the deposition of the di!erent markers are shown for the samples grown at di!erent temperatures. The measured o corresponds to the G!A4 radius of the circle de"ned by the center of the bottom (1 0 0) facet and the intersection between this facet and the two surrounding M3 1 1NA facets, as shown in Fig. 3 at the "fth marker. o inG!A4 creases slightly at 6253C and decreases slowly at 6003C, in contrast to the fast decrease observed in the case of ¹ "575 and 5503C. A self-limiting 4 is reached after depositing about value o4G!A4 40}50 nm GaAs nominal thickness (i.e. "ve decreases for demarkers) for all samples. o4G!A4 creasing temperatures: o4- +50, 42, 16 and G!A4 10 nm at ¹ "625, 600, 575 and 5503C, respective4 ly. Note that the initial pro"le of the groove is
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Fig. 4. Evolution of the GaAs growth front radius of curvature o as a function of the nominal GaAs thickness (markers), for G!A4 di!erent growth temperatures (550, 575, 600 and 6253C). The solid lines are the calculated evolutions using the analytic model developed in Ref. [19].
slightly di!erent for the di!erent samples investigated (di!erent initial o in Fig. 4); nevertheless, G!A4 a stable pro"le is obtained in all cases due to the strong driving force towards the self-ordered phase of the system. Moreover, when the self-limiting pro"le is perturbed by an InGaAs marker, o4- is G!A4 re-established within at most 10 nm (i.e. one marker period) for all temperatures. A similar behavior has been reported for Al Ga As grown on V0.47 0.53 grooved substrates; however, a growth temperature of 7003C was necessary in that case to obtain sub10 nm self-limiting radius of curvature [19]. In Fig. 5a, we show the GaAs growth front height from the initial groove surface, at the bottom (H , 7 "lled markers) and at the ridge (H , empty markers), 3 as a function of the nominal thickness, for ¹ "6253C (dashed line}circles) and ¹ "5503C 4 4 (solid line}squares). As a reference, we also present the usual linear dependence of the growth front height evolution on a planar surface (dashed}dotted line). The slope of the curve H versus nominal 7 thickness (full markers) and the one of H versus 3 nominal thickness (empty markers) are, respective-
Fig. 5. (a) GaAs height from the bottom of the V-groove above the groove (H , "lled markers) and above the ridge (H , empty 7 3 markers) at 5503C (solid line, squares) and 6253C (dashed line, circles) as a function of the nominal GaAs thickness; planarization points are marked by an arrow. GaAs height on a planar surface is shown as reference (dashed}dotted line). (b) Right scale: measured growth rate anisotropy R between the sidewalls and the ridge (squares) and exponential law of the temperature dependence of R (solid line). Left scale: normalized planarization height H /K as a function of the temperature (dashed lines) for 71 two groove geometries, obtained by a geometrical model taking into account the measured growth rate anisotropy R only.
ly, larger and smaller than the slope of the growth on a planar surface (dashed}dotted line), which indicates that more material is deposited in the groove than on the ridge, leading to the planarization of the corrugated surface. Notice that planarization occurs faster for the sample grown at low temperature than for the one grown at high temperature: a thickness from the bottom of the groove H of 215 nm (about 182 nm in nominal 71
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thickness) was necessary to reach planarization at ¹ "6253C, in contrast to only 148 nm (90 nm 4 nominal) at ¹ "5503C (arrows in Fig. 5a). 4 These observations above can be quantitatively understood when describing nonplanar growth in terms of both facet-dependent incorporation rate and subsequent surface di!usion of adatoms [22]. Incorporation rate anisotropy on the di!erent facets forming the V-groove originates mainly from local surface chemistry (bondings and steps density) which determines a facet-dependent activation energy for the decomposition of the precursors. In OMCVD with trimethyl precursors, the incorporation rate is typically maximum on the high-index sidewalls of the grooves and lower on the ridge (1 0 0) facets in this geometry [23]. Indeed, the measured ratio R between the growth rate on the high index sidewalls and the (1 0 0) ridge at ¹ "625, 600, 575 and 5503C, depicted as full 4 squares in Fig. 5b, amounts to 1.75, 1.85, 1.86 and 2.5, respectively.1 The solid line in Fig. 5b represents an exponential law of the temperature dependence of R based on values measured on many samples grown at di!erent temperatures, discussed in Ref. [19]. As the facet-dependant decomposition of the precursors is likely to be thermally activated, a di!erence in the activation energy of the decomposition between the high index sidewalls and the (1 0 0) ridge facet of 0.2 eV is observed. Thus, a decrease in temperature leads to an increase in incorporation rate anisotropy. Note that an incomplete decomposition of the precursors on the (1 0 0) ridge at 5503C may contribute to the large anisotropy observed (2.5) [24]. In addition to the growth rate anisotropy, surface di!usion stems from the lateral gradient in the surface chemical potential, which is lower in the groove that a!ects the surface evolution [25]. Thus, the migration of adatoms results in `capillaritya #uxes towards the bottom of the groove, which increase with increasing surface curvature [25]. A decrease in temperature results in smaller #uxes, as the di!usion length of GaAs
1 The values of R given were measured for facet lengths much larger than the self-limiting radius of curvature at the corresponding temperature, in order to neglect capillarity-related growth rate variations.
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species ¸ is thermally activated. Indeed, a di!uG!A4 sion barrier for GaAs on such high index sidewalls of 1.9$0.3 eV has been experimentally observed, from which we can deduce values of 58, 41, 28 and 19 nm for ¸ at ¹ "625, 600, 575 and 5503C G!A4 4 [19]. The observed evolution of the growth front pro"le towards self-ordering in the groove can be reproduced using the analytic model developed in Ref. [19], based on anisotropic incorporation rate and surface di!usion as the only ingredients. The calculated evolution of the radius of curvature o at the di!erent temperatures is depicted by G!A4 solid lines in Fig. 4, taking into account the temperature dependence of the incorporation rate anisotropy and of the surface di!usion; the calculated values are in quantitative agreement with the measured ones for all temperatures. At 5503C, during the "rst stage of growth, the fast linear decrease of o between 0 and 30 nm in nominal thickness is G!A4 mainly governed by the incorporation rate anisotropy; as the pro"le sharpens up, capillarity #uxes increase and o decreases more slowly between G!A4 30 and 40 nm; "nally, an equilibrium between incorporation rate anisotropy that tends to sharpen the pro"le and surface migration that tends to broaden it is reached after 40 nm, yielding a selflimiting value o4- "10 nm. When this equilibG!A4 rium is perturbed after the deposition of an InGaAs marker, capillarity #uxes will adjust in order to stabilize the pro"le back to its self-limiting value o4- . When the growth temperature is increased, G!A4 the same behavior holds. However, due to the decrease in incorporation rate anisotropy and the increase in di!usion length, the initial decrease of o is less pronounced and the equilibrium point G!A4 is shifted towards a larger value of o4- . Note that G!A4 the analytic model used is valid when the di!usion length of adatoms is much smaller than the size of the sidewalls, which is not the case at 6253C (58 and about 100 nm, respectively); however, the calculated evolution still agrees with the measured one. The planarization of the groove after thick enough material is deposited originates both from the incorporation rate anisotropy, as the smaller rate on the (1 0 0) ridge facets results in a gradual expansion of these facets at the expense of the
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other ones, and to surface di!usion leading to mass transport towards the bottom of the groove. However, the observed faster planarization at low temperature (¹ "5503C) than at high temperature 4 (¹ "6253C), indicates that the intuitive mecha4 nism of thermal smoothening of the nonplanar surfaces due to mass transport [25] plays a minor role in the planarization process. Furthermore, the dominant role played by the anisotropic incorporation rate has been clari"ed by a geometrical model predicting the evolution of the extension of the top (1 0 0) facet neglecting surface di!usion (see dashed lines in Fig. 3). The planarization thickness from the bottom of the groove H normalized by the 71 grating pitch K, is then related to the growth rate anisotropy R and to the length of the groove ¸ by: 7 H /K"[R/(R!1)][¸ /(2K)]. The modeled tem71 7 perature dependence of H /K is plotted in Fig. 5b 71 (dashed lines) for two di!erent groove geometries (¸ /K"0.5 and 0.8). The observed reduction of the 7 planarization thickness upon reducing the temperature is well reproduced with this model. Note, however, that the absolute height is overestimated by about 10% in all cases, as it does not take into account the acceleration of the planarization process visible in Fig. 3 between the two last markers before planarization. This acceleration is probably due to the predominant role played by surface di!usion during the very last stage before planarization. 3.2. Top surface morphology In order to address the issue of the regrowth of a DBR mirror on a V-grooved substrate, we have studied the GaAs top surface morphology after planarization for samples grown at di!erent temperatures (¹ "550, 575, 600 and 6253C) on sub4 strates patterned with a 0.25 lm pitch grating, and on planar substrates for reference. Fig. 6 displays atomic force microscopy (AFM) images of the top &&&&&&&&&&&&&&&&&&&&&c Fig. 6. (4]4 lm) AFM images of the GaAs top surface morphology after planarization for a reference sample grown at 5503C on a planar substrate, and for samples grown at 5503C and 6253C on a substrate patterned with a 0.25 lm pitch V-groove array.
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surface after planarization for the reference sample grown at ¹ "5503C, for the patterned sub4 strate grown at ¹ "5503C, and for the patterned 4 substrate grown at ¹ "6253C, respectively. All 4 samples show atomic terraces of di!erent size and shape. The planar reference sample exhibits a step #ow mode of growth, with lm-long steps oriented along the substrate unintentional miss-orientation; note that islands have nucleated on the terraces, indicating that the di!usion length of GaAs at 5503C on (1 0 0) facets is on the order of 500 nm, a value which is consistent with previously reported ones [26]. Both patterned substrates show a stable quasi-two-dimensional mode of growth after smoothening of the initial grooves, with a temperature dependent residual roughness. The terraces have no preferential orientation in this case, as the substrate miss-orientation has been lost during the V-groove etching process, resulting in a hills-andvalleys like morphology. At ¹ "5503C, a typical 4 height variation of about 20 monolayers on a 4 lm length scale is found on the surface. At ¹ "6253C, 4 a height variation of about 10 monolayers is found on the same length scale. The reduced roughness at 6253C can be explained by a better healing of the substrate after lithography, as the surface di!usion length for GaAs increases with temperature [25]. Note that in all cases, the observed roughness (a few nm) amounts to a small fraction of the wavelength of the light in the material (+250 nm) only, and should therefore not a!ect the optical properties of a DBR mirror grown on top of such a surface. This has been con"rmed by re#ectivity measurements on two nominally identical DBR mirrors, one grown on the planar reference surface and the other grown on the planarized surface of the GaAs
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patterned substrate, which yielded the same peak re#ectivity value. In summary, the growth conditions have been optimized by reducing the grating pitch and the growth temperature to, respectively, K"0.25 lm and ¹ "5503C, in order to yield stable GaAs 4 radius of curvatures compatible with electron quantum con"nement (+10 nm) and planarization thicknesses compatible with photon quantum con"nement (+250 nm). Moreover, these growth conditions lead to DBR mirrors on planarized V-grooved substrates with smooth interfaces.
4. InGaAs QWRs Periodic arrays of strained In Ga As/GaAs 0.15 0.85 QWRs with a nominal thickness of 5.8 nm and with di!erent grating pitches K"3, 0.5 and 0.25 lm were grown at 5503C, in order to study the e!ect of lateral packing density on the optical properties of the wires. Another similar array of QWRs with a nominal thickness of 4.7 nm was grown and used as a reference for the microcavity embedded QWRs discussed in the next section. 4.1. Structural properties A typical cross-sectional view of a dense (K"0.25 lm) array of QWRs, 4.7 nm in nominal thickness, is depicted in the TEM micrograph of Fig. 7. The QWRs in the bottom of the V-grooves, connected by side and top QWs, are highly uniform thanks to the seeded self-organization involved in the fabrication scheme used. The radius at the QWR interfaces expands from o "10 nm G!A4
Fig. 7. Dark-"eld TEM cross-sectional view of an array of strained In Ga As/GaAs QWRs connected by side and top QWs, 0.15 0.85 grown by low pressure OMCVD at 5503C on a 0.25 lm V-grooved substrate.
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(lower GaAs interface) to o "23 nm (upper I/G!A4 InGaAs interface), due to the larger di!usion length of In species towards the bottom of the groove as compared with Ga species [27], resulting in the typical crescent shape of the QWRs. Those sharp pro"les yield QWRs with vertical and lateral thicknesses at the crescent center of 10.6 and 19 nm, respectively. The faceting of the upper QWR interface is less pronounced than in similar GaAs/ Al Ga As V-groove QWRs grown at 6503C 0.42 0.58 [17], possibly due to previously reported straininduced height modulation along the wire axis on the M3 1 1NA facets [28]. The darker contrast at the crescent center in Fig. 7 suggests a non-uniform indium distribution in the InGaAs layer, owing to the di!erent di!usion length of Ga and In adatoms. An In content of 0.2 at the center of the wires and in the top QWs as compared with a value of 0.1 in the side QWs was measured by electron energy loss spectroscopy on a similar wire of 25 nm in nominal thickness [29]. This signi"cant segregation is expected to further enhance the lateral quantum con"nement in these structures. Note that QWRs with a nominal In content of 0.2 and a vertical thickness at the crescent center of 38 nm did not show any dislocation in high-resolution TEM images, suggesting a large critical thickness for systems of "nite lateral width, as already reported [7]. 4.2. Optical properties The samples were characterized by photoluminescence (PL) and photoluminescence excitation (PLE) spectroscopy with polarization analysis of the emission. The measurements were performed at low temperature (15 K), in a back scattering con"guration, using linearly polarized light from a tunable titanium}sapphire laser. The excitation energy was chosen so as to excite carriers in the GaAs barrier. Typical excitation power densities of 5 W/cm2 were used unless speci"ed otherwise. The emission was dispersed by a 0.85 m double spectrometer and detected by a germanium detector using standard lock-in techniques. The spectral resolution was "xed at 0.6 meV. Fig. 8 shows the normalized PL spectra of an array of nominally 5.8 nm thick QWRs with K"3,
Fig. 8. Normalized 15K PL spectra of strained In Ga As/ 0.15 0.85 GaAs QWR structures with di!erent grating pitches (K"3, 0.5 and 0.25 lm).
0.5 and 0.25 lm, respectively. For all samples, the lowest energy peak around 1.31 eV originates from the QWRs; the 20 meV energy distribution between the samples is attributed to inadvertent variations in growth parameters from run to run. The second peak at higher energy, around 1.37 eV, is assigned to emission from the top QWs at the ridge region. Note that the energy separation between the top QWs and the QWRs emission increases when decreasing the grating pitch (31 meV for K"3 lm and 51 meV for K"0.25 lm). A third peak still at higher energy around 1.39 eV, only visible in the spectrum of the K"3 lm sample, originates from the side QWs; no such emission was found for K"0.5 and 0.25 lm. These peak assignments were con"rmed by spatially resolved cathodoluminescence, PL of selectively etched samples where the top and side QWs were removed, and polarization analysis of the emission. The observed luminescence is attributed to excitonic transitions, as they dominate the low-temperature recombination process in LD nanostructures [1]. The high quality of the strained QWRs is attested to by a full-width at half-maximum (FWHM) as low as 8 meV. Moreover, the QWR line exhibits a symmetrical shape, as the low-energy tail present for similar GaAs/Al Ga As V-groove QWRs is absent 0.3 0.7 [18]. This indicates a di!erent localization dynamics, due to changes in interface and alloy disorders
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in the case of InGaAs/GaAs wires. The QWR emission peak intensity relative to that of the top QW, QWR/QW , increases from 0.025 for 501 K"3 lm up to 2.4 for K"0.25 lm. This increase in relative intensity when the grating pitch is reduced cannot be accounted for by the geometry of the samples solely (length of the QWs). In the case of sub-micrometer pitch gratings, more e$cient carrier capture into the wires plays a dominant role as con"rmed by measuring the e!ects of temperature and power excitation on the PL spectra. When increasing the pump power, the QWR/QW emis501 sion intensity ratio decreases as the QWR emission saturates before the QW emission due to the smaller joint density of states; this decrease is more pronounced for K"0.25 lm. When increasing the temperature, the QWR/QW emission intensity 501 ratio increases, due to more e$cient carrier transport into the wires as a consequence of the increased carrier mobility at higher temperature; this increase is less pronounced for K"0.25 lm. Both trends point towards a more e!ective carrier transfer into the QWRs in the case of sub-micrometer pitch grating, as this value is on the order of the exciton di!usion length at low temperature. For comparison purpose, we show in Fig. 10 the polarization analysis of the emission (dashed lines) and the PLE spectrum (dashed}dotted line) of the dense (K"0.25 lmN array of QWRs with a 4.7 nm nominal thickness, which will serve as a reference sample for microcavity QWRs. The dependence of the PL spectrum on the orientation of the linear polarization (parallel or perpendicular to the wire axis) reveals a striking anisotropy: the QWR emission is preferentially polarized along the wire with a degree of linear polarization P"(I1!3!I1%31)/ (I1!3#I1%31)"33%. This anisotropy is related to the details of the valence subband mixing [18] and to the anisotropic three-dimensional strain distribution [30] in the wires. Note that the top QW emission does not show any polarization anisotropy. The PLE spectrum was recorded under a pump power of about 100 W/cm2. It exhibits three heavy hole-like excitonic resonances associated with laterally con"ned subbands. The subband separation (about 40 meV) is larger than that observed in similar unstrained V-groove QWRs [18]. The observed small Stokes shift
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(9 meV) con"rms the high uniformity of the wires.
5. Microcavity QWRs Complete j planar Bragg microcavities incorporating similar arrays of In Ga As QWRs 0.15 0.85 with K"0.25 lm could thus be fabricated. As already mentioned, spatial and spectral tuning between the QWR emission and the cavity mode require an accurate and reproducible control of the location and of the emission energy of the QWRs on one hand, and of the cavity mode energy on the other hand. For that purpose, compositions and growth rates were carefully calibrated. The growth rate on V-grooved surfaces was found to be rather sensitive to the exact pro"le of the grating, and to the geometry of the deposition (substrate size and position). As a result, the tuning of the cavity turned out to be delicate. The microcavity structure was designed to be resonant at low temperature. 5.1. Structural properties Special care was taken during the patterning process so as to keep the AlAs-containing bottom mirror unaltered, as well as to obtain defect free interfaces. Fig. 9a shows a cross-sectional scanning electron microscope image of the microcavity structure prior to the second growth. A typical TEM cross-sectional view of the complete structure is depicted in Fig. 9b. The array of QWRs, similar to the reference one, is located in between the bottom and the top mirrors. Planarization of the GaAs growth front after the wires occurred before the deposition of the top DBR mirror, yielding #at interfaces. In spite of the regrowth process, no degradation of the structural properties has been introduced; a high structural quality is achieved as neither the regrown GaAs interface nor defects are visible in TEM cross-sections of the structure. The high structural properties of the microcavity were con"rmed by room-temperature re#ectivity measurements: a cavity mode with a linewidth of about 1 meV was observed, yielding a cavity quality factor Q of 1300.
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Fig. 9. (a) Typical SEM cross-sectional view of the bottom DBR mirror patterned with a grating of V-grooves. (b) Typical dark "eld TEM cross-section of a l planar AlAs/Al Ga As Bragg microcavity structure incorporating an array of In Ga As/GaAs 0.07 0.93 0.15 0.85 QWRs with a 0.25 lm pitch.
Fig. 10. 15 K PLE spectrum of a reference QWR sample without mirrors (dashed}dotted line); 15 K PL spectra of the reference QWR sample (dashed line) and of a microcavity QWR sample (solid line) of slightly di!erent thickness, for emission linearly polarized parallel or perpendicular to the wire axis.
5.2. Optical properties Fig. 10 shows the PL spectra of the microcavity (solid lines) for light linearly polarized parallel or
perpendicular to the wire axis. The corresponding spectra of the reference sample, already discussed, are also shown for comparison (dashed lines). Two photon energy scales are used in the plot as the two samples have a slightly di!erent nominal thickness. The microcavity spectra exhibit an intense and narrow peak at the `barea QWR emission energy (1.37 eV), and a much weaker and broader one at the `barea QW energy (1.41 eV). Note that the energy separation between these two peaks is exactly the same as the one measured on the reference sample. The peak at 1.37 eV exhibits the narrow (1 meV) characteristic linewidth of the cavity mode, instead of the broad (8 meV) emission linewidth of the reference QWRs. Furthermore, the on-axis emission intensity is enhanced by a factor of about 50 with respect to the reference QWR sample. Moreover, both structures show exactly the same emission polarization anisotropy, the emission polarized along the wires being favored. We interpret those results as a clear redistribution of the broad QWR spontaneous emission in free space into the narrow mode of the planar cavity, resulting from the resonant spatial and spectral coupling between the QWRs 1D carriers and the microcavity 2D photon states [8]. The optical properties of the system as a function of the detuning between the wire emission and the
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cavity mode are investigated in greater details in Ref. [31]. 6. Conclusions We have fabricated periodic arrays of strained In Ga As/GaAs V-groove QWRs with dens0.15 0.85 ities up to 4 QWRs/lm, and incorporated such dense arrays in a j thick (j+0.25 lm) planar Bragg microcavity. Prior to the fabrication of such demanding structures in terms of size, composition and reproducibility, we have studied quantitatively the GaAs growth processes on [0 1 11 ]-corrugated (1 0 0) substrates. Cross-sectional TEM analysis of the growth front evolution and AFM top surface morphology characterization established the key role played by the growth temperature, which was optimized to 5503C in order to meet the severe requirements of a microcavity QWR structure. All samples displayed high structural properties in spite of the two growth steps involved in the microcavity fabrication. The reference QWRs exhibited 1D narrow (8 meV) and intense emission due to e$cient carrier capture. The microcavity QWRs showed brighter (]50) emission with strongly reduced linewidth (1 meV) resulting from the resonant coupling between the wires and the microcavity. This e$cient light emission control demonstrates the possibility for novel light sources, combining the attributes of both QWRs (1D-DOS, exciton stability and preferential polarization) and microcavities (control of emission spectrum, pattern and rate). Future work on these structures include further investigations of the optical properties speci"c to this system, such as the emission rate and pattern, the polariton dispersion or cooperative e!ects between the wires [12]. Furthermore, this structure is a good candidate for investigating weak and strong coupling between 1D electrons and 1D or 0D photon states, realized by etching narrow ridge- or pillarlike structures on top of the microcavity [32,33]. Acknowledgements We thank A. Condo for help with the TEM work. This work was supported by the Swiss Priority Program Optics II.
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