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Impact of InGa As Capping Layer in Impeding Indium Desorption from Vertically Coupled InAs/GaAs Quantum Dot Interfaces Binita Tongbram, Saumya Sengupta, and Subhananda Chakrabarti ACS Appl. Nano Mater., Just Accepted Manuscript • Publication Date (Web): 17 Jul 2018 Downloaded from http://pubs.acs.org on July 18, 2018

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ACS Applied Nano Materials

Impact of InxGa(1-x)As Capping Layer in Impeding Indium Desorption from Vertically Coupled InAs/GaAs Quantum Dot Interfaces Binita Tongbram, †, ‡ Saumya Sengupta, ‡ and Subhananda Chakrabarti ‡* †, ‡

Centre for Research in Nanotechnology & Science, Indian Institute of Technology Bombay, Mumbai 400076, Maharashtra, India. ‡

Department of Electrical Engineering, Indian Institute of Technology Bombay, Mumbai 400076, Maharashtra, India. *Email: [email protected]

KEYWORDS: Quantum Dots, Self-assembled InxGa(1-x)As capping layers, Interface, Photoluminescence’s Measurement, High-resolution Transmission Electron Microscopy, High-resolution X-Ray Diffraction.

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ABSTRACT This study describes the effect of thin GaAs spacer of 4.5 nm thickness in bilayer coupled InAs quantum dot (QD) heterostructure. Here, we report the first demonstration of InAs/GaAs QDs capped by self-assembled InxGa(1-x)As layers. Self-assembled InxGa(1-x)As layers were introduced into each intermediate layer across the interface of InAs QDs and the GaAs layer in a verticalcoupled bilayer QD (VCBQD) heterostructure to prevent indium desorption from the QDs. The change in indium content in the seed layer InAs QDs changes the self-assembly position and modifies the InxGa(1-x)As layer thickness. A theoretical approach was presented to study the formation of self-assembled InxGa(1-x)As layers at each strain-free layer. We showed that strain energy at the 2nd intermediate (εzz2) is greater than that at the 1st intermediate (εzz1) layer; εzz2 depends on the vertical strain channel length. The impact of the InxGa(1-x)As layer thickness on strain energy was studied using high-resolution transmission electron microscopy; shorter strain channel length was found to facilitate the formation of more relaxed and larger sized selfassembled InxGa(1-x)As layer in the active layer. This InxGa(1-x)As layer formed at the intermediate layer acts as a capping layer or a protective shield for the indium adatoms, preventing their desorption from the InAs QDs. Furthermore, we studied the thermal stability of the self-assembled InxGa(1-x)As layer by annealing the VCBQD samples at 700°C and 800°C. This aspect has been investigated for the first time ever in the study of coupling efficiency between the InAs QDs and InxGa(1-x)As capping layer. A high-resolution in-plane (2θχ/ϕ) reciprocal space mapping (RSM) technique provided the connection between the in-plane reciprocal lattice point of InAs QDs and InxGa(1-x)As layers and revealed the strain and coupling between them. InAs QDs fully covered with the self-assembled InxGa(1-x)As layer enhanced the PL intensity by 77% and had an activation energy of 467 meV. 2 ACS Paragon Plus Environment

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1. Introduction Recently, vertical-coupled (VC) InAs/GaAs quantum dot (QD) heterostructures grown by molecular beam epitaxy (MBE) under the Stranski–Krastanov (SK) growth mode have shown a tremendous impact on the performance of optoelectronic devices such as solar cells, lasers, light emitting diodes, infrared detectors, and single-photon emitters.1–3 In addition, the optical, structural, and electrical properties of three-dimensionally confined InAs QDs can be tuned over a wide range by designing various multi-stacked InAs QD heterostructures with different material compositions. The GaAs spacer layer thickness can also be tailored to achieve the desired emission wavelength.4 Consequently, optimizing the VC stacked structures with GaAs spacer layers grown under different growth parameters is a challenging research topic. It opens new pathways for the application of InAs/GaAs QD heterostructures in telecommunication wavelengths ranging 1.3–1.55 µm.5 Recently, Stephan et al. studied the response of selfassembled InAs/GaAs QDs to strong terahertz pulses tuned to the s-to-p transition.6 In addition, a successful investigation on the lateral spacing of quantum coupling between InAs QDs demonstrated the maximum critical limit of indium migration distance to be ~30 nm; this study also presented the growth model of self-assembled InAs/GaAs QDs showing significant changes in the surface self-diffusion rate of indium adatoms.7 A past study8 reported that, in contrast to the VC stacked structures, uncoupled structures of self-assembled InAs QDs have significantly smaller dots with high QD density, randomly distributed in the GaAs matrix with non-uniform size distribution. These uncoupled QD structures degrade the optical quality of InAs/GaAs QDbased devices. Over the past years, extensive research has been carried out to improve the VC stacked structures by methods such as (i) optimizing the GaAs spacer thickness,2 (ii) introducing a strain-reducing layer,2–4 (iii) varying the material composition of III–V semiconductor alloys of 3 ACS Paragon Plus Environment


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the strain-reducing layer, such as GaAs,9 InGaAs,10, 12 AlGaAs,3, 10 InAlGaAs,2, 4 GaAsSb,10, 11 and InGaAsSb,10 and (iv) varying the monolayer coverage for the seed layer InAs QDs in order to achieve a high degree of correlation of vertically aligned QDs with stronger quantum confinement potential. We have previously investigated the advantages of using quaternary capping over ternary capping in multi-stacked InAs QDs.2,

4

In addition, Mohanta et al. have

discussed the benefits of the carrier relaxation process for different spacer layer thicknesses in multi-stacked InAs QDs.9 However, the VC structures also have some shortcomings that weaken the QD coupling. Different types of dislocation in the VC structures have been reported based on its density, including (i) interface roughness, (ii) vertical threading dislocation, (iii) threaded dislocation in the {111} plane, (iv) V-shaped dislocation, (v) lateral overlapping of multiple dots, (vi) poor crystallinity, and (vii) indium segregation by creating point defect in the QD structure.13 Shiramine et al. reported two general mechanisms underlying threading dislocations during the growth of GaAs capping layers by MBE—(i) dislocation initiating from the region of the overlap between two large islands, and (ii) dislocation caused by a misfit between the relaxed InAs islands and the GaAs layer, which results in the threading of the epilayer in the {111} plane.14 In addition, interface roughness due to the deposition of the capping layer is another major reason for the leakage of indium adatoms from the surface of InAs islands. To achieve an emission wavelength of 1.3 µm with narrow spectra for the application of single photon sources at O-band telecom wavelength, uniform deposition of a capping layer and consistent material composition are desired. However, no research results are published yet on the uniformity of capping layers and the deformation of the intermediate layer across the interface of the InAs QDs and the GaAs spacer layer. Immense scope for advanced research on InAs QDs based structure is open to be explored for the enhancement of the performance of optoelectronics devices. 4 ACS Paragon Plus Environment

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Considering the aforementioned literature related to QD heterostructures, in this study, we design VC bilayer QD (VCBQD) heterostructures. These VCBQD heterostructures were formed by sandwiching a fixed thin GaAs spacer of thickness 4.5 nm between two QD layers to achieve strong coupling between the QDs. In the intermediate layer between the QDs and the GaAs spacer layer, a strong hydrostatic compressive strain is developed that assists the formation of self-assembled InxGa(1-x)As layers under zero phase-transition temperature (T1→2 = 0), which prevents thermal-assisted In/Ga intermixing. This strategy of growing seed layer QDs and GaAs spacer layers at zero phase-transition temperature helps avoid threading dislocation in the QD layers. The main advantages of the proposed VCBQD heterostructures incorporated with strainassisted InxGa(1-x)As layers include the following: (i) a high degree of homogeneity of QDs in terms of their size, shape, and uniformity is achieved; (ii) larger quantum confinement effect is realized; (iii) indium desorption from the QDs is mitigated; (iv) photoluminescence (PL) intensity of the QDs is enhanced; (v) narrow full-width half maxima is achieved; (vi) crystallinity of the interface is improved; and (vii) threading dislocation is eliminated in the heterostructure. This study is the first approach to construct a self-assembled InxGa(1-x)As layer at the interface of QDs and the GaAs spacer layer to control the leakage of indium adatoms, thereby increasing the resistance to lattice deformation of InAs QDs, i.e., the layer’s original size and shape are retained. We also examine for the first time the characteristics of InAs QDs that are fully and partially covered with the self-assembly InxGa(1-x)As layer. Furthermore, the variation of indium content in the self-assembled InxGa(1-x)As layer and its effect on the optical properties of the InxGa(1-x)As layer is yet to be investigated.



Therefore, the VCBQD structure we propose herein is expected to enhance not only the optical emission wavelength of 1300 nm for O-band fiber optics in telecommunication system but also the performance characteristics to acquire the high degree of monochromaticity, highly compact, high thermal stability for future development of semiconductors lasers. Ultimately, such a VCBQD structure of high-quality single-photon source can be explored for the main applications such as optical fiber communication, high-speed data transfer, quantum key distribution, quantum repeaters, and quantum information science. Furthermore, we study the thermal stability of self-assembled InxGa(1-x)As layers that cap the InAs QDs as well as the fluctuation in QD size distribution and density. In this paper, we present a detailed report of the investigation on the effect of indium composition in the strain-assisted InxGa(1-X)As layer and the variation in InxGa(1-X)As layer thickness by rapid thermal annealing (RTA) at high temperatures (upto 800°C). Annealing the VCBQD samples at a moderate temperature for a suitable duration can repair certain types of defects and dislocations, which improves the overall crystallinity of the samples. The proposed VCBQD heterostructure was investigated in detail by using high-resolution transmission electron microscopy (HRTEM), and high-resolution X-ray diffraction (HRXRD) scan, both from in-plane and out-plane at (004) reflections. We present a novel technique to analyze indium concentration, strain profiles, and broadening of the self-assembled InxGa(1-x)As layer by using low incident grazing angle (0.25°) in-plane (2ߠ߯/ߔ) reciprocal space mapping (RSM). We also present a technique to perform a high-contrast color mapping of HRTEM images at a resolution of 5 nm to investigate the structural properties of the self-assembled InxGa1-xAs layer. This HRTEM image contrast mapping is the first approach to determine the strain broadening width in and around the QDs in the active layer. The width of compressive strain at the top interface (between GaAs and InAs 6 ACS Paragon Plus Environment

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QDs) directly relates to the thickness of the InxGa1-xAs layer. Moreover, using the in-plane RSM technique, we studied the coupling efficiency in our VCBQD samples by using a contour diffuse ring plot of the diffracted reciprocal lattice points of the InAs QDs, InGaAs layer, and GaAs layer. In-plane RSM at the (004) plane allows the determination of the indium content, residual strain, relaxation, surface roughness, and misalignment of epilayer thin films. We hereby present a new equation to show that the strain at the 2nd intermediate layer (εzz2) is inversely proportional to the strain at the 1st intermediate layer (εzz1). Moreover, we calculated the compressive and tensile strains from the out- and in-plane HRXRD measurements and investigated their correlation with the indium content in the self-assembled InxGa1-xAs layer. We also present the advantages of having a thicker self-assembled InxGa1-xAs layer in the active QD layer formed at 30°C phase transition temperature (T1→2 = 30°C). We investigated the difference in broadening between InAs QDs that are fully and partially covered with the self-assembled InxGa1-xAs layer. Finally, we studied the function of the InxGa1-xAs layer as a protective shield due to the strong coupling effect in the VCBQD samples, which correlates with the enhancement of the PL intensity as well as its activation energy. In this paper, we report the VCBQD structure emitting high purity single-photon sources at the emission wavelength of 1.3 µm with three impactful key features such as high thermal stability, high compact, and high monochromaticity. Such optimised strain patterning techniques have been adopted for bilayer InAs QDs structure capped by strain mediated InxGa(1-x)As layer structure to achieve a narrow linewidth with high uniformity of QDs size and shape. The present results suggest the strong potential for InAs QDs fully capped with the self-assembled InxGa(1x)As

layer based single photon sources emitting in the O-band fiber optics telecom system.



2. Materials and Methods Two types of VCBQD heterostructures were grown on (001)-oriented semi-insulating GaAs substrates by using a Riber SYS14020 Epineat III–V MBE system equipped with different effusion cells (i.e., Ga, In, Al, Si, and As crackers). A schematic representation of the proposed VCBQD heterostructure is shown in Fig. 1. Samples were classified into samples A and B based on the monolayer coverage (ML) of the seed layer QDs (see the inset table of Fig. 1 for details). The native-oxide desorption process at the GaAs (100) surface was executed at 620°C to provide a deoxidized smooth growth surface. After the deoxidization, a 200-nm-thick GaAs buffer layer was grown at 580°C, which serves as the base for subsequent GaAs layers. Next, 136-nm-thick iGaAs layers were grown, following which the growth temperature was dropped from 580°C to 490°C to achieve good crystallinity of InAs QDs. The seed QDs monolayer coverage for samples A and B were 2.5 and 3.2 ML respectively. The growth rate of InAs QDs were calibrated in monolayer (ML) per second. The InAs QDs were grown at a low growth temperature of 490°C. After achieving a critical thickness of 1.8 ML, InAs QDs were self-assembled on the GaAs substrate surface. Then, 4.5 nm thin GaAs spacer layers were deposited on the randomly formed seed layer InAs QDs. Subsequently, fixed 3.2-ML active layer InAs QDs were grown at a low temperature of 460°C for samples A and B to achieve a higher aspect ratio in the active layer. Lastly, GaAs capping layers of thickness 50 nm were grown at 490°C for both samples to minimize thermal-assisted In/Ga intermixing and to reduce indium desorption from the core of the InAs QDs. The difference in the growth temperature of the seed layer (490°C) and the active layer QDs (460°C) was maintained at 30°C to enhance the uniformity of QDs. The QD growth rate was also maintained at a fixed low value of 0.03 ML/s D for both samples in order to achieve a low density of InAs QDs (1E10/cm2). The VCBQD samples were subjected to post8 ACS Paragon Plus Environment

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growth RTA at 700°C and 800°C for 30 s in an argon atmosphere using an AnnealSys (model AS-One 150) rapid thermal processor system, keeping all other parameters constant, to explore the thermal stability of the self-assembled InxGa(1-x)As capping layer. The thermal treatment was performed by using GaAs proximity capping to avoid arsenic desorption from the sample surface. To study the structural properties of the VCBQD samples, cross-sectional transmission electron microscopy (XTEM) was used. The samples were prepared by a destructive technique that involved cleaving, grinding, polishing, and Ar ion milling for high electron transparency. High-quality HRTEM images were achieved by dimpling down the region of interest to 20 µm. After dimpling the samples, they were loaded inside a precision ion milling system for further thinning. The samples were milled for six hours with a 4° Ar+ beam at a pressure of 9×10−5 Torr and an acceleration voltage of 4 kV until they were perforated. Finally, the angle of the incidence beam was reduced to 2° with 2.0 kV, and the samples were milled for a longer period till a reddish diffracted spot was visible at the edges of the perforated region of the samples. TEM images were obtained using a Tecnai G2, F30 electron microscope at a point resolution of 2 Å and a line resolution of 1 Å, with a magnification of 58x to 1 million x and operated at 300 kV. To study the optical properties of the VCBQD samples, the temperature-dependent (18– 300K) PL measurement was performed using a closed-cycle liquid-helium (He) cooled cryostat (CCS-100/204N). The six (6) VCBQDs samples were mounted inside the cold finger of a continuous flow liquid-He cryostat and excited using a 532-nm diode-pumped solid-state laser at a power density of 1.1 kW/cm2. The signal was dispersed by a 0.75-m triple-grating monochromator and detected with a liquid N2-cooled InGaAs array or a charge coupled device.



A chopper and a lock-in amplifier is used to eliminate background noise. The measured data were recorded using WinSpec software. The crystallinity, strain profiles, and indium content of the prepared VCBQD heterostructure are studied by HRXRD (Rigaku smart lab). The HRXRD system was equipped with a twobounce Ge (220) monochromator that provides an incident optics, which extracts the CuKα1 (1.5405 Å) ray generated using a 9-kW-rotating anode source. In general, in and out-plane (2ߠ/߱) scan, the intensities of the GaAs substrate dominate the epilayer film intensities. A symmetric out-plane (2ߠ/߱) scan records the X-ray diffracted from the crystal lattice plane that are parallel to the sample surface. In other words, the out-plane measurement is used to obtain the lattice constant perpendicular to the sample surface (a⏊). Furthermore, an in-plane HRXRD scan at (004) reflection was performed by considering the highest intensity value of ϕ. In the 2ߠ߯/ߔ scan, the X-ray intensity depends on the X-ray incidence grazing angle of the sample. We performed the wide ϕ scan from 0° to 360° to determine the highest intensity of ϕ. We computed ϕ with the highest intensity manually and fixed both the incident angle and the detector angle at a low grazing incidence angle of 0.25°. This in-plane HRXRD scan, called the cross-section scan, of the VCBQD samples achieves the following: (i) characterization of the diffraction from the thin self-assembly InxGa(1-X)As layer, (ii) characterization of the diffraction of the InAs QDs from each layer, (iii) minimizing of the substrate X-ray intensities, (iv) measure of the lattice plane normal to the surface, i.e., the in-plane lattice constant, and (v) determination of material composition and film crystallinity, and (vi) measure of the twist angle and screw dislocation density of the epilayer films. The details of the in-plane HRXRD characterization technique are described in our previous publication.14 The in-plane high-resolution X-ray reciprocal space mapping at (004) reflection was performed to study the strain field around the InAs QDs and to 10 ACS Paragon Plus Environment

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analyze the existence of the self-assembled InxGa

1-xAs

layer as well as its corresponding

broadening width (β). In-plane (2ߠ߯/ߔ) RSM at (004) reflection was conducted by detecting each diffracted 2θχ lattice point map for every small-step change in ϕ at the (004) plane by using a monochromatic set up with the scanning point detector. The broadening of ϕ along the (110) plane represents the crystal properties and that along 2θχ indicates the fluctuation in the horizontal d-spacing. From the broadening of the diffuse ring of the diffracted scattering vectors of the self-assembled InxGa 1-xAs layer, the twist angle and interface roughness were determined. In addition, the RSM scattering intensity as a function of the magnitude of the scattering vectors for the (004) reflection was studied for the VCBQD samples. 3. Results and discussion Williamson et al. presented the detailed theoretical study of the three-dimensional strain field in and around pyramidal InAs/GaAs QDs.15 In addition, Stier et al. presented a numerical model for pyramidal InAs QDs formed on a GaAs (001) substrate.16 Considering the theoretical background of InAs QDs, we present a novel theoretical approach to determine the strain field along the [001] and [110] directions at each intermediate junction (ij = 1, 2, 3…, depending on the InAs QD layer). We also studied the strain evolution that leads to the formation of the selfassembled InxGa(1-x)As layer. Fluctuation in the strain propagation and its coupling potential depends on the vertical inter-QD spacing. We previously reported the advantages of having a smaller vertical inter-QD spacing (VIDS) in the coupled bilayer QDs and its correlation with the strain energy along the z-direction in association with the dot density distribution in the coupled InAs QDs.17 Higher the strain energy that propagates, larger the QD size in the active layer. We determined the efficiency of strain energy propagation in terms of the vertical strain channel 11 ACS Paragon Plus Environment


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length. In other words, vertical strain channel length (Ls) is defined as the strain propagation length from the seed layer QDs to the subsequent QD layer. The force applied to the next InAs QD layer through the strained channel also relies on the broadening of QDs in the [110] directions. A combination of a short strain channel length and a large aspect ratio (height/base) of seed layer InAs QDs becomes a pivotal parameter to increase indium content and the thickness of the self-assembly InxGa1-x)As layer. Thus, the strain channel length is inversely proportional to the QD size. Smaller the Ls, greater the aspect ratio in the active layer InAs QDs. Based on the abovementioned theoretical study of the strain field profile in VCBQD heterostructures, we first consider the Hamiltonian strain equation of 3D confined InAs QDs,18 which is given by

1 6d 2 H e = −a(ε xx + ε yy + ε zz ) − 3b × [(Lx − L2 )ε xx + c. p.] − [{Lx Ly }ε xy + c. p.] , 3 3

(1)

where ߝij denotes the components of the strain in the xyz direction and L is the angular momentum operator. The parameter a depicts the hydrostatic deformation potential and b and d depict the shear deformation potentials corresponding to the strain of tetragonal and rhombohedral symmetries, respectively; c.p. denotes cyclic permutations with respect to the indices x, y, and z, and the quantities in the curly brackets indicate the symmetries produced: {LxLy} = {(LxLy + LyLx}. The formation of the self-assembly InxGa(1-x)As layer was explained in terms of hydrostatic strain εh and biaxial strain εb, which are given by

ε h = ε xx + ε yy + ε zz .

(2)

ε b = 2ε zz − ε xx + ε yy .

(3) 12 ACS Paragon Plus Environment

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The hydrostatic strain is highly compressive around the pyramidal InAs QDs at each intermediate strain junction. It is shown that an intermediate unit cell exists where the lattice constant ranges 0.56532 < x < 0.60583 nm, depending on the concentration of indium and gallium adatoms. Fig. 2 (a) shows the systematic representation of the unit cell architecture of the proposed VCBQD heterostructure, demonstrating the formation of the strain-assisted self-assembled InxGa(1-x)As layer across each intermediate junction (GaAs/InAs QDs). The GaAs spacer lattice plane is compressed in the z-direction concerning the InAs QDs (εzz) and expanded in the inplane xy directions (εxx, εyy). The lateral base of the GaAs layer lattice constant is constrained to adopt the top surface lattice constant of the strain InAs QDs. The lattice constants of QDs along the [110] and [001] directions differ according to the size and shape of the QDs. In the growth process, the GaAs lattice constant expands in the horizontal direction and try to match with the InAs lattice constant, resulting a tensile stress along the xy directions (εxx, εyy), thereby the expanded lattice constant width (or interplanar spacing, d) is similar to the InGaAs alloy unit crystal, hence leads to the formation of InGaAs layer at each intermediate junction between InAs QDs/GaAs interface. This intermediate junction further forms the strain-assisted self-assembled InxGa(1-x)As layer. The bilayer VCBQD structure has two intermediate strain junction layers; one at the seed layer QDs and the 4.5 nm GaAs spacer and the other at the active layer QDs and the 50-nm GaAs capping layer. The strain contribution of the 2nd intermediate layer relating to the lattice deformation of the GaAs layer shown in Fig. 2 (a) can be expressed as

ε zz 2 + 1 = α∆T1→ 2 +

1 K2

  aGaAs    , aGaAs  4.5 − (ε zz1 + 1)    13 ACS Paragon Plus Environment

(4)


where ∆T1→2 is the difference in the growth temperature between the active layer QDs and the GaAs capping layer. Eq. (4), thus affirms our explanation for the formation of the self-assembled InxGa(1-x)As layer. It shows that the strain arouses at the 2nd intermediate layer (εzz2) is greater than the strain at the 1st intermediate layer (εzz1). Therefore, the type of deformed layer formed after the release of the strain energy depends on the strain potential at each InAs QDs/GaAs interface. The selfdeformed lattice assembled after the strain release forms the self-assembled InxGa(1-x)As layer. The lattice constant at the interface also changes with the applied growth temperature of the GaAs capping layer, as shown in Fig. 2 (a). Because of the misfit between the InAs QDs and the overgrown GaAs capping layer, the InxGa(1-x)As alloy was observed at a mixed phase of (1 × 3) reconstruction from Reflection high-energy electron diffraction (RHEED) patterns.19 This InxGa(1-x)As layer formed at the intermediate layer acts as a protective shield for InAs QDs, which lowers the desorption energy of indium, prevents the desorption of indium adatoms from InAs QDs and maintains the original lattice point. This shield is composed of high indium concentration of 78% on an average. The capability of shielding depends on the thickness of the self-assembled InxGa(1-x)As layer and the coverage area of the InAs QDs. InAs QDs fully covered with the self-assembled InxGa(1-x)As layer reduce the leakage of indium adatoms from the InAs QDs. Fig. 2 (b) shows the HRTEM images obtained before and after GaAs layer deposition on the InAs QD matrix of the proposed VCBQD heterostructure. Clearly, after the GaAs layer deposition, the gallium adatoms interdiffused into the outer ring of the InAs QDs, forming the InxGa1-xAs capping layer. In other words, the hydrostatic compressive strain at the intermediate junction is majorly responsible for the formation of the self-assembled InxGa1-xAs layer. As we 14 ACS Paragon Plus Environment

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discussed earlier, the formation of the self-assembled InxGa1-xAs layer also involves the thermal expansion of the lattice parameters of indium and gallium adatoms because α∆T1→2 = 30°C. Using the self-assembled InxGa1-xAs capping layer that contains high indium content enhances the quantum confinement effect and quantum efficiency. We observed that the InxGa1-xAs capping layer is circular in shape, which follows the circular pattern of the outer shape of the InAs QDs. The InxGa1-xAs capping layer covers the entire InAs QDs, as shown in Fig. 2 (b). The self-assembled InxGa1-xAs capping layer can minimize the indium desorption energy and maintain the QD size and shape. Furthermore, the intermediate layer thickness is directly proportional to the hydrostatic compressive strain. The variation in the composition of indium in the self-assembled InxGa1-xAs capping layer will fluctuate the protective shielding power. A thick self-assembled InxGa1-xAs protective shield can overcome the problems of leakage of indium from the InAs QDs. Figs. 3 (a)–(f) show the cross-sectional TEM images of the as-grown VCBQD samples A and B as well as these samples annealed at 700°C and 800°C. The inset figures show the HRTEM images at a resolution of 5 nm, indicating the formation of the self-assembled InxGa1-xAs capping layer. All the HRTEM images were recorded in bright-field conditions with optimized α and β tilting angles of the sample holder considering the highest transmission region to provide a strong contrast region between InAs QDs and the self-assembled InxGa(1-x)As layer. No threading dislocation, either V-shaped or lateral QD overlap, was observed in the VCBQD samples. Fig. 3 (a) clearly shows that all the QDs were coupled vertically, but the dot size distribution is not uniform in the case of as-grown sample A. With the increase in the seed layer QD monolayer coverage from 2.5 ML to 3.2 ML, as-grown sample B overcame the problems of non-uniformity of dot distribution as the vertical strain channel energy increased. In sample B, the InAs dots 15 ACS Paragon Plus Environment


formed with high aspect ratio and high uniformity, yielding a height of 6 nm. The strain energy values mediated to couple two QDs for sample A and B calculated were 6.6, 10 meV/atom. In addition, as-grown sample B had the maximum strain mediated vertical ordering of 100% between the seed layer and active layer QDs as compared to 75 % for sample A. This is also supported by the fact that the PL spectra have the same linewidth of 33 meV for the ground state (GS) peak and the first excited peak for Sample B. The average size of QDs for sample A and B were 15 nm and 20 nm with an error of ± 2.5 nm and ± 1.0 nm respectively. The average heights of the QDs in as-grown samples A and B were 3 and 6 nm, respectively. The corresponding inset figures of as-grown sample A shows that a thin self-assembled InxGa(1-x)As layer were formed and that the capping layer is misoriented toward the [011] direction, appearing on the right side edge above the QDs. These indicate that the InxGa(1-x)As layer partially capped the QDs from the side facet of the InAs QDs. In the case of as-grown sample B, as shown in the inset of Fig. 3 (b), the thickness of a uniform capping layer at the second intermediate junction is much larger than that of the thickness of seed layer InAs QDs. The self-assembled InxGa(1-x)As layer capped perfectly and oriented toward the [001] direction. This indicates that there was no misalignment between the coupled QDs. The thickness of the InxGa(1-x)As layer can be explained by the value of hydrostatic compressive strain; larger the hydrostatic compressive strain, larger the thickness of the self-assembled InxGa(1-x)As layer. The thicknesses of the self-assembled InxGa(1-x)As layers forming the 1st intermediate junction were 1.0, and 1.5 nm and those at the 2nd intermediate junction were 2.6 and 3.7 nm for as-grown samples A and B, respectively with an error of ± 0.5 nm. To check the thermal stability of the InxGa(1-x)As layer along with InAs QDs, we annealed the VCBQD samples A and B at 700°C and 800°C. In sample A_700°C, we observed that the self-assembled InxGa(1-x)As layer disappeared because of the following reasons: 16 ACS Paragon Plus Environment

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first, annealing at a high temperature improved the crystallinity, reducing the strain at each intermediate layer, and, second, the segregation rate of indium atoms increased by thermalassisted indium out-diffusion, which resulted in the diffusion of more gallium atoms into the InAs QDs and, thereby, reduction in the QD size. In this context, the surface mobility of the gallium atoms plays a major role in smoothing this strain-free non-planar GaAs surface. Gallium migration induced by the strain field at the intermediate junction in association with the partially strain-relaxed InAs islands introduces a surface roughness in the InAs islands covered by a thin GaAs capping layer.20 In sample A (700°C), we observed non-coherent QDs with high indium agglomeration in the InAs QDs layer. In this sample, indium desorption energy increased, the strain energy decreased, and the neighboring QDs started overlapping, creating defects and dislocation. The desorption rate also increased because of the high annealing temperature, which reduced the size of the QDs. Further annealing of sample A at 800°C led to the disappearance of all the QDs from both the seed and active layers; consequently, a fast transition occurred from QDs to the InAs wetting layer (2D), as shown in Fig. 3 (e). The geometrical shape and size of the QDs changed upon annealing in both samples. These indicate that the strain energy were lowered the condition required for the formation of 3D QDs from 2D. Reduction in the carrier localization energy stimulates indium evaporation from the InAs QD structures at a high temperature.21 Annealing sample B to a high temperature of 700°C increased the height of QDs in the active layer, and the seed layer QDs greatly decreased in size simultaneously (see Fig. 3 (d)). The thickness of the self-assembled InxGa(1-x)As layer increased due to the effect of 700°C annealing temperature, and the corresponding increase in the indium composition of the capping layer resulted in a strong coupling of vertical QDs. The change in indium composition can be 17 ACS Paragon Plus Environment


examined by the contrast pattern of the HRTEM images. In sample B_700°C, the intense darkshaded circular area appeared above the InAs QDs, indicating a high indium content with a thicker InxGa(1-x)As layer. Furthermore, for sample B_700°C (Fig. 3 (d)), a thin InxGa(1-x)As capping layer of ~60 Å was formed in the 1st intermediate junction, whereas a larger thickness of ~2.5 nm was formed in the 2nd intermediate junction. A more relaxed InxGa(1-x)As layer was formed in the 2nd intermediate junction because of the short distance migration of indium and gallium adatoms induced upon annealing. As observed from the TEM images, because of the diffusion of indium into larger QDs and the consequent decrease in the density, sample B_700°C had larger QD size distribution than sample A_700°C. The lateral spacing between the two QDs increased because of the effect of the 700°C annealing temperature in sample B. The properties of QDs deteriorated in both samples A and B upon annealing at a high temperature of 800°C; this degraded the optical properties. Self-assembled InxGa(1-x)As capping layers did not develop above the QDs because the strain energy at the intermediate layers decreased for sample A_800°C and sample B_800°C. However, sample B_800°C had smaller QDs in each dot layer with low dot densities; thus, small QDs of height 2 nm were nonuniformly distributed, as shown in Fig. 3 (f). The coupling potential between two dots decreased upon annealing at a high temperature. In conclusion, the InAs QD size of sample B had the highest thermal stability; in other words, the self-assembled InxGa(1-x)As layer was not affected, and the geometrical shape and size of the QDs were maintained till the limiting temperature of 700°C for sample B because of the occupancy of a thicker InxGa(1-x)As capping layer. Figs. 4 (a)–(c) present the out-plane HRXRD results for the as-grown samples as well as the samples annealed at 700°C and 800°C. The out-plane 2θ/ω HRXRD measurements at (004) reflections were performed to obtained detailed information of In/Ga interdiffusion and 18 ACS Paragon Plus Environment

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intermixing in the VCBQD heterostructures. The zeroth-order satellite peak appears adjacent to the Bragg’s peak in the diffraction patterns of the GaAs substrate. The following are the main reasons for this: (1) presence of highly compressed (shrink) or tensile (expand) interplanar dspacing, (2) tilt angle in the 001 plane of the strain GaAs/InGaAs/InAs layer, and (3) misfit dislocation introduced while forming 3D confined QDs. The stacked QD system is considered to be under a compressive strain when the zeroth-order peak lies on the left side of the main Bragg’s peak of the GaAs substrate, i.e., θs > θ B ; in contrast, for a tensile strain, θs < θ B . All the as-grown VCBQD samples and the annealed samples show a compressive strain. On the right side of the 2θ/ω scan, fewer satellite peaks appear because of the destructive interference effects from the envelope function, depending on each individual layer thickness.22 The indium content can be expressed as ∆θ ∝ Χ .

(5).

From Figs. 4 (a)–(c), the zeroth-order satellite peak appearing on the left side of the Bragg’s peaks indicates that the system was under a compressive strain (since θ ∝

1 1 ∝ ). d a0

The indium content (X) in the entire VCBQD heterostructure was determined on the basis of two factors—(1) the separation angle (∆θ° = (θB-0th)° ) between the zeroth-order satellite peak (0th) and Bragg’s peak of GaAs (θB); (2) the amplitude of the oscillation of satellite peaks due to the difference between the densities of InAs QDs and GaAs substrates; higher the concentration of indium, greater the amplitude of the oscillation. The zeroth-order peak positions were at 65.587°, 65.4719°, 65.6°, 65.58°, 65.6°, and 65.592° for samples A_asg, B_asg, A_700°C, B_700°C, A_800°C, and B_800°C, respectively. These values revealed that sample B_asg is highly compressive with a high degree of periodicity. Therefore, compressive strain along the z19 ACS Paragon Plus Environment


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direction is determined by the shift in the zeroth-order peak position toward the GaAs Bragg’s peak. In Fig. 4 (a), sample B shows a large fluctuation in perpendicular d-spacing along the 001 plane owing to the strain accumulation in the stacked QD system. To correlate crystallinity, in terms of the broadening width β, with the ith-order satellite peaks, consider θi (i = 0, 1, 2,…) as the number of satellite peaks. Then, we have ∝

∑θ

i

∝β .

(6)

i =1

The number of oscillations of the diffracted satellite peaks will determine the sample crystal quality. The higher the rate of periodicity or number of oscillations of the diffraction peak, greater the broadening (β) due to the significant fluctuation of the d-spacing. The broadening widths (β) were 52.848, 32.868, 46.18, 38.2, 39.0, and 34.96 arcsec for samples A_asg, B_asg, A_700°C, B_700°C, A_800°C, and B_800°C, respectively. Sample C_800°C did not exhibit in any satellite peak in the second oscillation, as shown in Fig. 4 (c), because of the effect of high annealing temperature. The reduction in the amplitude of the satellite peak oscillation implies a decrease in the QD aspect ratio. Annealing at a high temperature lowered the interface roughness, as shown in Fig. 4 (b). The d-spacing values of the zeroth-order peaks were 1.42211, 1.4244, 1.4218, 1.4222, 1.42092, and 1.4219 Å for samples A_asg, B_asg, A_700°C, B_700°C, A_800°C, and B_800°C, respectively. Sample B_asg achieved a larger perpendicular lattice constant (a⏊) relative to the lattice constant of the InAs QDs. We affirmed that the VCBQD samples annealed at 800°C acquired a uniform distribution of d-spacing, indicating that sample A_800°C moved toward the Bragg’s peak position of GaAs substrate. Fig. 4 (d) shows the indium content profile for the as-grown samples and the samples annealed at 700°C and 800°C. We calculated the indium content using Eq. (6), and we found that 20 ACS Paragon Plus Environment

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sample B has the highest indium content, indicating the QDs are large and uniformly distributed. With the increase in the annealing temperature, the thermal atom migration increases, leading to the diffusion of gallium adatoms to the indium atom vacancies, and thereby increasing the GaAs thickness. Therefore, because the indium content decreases due to thermal effect, the in-flow of gallium adatoms to the InAs QDs can be controlled by the annealing temperature. The fast dissolution of the SL peaks under the compressive strain is due to the strain-induced In/Ga interdiffusion. The In/Ga intermixing across the interface between the QDs and the GaAs spacer layer changes the QD size and the composition of the QDs upon annealing. In/Ga intermixing reduces the QD size and cause a blueshift of the emission spectra. Lower intensities of the higher-order SL peaks reveals that the system has degraded the crystalline quality upon the annealing, thereby specifying a strong evidence of thermally assisted In/Ga interdiffusion at the interface in the heterostructure system.23 Figs. 5 (a) and (b) show the in-plane 2θχ/ϕ HRXRD measurement for samples A and B and the samples annealed at 700°C and 800°C to study the indium concentration in the intermediate selfassembly InxGa(1-x)As layer. In the in-plane measurement, the scattering vector lies parallel to the epilayer QD surface. The magnitude of the scattering vector depends on the in-plane d-spacing driven by non-uniform QD distribution. The low grazing incidence angle of 2θχ/ϕ were performed by considering a high intensity ϕ value to detect the In/Ga interdiffusion and interface roughness. We analyzed the diffraction peaks of the InAs QDs, GaAs, and self-assembly InxGa(1x)As

layer at (004) reflection. In Fig. 5 (a), as-grown sample A has a diffraction peak of self-

assembly InxGa(1-x)As layer at 62.18°. No peak was observed for InxGa(1-x)As layer when sample A were annealed at 700°C and 800°C. A shoulder peak at 63° was visible at the tail of the GaAs



diffraction peak, indicating the interface of InAs and GaAs. For sample A annealed at 800°C, a diffraction peak of InAs appeared at ~59° because of a thicker InAs wetting layer. We also observed that the XRD intensity of the GaAs layer was enhanced upon annealing because of increased Ga diffusion into indium (III) atom vacancies, and the InAs peak intensity decreased because of the high rate of indium desorption. The significant increase in compressive strain observed at the 2nd intermediate junction was mainly due to the incorporation of floating indium atoms.24 The supply of gallium adatoms is larger during the growth of the GaAs capping layer, which strongly modifies the phase equilibrium transition between the InAs QDs and GaAs. Thus, indium incorporated with gallium indicates a mixed phase in the form of the InxGa(1-x)As alloy.24 The self-assembly InxGa(1-x)As layer peak was observed at 62.18° at a lower angle than that of Bragg’s GaAs peak for as-grown sample A. The strain released from the 2nd intermediate junction is directly related to the concentration of indium atoms involved in the formation of the InxGa(1-x)As layer. Higher the strain, higher the indium content in the capping layer. The dspacing of the self-assembly InxGa(1-x)As layer can be tuned in the range 1.413–1.546 Å, depending on the strain energy released from the 2nd intermediate junction. High-intensity diffraction peaks of the InxGa(1-x)As layer appeared at 62.14° and 61.95° for samples B_asg and B_700°C, respectively. The self-assembly InxGa(1-x)As layer structure is not affected till 700°C, which signifies that the strain energy emitted from the 2nd intermediate junction is very high. Beyond 700°C, the self-assembly InxGa(1-x)As layer peak was not observed because of the high rate of indium desorption energy. One of the significant features involved in the self-assembly of the InxGa(1-x)As layer is the peak shift toward the InAs QDs showing high indium concentration of 82.3%. The indium concentration increased to 4% after annealing sample B at 700°C. 22 ACS Paragon Plus Environment

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Annealing sample B at 800°C increased the magnitude of the scattering vectors of GaAs and InAs QDs due to high crystalline quality. This reversibility offers a new approach toward the study of the evolution of the strain-induced islanding process and provides a new method to manipulate the 3D island sizes. According to the Heyn et al., two possible effects of indium desorption are (1) the total amount of material stored in the QDs will be reduced and (2) the QDs will become more gallium rich. Both these effects crucially modify the electronic properties of the QDs.25 The indium desorption is exponentially dependent on the growth temperature and the uniformity of the capping layer. Beyond 700°C, the indium desorption decreases the homogeneity of the dot size. A sufficiently thick strain-relaxed InGaAs layer (~ approximately 3.0 nm) reduces indium desorption, thus enhancing the indium sticking coefficient and the post-growth stability of InAs QDs Fig. 6 (a)–(d) presents the results of the in-plane RSM at (004) reflection for samples A and B as well as these samples annealed at 700°C to study the coupling of the InxGa(1-x)As capping layer with the InAs QDs. We changed the axis to the in-plane d-spacing d∥(Å) from the scattering vector of qx, qy to study strain relaxation and misorientation of the self-assembly structures (InxGa(1-x)As layer) formed at the interface. From the basic Bragg’s law, we have

d || =

2π 2

qx + q y

2

=

nλ . 4π sin(θ χ )

(7)

In Fig. 6 (a)–(d), we marked the scattering peak position along with the values for the InAs QDs, GaAs layer, and self-assembled InGaAs layer. In sample A (Fig.6 (a)), the contour lines for the InxGa(1-x)As layer that connect the GaAs and InAs were found to be misoriented on the negative plane of ϕ of the order of −1°. The RSM of 23 ACS Paragon Plus Environment


the GaAs layer shows that the diffuse contour lines are not relaxed, but elongated in the lateral and the vertical {111} direction, as shown in Fig. 6 (a). The d-spacing values of the InAs QDs, InxGa(1-x)As layer, and GaAs layer were 1.563, 1.491, and 1.412 Å for as-grown sample A. We observed from the RSM mapping that for sample A_700°C, strain-assisted InxGa(1-x)As layer was not formed (shown in Fig 6.(c)). At a high temperature, the diffused scattering of GaAs elongated toward InAs QDs, indicating a high rate of thermal-assisted Ga/In intermixing. This also signifies that the gallium atoms diffused more into the InAs QDs. We also found that the dissolution of QDs not only is due to high temperature but also depends upon the thickness of self-assembled InGaAs layer formed by strain assisted In/Ga interdiffusion, as shown in Fig. 6 (c). In Fig. 6 (b), the contour plot of RSM for sample B shows the connectivity between the InAs QDs and GaAs substrate. The d-spacing values of the InAs QDs, InxGa(1-x)As layer, and GaAs layer were 1.564, 1.493, and 1.411 Å for as-grown sample B. We found that all the diffused scattering vectors of the InAs QDs, InGaAs layer, and GaAs layer were perfectly aligned at a ϕ value of 0°. The self-assembly formed at the intermediate layer connected to InAs QDs, which indicates complete capping to the InAs QDs and strong coupling between the GaAs capping layer and the InxGa(1-x)As layer. This intermediate layer helped in preventing indium migration from the InAs QDs. This high coupling indicates the existence of a high strain in the system. In Sample A_700°C, the coupling between the InAs QDs and GaAs layer was less as shown in Fig. 6 (c). The contour maps of the GaAs layer showed elongation in all the directions, indicating a huge migration of gallium atoms toward indium (III) vacancies. This decreases the size of the QDs and makes them less confined. The d-spacing values for the InAs QDs and GaAs layer were 1.563 and 1.412 Å for sample A_700°C. In contrast to sample A_700°C, sample B_700°C still 24 ACS Paragon Plus Environment

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had the InGaAs layer at a d-spacing of 1.497 Å as shown in Fig. 6 (d). The coupling of the InxGa(1-x)As layer with the InAs and GaAs layers was slightly misoriented toward the positive ϕ plane. Due to the effect of high temperature, sample B_700°C had a more relaxed lattice plane, and the contour diffused diffraction peak had elongated in the vertical direction, indicating lower strain in sample A_700°C. 4.1. Out-plane strain calculation The average compressive strain and tensile strain were calculated26 using the following formulas: Compressive strain:

ε =

sin(θs ) −1 , sin(θs + ∆θ )

(8)

Tensile strain:

ε⊥ = −

2C12 ε || , C11

(9)

where θs is the Bragg’s angle of the GaAs substrate peak, ∆θ is the difference between the substrate and zeroth-order peaks, and C11 and C12 are the elastic constants. The vertical strain channel affects the deformation of the InAs QDs in the active layer. The out-plane compressive strain calculated for sample A and B were 1.3% and 1.6% respectively. The tensile strain calculated from the Eq. (9) were 1.4% and 1.7% for sample A and B respectively. Sample B has the highest compressive strain. We found that with larger QD size and shape, shorter the strain channel length (Ls), thereby increased the strain magnitude. Higher strain at the intermediate junction led the formation of thicker self-assembled InGaAs layer with the possibility of fully capped InAs QDs. 4.2. In-plane strain calculation 25 ACS Paragon Plus Environment


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We determined the in-plane strain at the intermediate junction and the total strain energy of the InAs QDs with capping layers of different thicknesses. The in-plane strain at the intermediate junction (bulk GaAs on InAs QDs) along the [110] direction is given as Tensile strain: ε ||

( GaAs / InAsQDs )

Compressive strain: ε ||

= + ve =

( InAsQDs / GaAs − layer )

E strain

a|| ( InAs _ QDs ) − a|| (bulk _ GaAs ) a|| ( InAs _ QDs )

= −ve =

a|| (GaAs _ layer ) − a|| ( InAs _ bulk ) a|| (GaAs _ layer )

( C 11 ) 2 + C 11 C 12 − ( 2 C 12 ) 2 2 =ν ε || . C 11

(10)

(11)

(12)

The in-plane tensile strain is greater than the compressive strain, which is true observation. The strain energy was maximum when the in-plane strain was high, and the vertical strain channel length was small. A higher percentage of compressive strain resulted in the formation of the InGaAs layer. Sample B had maximum in-plane biaxial compressive of 7.4%, thereby forming an InGaAs layer of 1.5 nm in the 1st intermediate junction and that of 3.7 nm in the 2nd intermediate junction. Higher the strain, greater the thickness of the InGaAs capping layer and lower the indium segregation rate. The measured strain energy (meV/atoms) values calculated from the above Eq. 12 were 6.6, 10, 6.2, 6.6, 6.1, 6.3 meV/atoms for A (asg), B (asg), A (700°C), B (700°C), A (800°C), B (800°C) respectively. Figure 7 shows the correlation of indium content in the self-assembly InxGa(1-x)As layer with the thickness of the InxGa(1-x)As layer, tilt angle, and twist angle, and the QD size distribution is compared with respect to indium content. We observed that a higher indium content of 78% in the InxGa(1-x)As capping layer enhances uniform distribution of QDs. Tilt and twist angles were 26 ACS Paragon Plus Environment

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derived from the out- and in-plane rocking scan at (004) reflection plane. The tilt and twist angles were found to be minimum for thick capping layers, which implies that with an increase in the capping layer thickness, the system becomes more relaxed, achieving fewer defects and dislocation. The broadening of the capping layer also gives information on the crystallinity. Larger the broadening, higher the tilt and twist angles. Lower twist and tilt angles in the VCBQD heterostructure indicate minimum leakage of indium atoms from the InAs QDs, thus imparting resistance to deformation. Figs. 8 (a) and 8 (b) shows the color mapping HRTEM images of sample A and B, which helps distinguish the shape of the self-assembled InxGa(1-x)As capping layer or the protective shield of InAs QDs. This mapping also provides information on the vertical coupling between the QDs of the seed and the active layers. We found that in sample A, the QDs in the seed layer is misaligned vertically to the active layer QDs. In contrast to sample A, sample B had a strong coupling between the QDs of the two layers; this indicates that the thicker InGaAs protective shield layer had fully covered the active layer InAs QDs. The weak coupling of the InAs QD structures (sample A) led to the formation of the InGaAs layer on the side facets of QDs along the [011] direction, which is also affected by the agglomeration of indium atoms inside the InAs QDs. The pyramidal side of the QDs, which had a high indium content, led to the formation of the self-assembled layer because of the high strain at the interface. Therefore, indium desorption is increased exponentially. In support of structural characterization data, we have presented a low temperature (18K) PL measurement results of as-grown sample A and B as depicted in Fig. 8 (c) and 8 (d). Ground state PL peak recorded at 18K for sample A and B were observed at 1181 and 1162 nm respectively, and the corresponding first excited state emission peaks were located at 1128 nm 27 ACS Paragon Plus Environment


and 1114 nm for sample A and B respectively. In the case of sample A, multiple peaks were visible due to non-uniform QD size distribution (as proven from TEM images); and the PL emission peaks intensities were too weak as compared to sample B. This is because of poor quantum confinement of QDs. The linewidth calculated for ground state (GS) and 1st excited state (ES1) PL peaks at 18K were 47 and 44 meV respectively for Sample A, which reveals inhomogeneity in QD size and shape. On the other hand, Sample B exhibited the identical linewidth of 33 meV for both the ground state (GS) and the first excited state (ES1) PL peaks, thereby indicating a high degree of uniformity of QD shape and size distribution. The enhancement in the PL intensity of sample B is due to the high stability of QDs. InAs QDs fully covered with the self-assembled InxGa(1-x)As layer enhanced the PL intensity by 77%. It should be noted that sample A is capped with thinner self-assembled InGaAs layer due to smaller dot size, and hence lesser strain at the 1st intermediate layer as compared to Sample B. Further, we observed that sample A annealed at 700°C and 800°C exhibited a blueshift of about 91 nm from 1116 nm to 1025 nm, recorded at 18K. However, in case of sample B annealed at 700°C and 800°C, we observed a lesser blueshift of 52 nm from 1117 to 1065 nm, at 18K. We also observed the existence of 1st excited state peaks at 1070 and 1025 nm for 700°C and 800°C respectively in Sample B. There is no evident signature of 1st excited state peaks in PL spectra in the case of sample A annealed at 700°C and 800°C owing to high Indium dissolubility in the active layer. Hence it also signifies that there is a transition from 3D QDs to 2D layer as we annealed the sample A from 700°C to 800°C, which deteriorated its optical properties.


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To determine the mechanism of the PL thermal quenching effect which is usually attributed to the non-radiative relaxation. The non-radiative transition probability is strongly dependent on temperature resulting in the decrease of emission intensity. Temperature-dependent PL measurements were taken to understand the carrier-escape mechanism at various temperatures. The relation between integrated PL intensity and temperature is represented by the following equation27, 28

I (T ) =

I (0) , [1 + Ce − Ea / KbT ]

(13)

where I(0) is the integrated PL intensity at 18 K, C is the ratio of the thermal escape rate to the radiative recombination rate, Ea is the activation energy, and KbT is the Boltzmann constant (8.61733 × 10-5 eV⋅K-1). Temperature-dependent PL for annealed VC QDs structures are crucial in determining the segregation rate of (In, Ga) adatoms in the system. Saumya et al. studied the effect of postgrowth annealing on bilayer QDs structure using thin spacer layer.28 The reduction of PL intensities with an increase in temperature is mainly attributed to thermal activation of carriers out from the InAs QDs to the self-assembled InGaAs layer. Thermal activation energies for samples A and B calculated using a typical Arrhenius plot were approximately 197 and 467 meV, with an error of ±20, respectively. The activation energies for samples A_700°C and B_700°C were 169 and 426 meV, whereas the activation energies for sample A_800°C and B_800°C were 141 and 412 meV, respectively. The Ea value obtained from the InAs QDs with fully capped self-assembled InGaAs layer is almost twice times larger than the InAs QD with partially capped InGaAs layer. This result indicates that InAs QDs with 29 ACS Paragon Plus Environment


fully capped InGaAs layer has a strong quantum confinement effect. Higher activation energies in sample B also reveals that the system are inbuilt with larger QDs size and high dots uniformity. We observed a similar trend of the activation energy (with respect to asg < 700°C < 800°C) for samples A and B. The reduction of Ea with an increase of annealing temperature (concerning to 700°C < 800°C) is attributed to the reduction in quantum confinement effect which is due to In/Ga interdiffusion at the dot/GaAs barrier layer interface. Moreover, reduction of activation energy for annealed VCBQD samples also reveals that there is an enhancement of gallium adatoms occupying the vacancies of Indium adatoms of the InAs QDs. Annealing sample A at 700°C results from weak QDs coupling potential and non-uniform dispersion of smaller dots. We also observed that the quality of QDs degraded due to high In/Ga interdiffusion, and created maximum defects and dislocation in the structure. Thereby reducing the activation energy from 197 to 169 meV. As we further annealed the sample A at 800°C, there is a transition of InAs QDs to InAs Quantum well, that attributes to a high rate of indium out diffusion occurred in the system. Hence, a massive reduction of the activation energy. As regards the activation energy for as-grown sample B, it was found to be 467 meV, which shows the maximum vertical ordering of 100% with uniform QDs dispersion. On annealing the sample B at 700°C, Ea reduces to 427 meV from 467 meV due to reduction in QDs size. Note that sample B(700°C) attains 100% vertical stacking with QDs capped by strain assisted self-assembled layer. On further annealing sample B at 800°C, self-assembled InxGa(1-x)As capping layers did not develop above the QDs because of the decrease in strain energy at the intermediate junction. However, sample B_800°C had smaller QDs in each dot layer with low dot densities, thereby reducing the activation energy to 412 meV from 426 meV.


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The scenario of sample A and B annealed at 700°C is entirely different from one another, one with thick InGaAs capping layer (sample B, Ea = 427 meV) and other without InGaAs capping layer (sample A, Ea = 169 meV). We also observed a high PL intensity for VC bilayer QDs heterostructures inbuilt with the prominent features of self-assembled InxGa(1-x)As layer. Sample B showed the highly stable InGaAs layer at each intermediate layer as compared with Sample A. This is attributed to fewer defects in the system and also highly controlled interdiffusion (In/Ga) across the interface between QD and GaAs barrier. The absence of self-assembled layer in samples A_700°C and A_800°C reveals excessive dissolubility of indium adatoms into the GaAs layer, thereby resulting in the transition from 3D QDs to 2D layer. 4. Conclusion By analyzing the properties of InAs QDs with the self-assembled InxGa(1-x)As capping layer, the main factors that determine the enhancement of PL intensity can resolved. Fully capped InAs QDs with the self-assembled InxGa(1-x)As layers resulted in a low linewidth of 33 meV, while partially capped InAs QDs exhibited a broader linewidth of 48 meV. The fully capping layer induced a high indium composition of 78% and acquired activation energy of 467 meV. Strain relaxation involves higher indium compositions in the capping layer subjected to annealing at 700°C. Partially capped layers were formed at the side face of the pyramidal QDs and featured a high ratio of tilt and twist angles. The temperature-dependent PL analysis of InAs QDs fully covered with self-assembled InxGa(1-x)As layer containing a high indium content showed a strong coupling between the QDs of the seed and active layers, and the uniformity of QDs enhanced to 100% compared with those of partially capped QDs. The 100% vertical ordering of the coupled 31 ACS Paragon Plus Environment


QDs of sample B were evidence by the HRTEM images, in- and out-plane HRXRD, and PL spectra at 18°C. Annealing of the VCBQD at a high temperature (800°C) resulted in the formation of QDs without the InxGa(1-x)As capping layer, thereby changes from 0D QD to the 2D quantum well. This is due to the reduction in strain energy in the system, driven by the transition from QDs to the thicker 2D InAs wetting layer. Hence, QDs without the InxGa(1-x)As capping layer lower the degree of confinement and coupling between the QDs. High thermal stability, very compact, high degree of monochromaticity were the three most impactful contribution of our proposed VCBQDs structure to the future development of single photon emitter based semiconductor laser. Thus, we conclude that VCBQDs structure inbuilt with self-assembled InGaAs layer is best suitable for O-band fiber optics in telecom application with minimum attenuation. Our VCBQDs work provides a novel way of preparing high-purity single photons. These VCBQDs based single photon emitter can support high data rate transfer and more reliable communication. Furthermore, our design VCBQDs structure has 100% vertical ordering of which it will get a better stable monochrome output. Uniform dispersion of photon energy emitted by our proposed coupled QDs structures makes the reliable production of such devices. By increasing the iteration of QDs layer in such VCBQDs systems would be highly advantageous for 1550 nm telecommunication window, with a narrower FWHM.

Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. ‡These authors contributed equally. 32 ACS Paragon Plus Environment

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FUNDING SOURCES The authors acknowledge the funding received from the Centre of Excellence in Nanoelectronics, Department of Information and Technology, Government of India, National Center of Excellence in Technology for Internal Security, IITB and IITB Nanofabrication facility (IITBNF), Department of Electrical Engineering, IITB. The authors also acknowledge the financial support provided by the Department of Science and Technology and Department of Science and Technology Nano Mission, India, as well as the funding received from the Indian Space Research Organization. These funding organizations have no role in the study design, collection, analysis, and interpretation of data, writing of the report, and the decision to submit the article for publication. The authors declare no conflict of interest. ACKNOWLEDGMENT The authors thank the TEM facility at Sophisticated Analytical Instrument Facility, Indian Institute of Technology, Bombay (IITB) and the HRXRD facility at the Department of Physics, IITB.

ABBREVIATIONS Ls

Strain channel length

HRTEM

High-resolution Transmission Electron Microscopy

HRXRD

High-resolution X-Ray Diffraction

PL

Photoluminescence

QD

Quantum dot

RSM

Reciprocal space mapping

VCBQD

Vertical coupled bilayer quantum dot 33 ACS Paragon Plus Environment


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Figure 1. A schematic representation of the proposed vertical-coupled bilayer InAs QD (VCBQD) heterostructure inbuilt with a self-assembled InxGa1-xAs layer. The green arrows indicate the vertical inter-QD strain channel. The inset table shows the growth specification and classification of the VCBQD samples.

B. Tongbram et al. Fig. 1/8



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(a)

(b) Figure 2. (a) The schematic unit cell architecture showing each interface and defining the formation of the strain-assisted self-assembled InxGa1-xAs layer. (b) High-resolution TEM images of the InAs QDs showing the growth process of the self-assembled InxGa1-xAs layer before and after the deposition of the GaAs layer. 40


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(a)

(b)

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Figure 3. (a)–(f) Cross-sectional TEM images of the as-grown VCBQD samples A and B and their respective samples annealed at 700°C and 800°C obtained at a 20 nm resolution; the inset figures show HRTEM images at a 5 nm resolution; they indicate the formation of the selfassembled InxGa1-xAs layer at the intermediate junction between the InAs QDs and GaAs capping layer/spacer layer.



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(b) (a)

(d)

(c)

Figure 4. (a)–(c) The out-plane (2θ/ω) HRXRD measurement at (004) reflection of the as-grown VCBQD samples A and B and their respective samples annealed at 700°C and 800°C. (d) The indium profiles of the as-grown and annealed VCBQD samples.




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(b)

(a)

Figure 5. The in-plane (2θχ/ϕ) HRXRD measurement at (004) reflection for all the annealed VCBQD samples along with the as-grown samples shown in (a) Sample A, (b) Sample B. Asgrown samples show the existence of the self-assembled InxGa1-xAs layer.

B. Tongbram et al. Fig.5/8


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Figure 6. In-plane high-resolution reciprocal space mapping (RSM) at (004) reflection for the as-grown VCBQD samples (a) Sample A, (b) Sample B, and the samples annealed at 700°C shown in (c) Sample A_700°C, (d) Sample B_700°C.

B. Tongbram et al. Fig.6/8


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Figure 7. Indium profile of InxGa(1-x)As capped layer with correlation to tilt and twist angles of the VCBQD samples, size distribution of QDs (FWHM), and with the capping thickness of the 2nd intermediate InxGa(1-x)As layer.




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(a)

(b)

(c)

(d)

Figure 8. A color contrast mapping of XTEM images of as-grown (a) sample A, and (b) sample B to explore the characteristics of fully and partially capped QDs. A low temperature photoluminescence spectra at a power density of 1.1 kW/cm2 for as-grown (c) sample A, and (d) sample B; The insets show the Gaussian fitting lineshapes of each PL spectrum recorded for samples A and B, yielding a full width at half maximum (FWHM).

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