Multi-stacked InAs/GaAs quantum dots grown with different growth modes for quantum dot solar cells

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1 Multi-stacked InAs/GaAs quantum dots grown with different growth modes for quantum dot solar cells Yeongho Kim a), Keun-Yong Ban b), and Christiana B. Honsberg School of Electrical, Computer and Energy Engineering, Arizona State University, Tempe, Arizona 85287, USA We have studied the material properties and device performance of InAs/GaAs quantum dot solar cells (QDSCs) made using three different QD growth modes: Stranski Krastanov (S K), quasi-monolayer (QML), and submonolayer (SML) growth modes. All QDSCs show an extended external quantum efficiency (EQE) at near infrared wavelengths of nm from the QD absorption. Compared to the S K and SML QDSCs, the QML QDSC with a higher strain exhibits a poor EQE response in the wavelength region of nm due to increased non-radiative recombination. The conversion efficiency of the S K and SML QDSCs exceeds that of the reference cell (13.4 %) without QDs due to an enhanced photocurrent (> 16 % increase) produced by the silicon doped QD stacks. However, as expected from the EQE of the QML QDSC, the increase of strain-induced crystalline defects greatly degrades the photocurrent and open-circuit voltage, leading to the lowest conversion efficiency (8.9 %). Self-assembled quantum dots (QDs) have attracted considerable interest as a promising active medium for quantum devices such as lasers, optical switches, and photovoltaic devices due to a high quantum efficiency and wide wavelength tunability. 1,2,3 Recently, quantum dot solar cells (QDSCs) using InAs/GaAs QDs have been studied extensively in an effort to exceed the detailed balance efficiency limit of single-junction solar cells. 3,4,5 They have been demonstrated to have improved short-circuit current density (J sc ) through sub-bandgap photon absorption whereas the open-circuit voltage (V oc ) is often degraded by increased non-radiative recombination mainly due to the creation of crystalline defects. Consequently, the photovoltaic conversion efficiencies of QDSCs have been much lower than that of the reference GaAs solar cells. 3,5 In order to grow self-assembled InAs/GaAs QDs as a sub-bandgap photon absorber for QDSCs, any of three different growth modes can be employed. In the widely employed Stranski Krastanov (S K) growth mode, three-dimensional (3D) QDs with a thin wetting layer are formed by depositing InAs with a nominal coverage (> ~1.75 ML) on GaAs through a strain relief process. 6 Alternatively, self-assembled QDs can be a) Electronic mail: ykim172@asu.edu. b) Permanent address: Applied Materials, Inc., Santa Clara, CA , USA 1

2 grown by cyclic deposition of InAs (~1 ML) partially capped with a very thin GaAs layer (~1 ML) in the quasimonolayer (QML) growth mode. 7 Finally, the multi-stacking of ML thick InAs fully capped with a few monolayers (2 4 ML) of GaAs forms quantum-dot-quantum-well-like islands slightly elongated towards the [ 110] direction, and is called the sub-monolayer (SML) growth mode. 8,9 Unlike the S K and QML QDs the SML QDs have no InAs wetting layers because of the highly reduced deposition amount of InAs. The potential fluctuation of the wetting layers with two-dimensional (2D)- like density of states influences the carrier trapping and escape. 10 Moreover, the SML QDs have been reported to have a high areal density (~10 11 cm -2 ) and uniform size distribution of QDs which are advantageous to increasing the sub-bandgap photon absorption. 9 So far, there have been no systematic reports investigating the effect of different QD growth modes on the device performance of QDSCs. In this letter, the S K, QML, and SML QD growth modes have been utilized to realize multi-stacked InAs/GaAs QDSCs. The S K and SML QDSCs show improved conversion efficiencies due to the high J sc induced by the enhanced material properties and the absorption of sub-bandgap photons by the QDs. The high strain and resulting defects in the QML QDSC cause strong non-radiative recombination of the photo-generated carriers, leading to a low J sc and V oc and thus degraded conversion efficiency. All samples were grown by a solid-source molecular beam epitaxy system on p + -GaAs (001) substrates. The reference GaAs solar cell without the QDs had a p-i-n structure with a 50 nm n + -GaAs contact layer, a 50 nm n + -Al 0.8 Ga 0.2 As window layer, a 150 nm n + -GaAs emitter layer, and an 1 m undoped GaAs base layer. For the QDSCs, three different QD stacks consisting of InAs (2.0 ML)/GaAs (10 nm), InAs (1.2 ML)/GaAs (1.2 ML), and InAs (0.5 ML)/GaAs (2.5 ML) were grown at the center of the base layer of the reference cell at 500, which correspond to the S K, QML, and SML QDSCs, respectively. To maintain the same level of the total InAs coverage of ~20 ML in the stacks, the numbers of QD stacks were set to 10, 20, and 40 for the S K, QML, and SML QDSCs, respectively. The QD layers were delta (δ)-doped with silicon (Si) at an areal density of cm -2 to supply electrons into the QDs. During the InAs deposition, the in-situ reflection high-energy electron diffraction showed chevron patterns, indicating a 3D growth, for the S K and QML QDSCs while it displayed sharp streaky patterns, the sign of a 2D growth, for the SML QDSC. The double-crystal (DC) ω-2θ and triple-crystal (TC) ω rocking curves (RCs) were measured to evaluate the elastic strain and crystal quality using high-resolution x-ray diffraction (XRD) with CuK α1 radiation 2

3 ( = Å). The photoluminescence (PL) measurement was performed at 10 K with a 532 nm laser and liquid-nitrogen cooled Ge detector to study optical transitions in the QDs. The atomic force microscopy (AFM) measurement was conducted using a Veeco Dimension 3100 AFM under a tapping mode to characterize the QD size and density of the uncapped S K and QML QDs grown under the same growth conditions for the QD stacks of the S K and QML QDSCs. For device characterization, Ti/Pt/Au and Ge/Au/Ni/Au were deposited on the backside and frontside of the as-grown samples and annealed in nitrogen for p-type and n-type ohmic contacts, respectively. An antireflection coating was not applied to the samples. The external quantum efficiency (EQE) spectra were measured at room temperature using a QEX10 system with a xenon arc lamp source dispersed by a dual-grating monochromator. The current-voltage (I-V) characteristics were taken under one-sun illumination (100 mw/cm 2 ) at AM 1.5G in an Oriel Class A solar simulator. Figure 1(a) shows the DC ω-2θ RCs in the vicinity of the (004) reflection for the as-grown samples. Compared to the reference cell, the periodic superlattice (SL) peaks (denoted by the numbers 0, ±1, ±2, etc.) of the QDSCs are clearly observed, originating from the multi-stack of InAs/GaAs QD layers. The zeroth-order (0 th ) SL peaks of the QDSCs are located on the left side of the GaAs substrate peak, indicating the QD stacks are inplane compressively strained. The in-plane strain (ε ) for the QDSCs is estimated from the out-of-plane strain (ε ) measured by the peak separation between the 0 th -SL peak and the substrate peak using Poisson s ratio for the biaxially strained system. 11,12 (1 ) (1 ) sin B 1 (1) 2 2 sin( ) where is Poisson s ratio of the SL, θ B is the substrate Bragg angle, and Δθ is the angle separation between two peaks of the 0 th -SL and substrate. The values of ε are determined to be -0.74, -1.55, and % for the S K, QML, and SML QDSC, respectively. The large compressive ε of the QML QDSC causes an increase in tensile strain along the growth direction by the Poisson's effect, increasing the aspect ratio (QD height-to-base ratio) of the QDs. The sharp interference fringes between neighboring SL peaks in the ω-2θ RCs for the S K and SML QDSCs are observed whereas they are not found for the QML QDSC, which implies that the crystal quality of the QML QDSC is deteriorated because of the increased strain. To further investigate the crystallinity of the QDSCs, TC ω RCs of the 0 th -SL peak position are measured as seen in Fig. 1(b). The different XRD intensities of the coherent central peaks of the ω RCs are due to the change in the diffraction volume of the QD structures, B 3

4 together with the crystallinity of the QDSCs. The full-width-at-half-maximums (FWHMs) of the central peaks of the ω RCs are derived to be 10.4, 214.8, and 23.4 arcsec for the S K, QML, and SML QDSC, respectively. From the FWHMs, the densities of dislocation loops in the volume of the structures are calculated to be , , and cm -2 for the S K, QML, and SML QDSC, respectively. 13 The highly increased dislocation density of the QML QDSC is due to the increase of the misfit strain as confirmed by the strain analysis. Figure 2 shows the PL emission spectra at 10 K from the QDSCs with different growth modes. The PL peak of the ground state (GS) emission for the QML QDSC is positioned at ev, which is redshifted by 65 mev with respect to that (1.141 ev) for the S K QDSC. This is explained as follows. The AFM images in the insets of Fig. 2 exhibit that the average diameter (height) and areal density of the S K QDs are 28.9 (5.8 nm) and cm -2 while those of the QML QDs are 43.3 (14.6 nm) and cm -2. In case of the SML QDs, the QD-like features are not observed since the AFM used is unable to measure an extremely small height (~1 ML) due to its resolution limit. From the AFM analysis the aspect ratios of the S K and QML QDs are determined to be 0.20 and 0.34, respectively. Thus, the longer wavelength emission from the QML QDSC is attributed to the higher aspect ratio of the QDs. This is in accordance with the report that the PL peak position is redshifted with increasing aspect ratio of dots due to the large QD size, especially the QD height. 14 Meanwhile, the PL emission from the SML QDSC exhibits a single sharp QD peak located at ev, which is near the GaAs band-edge (~1.52 ev at 10 K). The GS peak shift of the SML QDSC towards the GaAs band-edge is caused by quantum size effects. Unlike the S K and QML QDs, the SML QDs are ~1 ML high 2D islands with a lateral dimension of 5 10 nm, which is much smaller than the exciton Bohr radius (~34 nm) of bulk InAs. 15 The excitonic transition energy of such smaller QDs is increased because of the increased energy level separation between conduction and valence band states with a significant decrease of QD sizes. 16 Moreover, the PL spectrum of the SML QDSC has the narrowest FWHM of ~6 mev, which is nearly eight times less than that of the S K and QML QDSCs. F. Liu et al. have reported that the surface strain accelerates the 2D island-island interactions via island edge diffusion, resulting in a narrow size distribution of dots. 17 Thus, the strain-driven island migration is responsible for the very sharp PL linewidth of the SML QDSC. In addition, it is noted that the PL intensity of the S K QDSC is stronger due to the higher dot density together with improved crystal quality as compared to the QML QDSC. Further, the PL intensity of the SML QDSC is five times stronger than that of the S K QDSC. The crystallinity of the SML QDSC is degraded by the higher compressive strain (ε = - 4

5 0.78 %) in the QD stacks as confirmed by the XRD analysis. The increased misfit strain can generate nonradiative recombination centers, leading to the PL quenching. However, the strong PL intensity can be obtained by the highly increased radiative recombination of carriers in the SML QDs which may have smaller dot size and ultra-high dot density on the order of ~10 11 cm -2 under the similar growth conditions. 9 To study the effect of different QD growth modes on the EQE of the fabricated cells, the EQE spectra of the reference cell and QDSCs are measured at room temperature. As seen in Fig. 3, all of the QDSCs exhibit an extended spectral response beyond the GaAs band-edge (880 nm) due to the sub-band gap photon absorption by the QDs. The absorption edges in the near infrared region are 1070, 1050, and 950 nm for the S K, QML, and SML QDSC, respectively. It is stressed that the absorption edge of the SML QDSC is reduced due to the decreased QD size. This is analogous with the PL data shown in Fig. 2. Its EQE is higher in the wavelength region of nm than those of the S K and QML QDSC. This behavior is ascribed to a higher dot density and the absence of the wetting layers resulting in efficient carrier transport. 9,18 Furthermore, it is found that the EQE in the wavelength region of nm is greatly affected by the different QD growth modes. The inset of Fig. 3 shows the EQE ratio of the QDSCs normalized to that of the reference cell. In the wavelength region of nm, the EQE ratio of the S K and SML QDSC is greater than one. This means that the EQE of the QDSCs is slightly improved compared with the reference cell, which is probably due to the Si doping effect by which the excess strain can be mitigated, leading to the suppression of the non-radiative recombination. 19 On the other hand, a relatively significant reduction by above 10 % in the EQE of the QML QDSC is observed over the wavelength region, which is indicative of a severe degradation in the emitter and base layer caused by dislocation propagation. 20 We attribute the degraded EQE in the QML QDSC to the increased misfit and threading dislocations acting as non-radiative recombination centers as evidenced by the broadening of the FWHM of the TC ω RCs (Fig. 1(b)) and weak PL intensity (Fig. 2). The I-V characteristics of the QDSCs under one-sun illumination condition (100 mw/cm 2, AM 1.5G) are shown in Fig. 4. The open-circuit voltage (V oc ), short-circuit current density (J sc ), fill factor (FF), and conversion efficiency ( ) are listed in Table I. Compared to the reference cell, all of the QDSCs have a low V oc of V due to the introduction of InAs QDs with a narrow effective bandgap into the base region of the QDSCs. The V oc can be further reduced with an increase of non-radiative defects in the QD structures. Among the QDSCs, the QML QDSC has the lowest V oc and J sc as a result of the high non-radiative recombination of photo-generated carriers. The degradation in both V oc and J sc leads to a highly deteriorated FF 5

6 of 63.1 % for the QML QDSC. To analyze the extent of the non-radiative recombination, the experimental data of the I-V curves are fitted with the double diode model of 21 J q(v JR ) q(v JR ) V JR s s s J ph Jdiff exp 1 J nr exp 1 (2) n1k BT n 2k BT R sh where J ph is the illuminated photocurrent density, J diff is the diffusion current density, J nr is the non-radiative recombination current density through crystalline defects, R s is the series resistance, R sh is the shunt resistance, n 1(2) is the diode ideality factor (n 1 < n 2 ), and k B T is the thermal energy (25.9 mev), respectively. The J diff and J nr of the QML QDSC, as listed in Table I, are the highest among the samples due to the enhanced non-radiative recombination in the QDs. These high non-radiative recombination losses result in the poor J sc and V oc of the QML QDSC. Interestingly, the J sc of the S K and SML QDSCs is markedly improved by 16 and 20 % with respect to that of the reference cell, respectively. The increased J sc is ascribed to the sub-bandgap photon absorption by the QDs in conjunction with the EQE improvement in the wavelength region below the GaAs band-edge. However, the J sc increment of the QML QDSC is marginal owing to the deteriorated crystal quality and insufficient sub-bandgap photon absorption limited by the low QD density ( cm -2 ). Another possible cause of the J sc variation of the QDSCs is the equivalent number of electrons per dot from the Si doping. The QML QDs have a higher electron density of 3.5 e - /dot compared with that (1.7 e - /dot) of the S K QDs. We have recently found that the carrier lifetime of InAs QDs, which were highly δ-doped with Si, was reduced due to the strong carrier-carrier scattering. 22 Therefore, we think that the J sc of the S K QDSC can be further increased since the moderate electron density is beneficial to improve the carrier lifetime as well as to inject electrons to the dots for the efficient QD absorption. The highly improved J sc of the S K and SML QDSC makes the higher by 2.4 and 0.7 % than that of the reference cell by compensating for the reduced V oc, respectively. Thus, it is demonstrated that the different QD growth modes affect the device performance of the QDSCs and the can be boosted in a way that a highly improved J sc offsets a degraded V oc. To conclude, the effects of the different QD growth modes on the material and device characteristics of InAs/GaAs QDSCs have been investigated. The XRD results reveal that the highly accumulated strain in the QML QDSC facilitates the formation of crystalline defects in the structure. With respect to the PL peak position of the S K QDSC, the QML and SML QDSCs show the opposite PL shifts (redshift and blueshift, respectively) caused by increased tensile strain and much reduced dot size. It is also found that the EQE of the S K and SML 6

7 QDSCs is enhanced due to the improved crystal quality and high sub-bandgap photon absorption. However, the EQE of the QML QDSC in the wavelength region below the GaAs band-edge drops because of the increased crystalline defects. This crystal imperfection leads to the low (8.9 %) of the QML QDSC. In contrast, the high crystallinity of the S K and SML QDSCs results in improved which is higher than that (13.4 %) of the reference cell. Therefore, we demonstrate that the strain accumulation in the stacked QD structures highly affects the material properties and device performance and it can be controlled by using different QD growth modes. This material is based upon work primarily supported by the National Science Foundation (NSF) and the Department of Energy (DOE) under NSF CA No. EEC Any opinions, findings and conclusions or recommendations expressed in this material are those of the author(s) and do not necessarily reflect those of NSF or DOE. We gratefully acknowledge the use of facilities within the LeRoy Eyring Center for Solid State Science at Arizona State University. We also acknowledge Dr. Yong-Hang Zhang for providing the photoluminescence system and solar cell probe station. 7

8 FIG. 1. (a) Experimental DC ω-2θ RCs measured around the symmetric (004) reflection for the as-grown samples. The arrows for the QDSCs indicate the 0 th -SL peaks. (b) Experimental TC ω RCs scanned at the 0 th -SL peaks of the QDSCs. 8

9 FIG. 2. PL spectra of the S K, QML, and SML QDSCs measured at 10 K. The PL spectrum of the S K and QML QDSC is decomposed into Gaussian bands consisting of a ground state emission (indicated by an arrow) and three excited state ones. The insets show the AFM images (1 1 m 2 ) of the uncapped S K and QML QDs. 9

10 FIG. 3. EQE spectra of the reference solar cell and QDSCs. The inset shows the EQE ratio between the reference cell and QDSCs. 10

11 FIG. 4. Illuminated I-V characteristics for the reference solar cell and QDSCs under one-sun illumination (100 mw/cm 2 ) at AM 1.5G. The experimental data are plotted as open circles and the fitting results are plotted by solid lines. 11

12 TABLE I. Photovoltaic device parameters and diffusion/non-radiative recombination current density extracted from the double diode model for the fabricated solar cells. Sample V oc (V) J sc (ma/cm 2 ) FF (%) (%) J diff (ma/cm 2 ) J nr (ma/cm 2 ) GaAs Ref S K QDSC QML QDSC SML QDSC

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