Perovskite-Based Artificial Multiple Quantum Wells - Nano Letters

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Perovskite-Based Artificial Multiple Quantum Wells Kwang Jae Lee, Bekir Turedi, Lutfan Sinatra, Ayan A. Zhumekenov, Partha Maity, Ibrahim Dursun, Rounak Naphade, Noor Merdad, Abdullah Alsalloum, Semi Oh, Nimer Wehbe, Mohamed Nejib Hedhili, Chun Hong Kang, Ram Chandra Subedi, Namchul Cho, Jin Soo Kim, Boon S. Ooi, Omar F. Mohammed, and Osman M. Bakr Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.9b00384 • Publication Date (Web): 22 Apr 2019 Downloaded from http://pubs.acs.org on April 22, 2019

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Perovskite-Based Artificial Multiple Quantum Wells Kwang Jae Lee,†,¶ Bekir Turedi,†,¶ Lutfan Sinatra,⊥ Ayan A. Zhumekenov,†,¶ Partha Maity,† Ibrahim Dursun,†,¶ Rounak Naphade,†,¶ Noor Merdad,†,¶ Abdullah Alsalloum,†,¶ Semi Oh,§ Nimer Wehbe,‡ Mohamed N. Hedhili,‡ Chun Hong Kang,Δ Ram Chandra Subedi,Δ Namchul Cho,∥ Jin Soo Kim,○ Boon S. Ooi,Δ Omar F. Mohammed,† & Osman M. Bakr*,†,¶ †Division

of Physical Sciences and Engineering, ¶KAUST Catalyst Center (KCC), ΔPhotonics

Lab, ‡Core Lab, King Abdullah University of Science and Technology (KAUST), Thuwal 23955-6900, Kingdom of Saudi Arabia ⊥Quantum

∥Department

Solutions LLC, Thuwal 23955-6900, Kingdom of Saudi Arabia

of energy systems engineering, Soonchunhyang University, Asan 31538, Republic of Korea

§School

of Materials Science and Engineering, Gwangju Institute of Science and Technology, Gwangju 61005, Republic of Korea

○Division

of Advanced Materials Engineering and Research Center of Advanced Materials

Development, Chonbuk National University, Jeonju 54896, Republic of Korea

Keywords: Perovskite, quantum well, bandgap engineering, CsPbBr3, hot carrier, femtosecond spectroscopy


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Abstract

Semiconductor quantum well structures have been critical to the development of modern photonics and solid-state optoelectronics. Quantum level tunable structures have introduced new transformative device applications and afforded myriad groundbreaking studies of fundamental quantum phenomena. However, non-colloidal, III-V compound quantum well structures are limited to traditional semiconductor materials fabricated by stringent epitaxial growth processes. This report introduces artificial multiple quantum wells (MQWs) built from CsPbBr3 perovskite materials using commonly available thermal evaporator systems. These perovskite-based MQWs are spatially aligned on a large-area substrate with multiple stacking and systematic control over well/barrier thicknesses, resulting in tunable optical properties and a carrier confinement effect. The fabricated CsPbBr3 artificial MQWs can be designed to display a variety of photoluminescence (PL) characteristics, such as a PL peak shift commensurate with the well/barrier thickness, multi-wavelength emissions from asymmetric quantum wells, the quantum tunneling effect, and long-lived hot-carrier states. These new artificial MQWs pave the way toward widely available semiconductor heterostructures for light-conversion applications that are not restricted by periodicity or a narrow set of dimensions.


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Introduction A multiple quantum well (MQW) is a repetitive semiconductor heterostructure in which one thin active layer (< 10 nm; comparable to the de Broglie wavelength of the charge carriers) is bound by potential barriers.1 MQWs are marked by discrete energy levels with a step-like density of states (DOS), spatial overlap between the electron-hole wave functions, and strong carrier quantumconfinement effects. These structures have technologically revolutionized traditional lightemitting diodes (LEDs),2 laser diodes (LDs),3 optical modulators,4 and switching sensors.5 They have also facilitated conceptually new iterations of device applications, including quantum cascade lasers,6 super-luminescent diodes,7 and structures that sustain terahertz charge oscillations.8 Additionally, MQWs have led to groundbreaking studies of peculiar quantum phenomena, such as the quantum-confined Stark effect (QCSE), the optical Stark effect, and Bose-Einstein condensation of exciton polaritons.9-12 Despite their triumph in many research areas of significance, existing MQWs have been mostly restricted to conventional semiconductor materials, such as InxAlyGa1-x-yAs/GaAs, InxGa1xN/GaN,

and GaN/AlxGa1-xN,13-17 which demand time-consuming and capital-intensive processes

such as metal-organic chemical vapor deposition (MOCVD) and molecular-beam epitaxy (MBE). Specifically, these methods require high growth temperatures (>800 °C) and substrates that match the crystallographic orientation of the epitaxial thin films, as well as high single-crystal thin-film quality for low densities of deep-level traps.13, 16, 17 These requirements have inhibited the scope of MQWs to niche and high-end applications and prevented them from becoming a ubiquitous part of energy, lighting, and sensing applications.


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Halide perovskites are defect-tolerant semiconductors18,

19

that can be fabricated into

numerous material states ranging from bulk single crystals20-22 to low-dimensional structures (nanowires,23, 24 quantum dots,25 nanocrystals,26-30 etc.) by either vacuum deposition or solutionbased techniques. This tremendous variety of structures, phases and dimensionalities makes perovskites an excellent candidate for a broad array of applications, such as solar cells, photodetectors, LEDs, scintillators, and amplified spontaneous emission (ASE); thus, perovskites offer immense potential for realizing MQWs that could overcome the limitations of conventionally fabricated semiconductor heterostructures. For instance, several studies have recently introduced solution-processed Ruddlesden–Popper perovskites (RPPs) as chemically synthesized natural MQWs composed of atomically precise perovskite layers separated by organic spacers.31-34 These repetitive structures of two-dimensional (2D) perovskite sheets can be structurally considered MQWs. Furthermore, heterostructures of different 2D RPPs with different well thicknesses can be formed by the bottom-up nanostructure synthesis of 2D perovskite sheets and the epitaxial growth of perovskites on suitable substrates by CVD.35, 36 Here, we demonstrate a facile approach for fabricating perovskite-based artificial MQWs using commonly available thermal evaporator systems. These MQWs are directly built from CsPbBr3 perovskite layers and 1,3,5-tris(N-phenylbenzimidazol-2-yl)benzene (TPBi) barriers. These CsPbBr3/TPBi artificial MQWs, of type-I band alignment, enable precise and arbitrary control over quantum wells/barriers and their spatially aligned multiple-stacking configuration, without requiring specific substrates as epitaxial growth normally entails.37, 38 The CsPbBr3/TPBi artificial MQWs, which are polycrystalline in nature, exhibited tunable optical characteristics related to the well/barrier thicknesses and stacking configurations, including thickness-dependent


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carrier transitions and confinement, multi-wavelength emissions, the quantum tunneling effect (in temperature-dependent photoluminescence, TDPL), and long-lived hot carriers.

Results and Discussion Fabrication of perovskite-based artificial MQWs To design perovskite-based artificial MQWs, we chose TPBi as a quantum barrier material. Quantum barriers serve to confine electrons and holes in a well. It is noted that the conduction and valence bands of quantum barrier should be offset with respect to the bands of the well material. To confirm band offsets (ΔE) between CsPbBr3 (the well) and TPBi (the barrier), we determined the core-level binding energy and valence-band maximum (VBM) of 50-nm-thick CsPbBr3 (138.28 eV and 1.20 eV, respectively) and 50-nm-thick TPBi (398.38 eV and 1.40 eV, respectively) by analyzing their X-ray photoelectron spectra (XPS) using Kraut’s method39 (Figure S1a-b). For a CsPbBr3 (3 nm)/TPBi (50 nm) double-layer structure, we extracted a 260.41 eV binding energy difference between these two reference core levels (Pb 4f7/2 and N 1s) (Figure S1c). We estimated the bandgap energies of CsPbBr3 (2.36 eV) and TPBi (3.24 eV) from their PL peak positions (Figure 1a). Lastly, the ΔE values for the conduction band and valence band were calculated to be 0.37 eV and 0.51 eV, respectively, by combining all measurement values (Figure 1b), indicating type-I band alignment between CsPbBr3 and TPBi.


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Figure 1. Structuring of perovskite-based artificial MQWs. (a) PL (325 nm excitation) spectra of TPBi and CsPbBr3 perovskite powders. (b) Bandgap alignment between CsPbBr3 and TPBi thin films. (c) Schematic diagram of MQW structuring method using thermal evaporation. (d) Elemental distribution with depth obtained using secondary-ion mass spectrometry (SIMS). The ion signals, ascribed to Br, Cs, Pb, and Si, are plotted versus the sputtering depth. (e) 3D mapping of Br and Si signals by SIMS acquired from each pixel as a function of depth by scanning the ion beam over a surface area divided into 100 × 100 pixels. (f) Cross-sectional TEM image of 5-stacked CsPbBr3 (3 nm)/TPBi (7 nm) MQWs.

We fabricated type-I MQWs with CsPbBr3 and TPBi layers by thermal evaporation. Figure 1c depicts a schematic diagram of the MQW fabrication process. First, TPBi powder was evaporated onto a glass substrate at an exquisitely optimized evaporation rate of 0.20 Å/s under a chamber pressure of 4 × 10-6 Torr. A CsPbBr3 powder, prepared by the inverse temperature crystallization (ITC) method20 (Figure S2), was evaporated onto the TPBi thin film at an evaporation rate of 0.15 Å/s under the same chamber pressure. The pair of TPBi and CsPbBr3 thin films was sequentially stacked multiple times on a glass substrate and terminated with a TPBi layer on top, resulting in the CsPbBr3/TPBi MQW structure. The phase purity of CsPbBr3 before evaporation and after depositing 20-nm-thick thin films was confirmed by X-ray diffraction (XRD) (Figure S3).


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After evaporating 5-stacked CsPbBr3/TPBi MQWs on the glass substrate, we analyzed the elemental distribution through the depth of the stacking structure using secondary-ion mass spectrometry (SIMS) (Figure 1d and S4). The ion signals, ascribed to CsPbBr3 (Br, Cs, and Pb), TPBi (CN), and glass (Si), were plotted versus the sputtering depth. Indeed, all ion signals (Br, Cs, and Pb) assigned to CsPbBr3 exhibited strong enhancement when the CsPbBr3 single layer was reached and dropped sharply once it was crossed. This observation is a good indication of the interface quality between the CsPbBr3 and TPBi layers. We also performed 3D mapping by scanning the ion beam over a surface area divided into 100 × 100 pixels as a function of depth (Figure 1e). Collected Br and Si signals from each pixel demonstrated the presence of 5-stacked CsPbBr3/TPBi MQW, confirming successful thin-film growth. Cross-sectional transmission electron microscopy (TEM) images also verified the 5-stacked structure of the MQW, in line with the SIMS results (Figure 1f). We confirmed the thickness of the quantum wells and quantum barriers to be 3 nm and 7 nm, respectively.

Optical properties of perovskite-based artificial MQWs Perovskite-based artificial MQWs showed variable optical characteristics through bandgap engineering. To confirm the quantum-confined carrier transition, we first fabricated a series of CsPbBr3/TPBi single QWs (SQWs) with various CsPbBr3 quantum well thicknesses (LQW). With the quantum barrier thickness (LQB) fixed at 7 nm, LQW was varied among the values 20 nm, 10 nm, 5 nm, 3 nm, and 2 nm. Figure 2a shows the normalized PL spectra of the corresponding structures. With a decrease in LQW, the PL peak position was systematically blue-shifted. Figure 2b displays the change in the PL peak position from 2.387 eV for a 20-nm-thick SQW to 2.399 eV for a 10-nm-thick SQW, 2.426 eV for a 5-nm-thick SQW, 2.483 eV for a 3-nm-thick SQW,


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and 2.528 eV for a 2-nm-thick SQW. This variation in the transition energy with LQW is attributed to quantum confinement, hence obeying the equation Ee-h = ħ2π2/2mL2, as described for a “particlein-a-box” model.38 In addition, the absorption spectra of these CsPbBr3 SQWs showed a spectralblue shift with a decrease in LQW (Figure 2c), although the rate of the blue-shifting decreases below an LQW of 5 nm due to the increase in the FWHM of the PL spectrum. The artificial MQWs have the advantage of multiple stacking with various types of bandgap structures. Figure 2d shows the change in integrated PL intensity observed by stacking 5-nm-thick quantum wells. As shown, 3- and 5-stacked CsPbBr3 (5 nm)/TPBi (7 nm) MQWs displayed integrated PL intensities 3.57 and 5.39 times higher than the intensity of the SQW. In contrast, the corresponding PL spectra showed a constant FWHM (120 meV ± 20 meV) with a similar PL peak position (2.50 eV ± 15 meV) regardless of the stacking number (inset of Figure 2d).

Figure 2. Optical properties of perovskite-based artificial MQWs. (a) PL spectrum of CsPbBr3/TPBi (7 nm) SQW as a function of LQW. (b) PL peak position and FWHM with variation in LQW. (c) Absorption spectra of SQW as a function of LQW. (d) PL spectra of MQWs with increasing stacking number of quantum wells.


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We also constructed asymmetric MQWs with the thickness of each well independently controlled. Namely, we fabricated TPBi (7 nm)/CsPbBr3 (5 nm)/TPBi (7 nm)/CsPbBr3 (15 nm)/TPBi (7 nm) asymmetric dual QW and TPBi (7 nm)/CsPbBr3 (3 nm)/TPBi (7 nm)/CsPbBr3 (10 nm)/TPBi (7 nm)/CsPbBr3 (20 nm)/TPBi (7 nm) asymmetric triple QW. Figure 3a-b depicts the PL spectra and the corresponding bandgap structures of the asymmetric dual QWs and triple QWs, respectively. Both MQWs exhibited multiple-wavelength emissions, resulting from the copresence of wells with different LQW values. The red, green, and blue dotted lines (E1,e and E1,h) signify the energy levels for carrier transition (QW1, QW2 and QW3, respectively) obtained by deconvoluting the multiple-wavelength emission spectra using Gaussian functions.

Figure 3. Optical properties of asymmetric perovskite-based artificial MQWs. (a) PL spectrum of TPBi (7 nm)/CsPbBr3 (5 nm)/TPBi (7 nm)/CsPbBr3 (15 nm)/TPBi (7 nm) dual QWs with the bandgap structure. (b) PL spectrum of TPBi (7 nm)/CsPbBr3 (3 nm)/TPBi (7 nm)/CsPbBr3 (10 nm)/TPBi (7 nm)/CsPbBr3 (20 nm)/TPBi (7 nm) triple QWs with the bandgap structure.

To gain a deeper understanding of the excitonic features in the artificial MQWs, we performed TDPL measurements to probe their emission characteristics. Figure 4a shows normalized TDPL


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contour maps of CsPbBr3 powder, a 20-nm-thick CsPbBr3 SQW, and a 3-nm-thick CsPbBr3 SQW. Figure S5 displays the TDPL spectra of these structures with the PL intensity varied. By increasing the temperature from 80 K to 300 K, a blue shift in the PL peak was observed for all structures, along with a decreasing PL intensity and increasing FWHM (Figure 4b). To estimate the exciton binding energy, one-channel Arrhenius plots were constructed using the following equation:40 I(T) = I0/(1+Aexp(Eb/kBT)), where I0 is the estimated integrated PL intensity as a scaling factor at 0 K, A is a constant related to the density of non-radiative recombination centers, kB is the Boltzmann constant, and Eb is the exciton binding energy determined from the TDPL intensities of CsPbBr3. The values of A and Eb calculated using the equation were 48.5 and 59.7 meV for the CsPbBr3 powder, 46.1 and 60.4 meV for the 20-nm-thick CsPbBr3 SQW, and 50.1 and 85.2 meV for the 3nm-thick CsPbBr3 SQW, respectively. In particular, the 3-nm-thick CsPbBr3 SQW showed a high exciton binding energy due to the spatially confined electronic states. In other words, the results indicate that these artificial quantum wells exhibit a tunable bandgap and a carrier confinement effect. The blue shift in the PL peaks, described by the temperature coefficient (α = dE/dT), can be explained by the interplay between the thermal expansion of the crystal lattice and electron-phonon renormalization.41 The value of α was calculated to be 0.077 meV/K for the 3D CsPbBr3 powder, 0.091 meV/K for the 20-nm-thick CsPbBr3 SQW, and 0.037 meV/K for the 3-nm-thick CsPbBr3 SQW. The broadening of the FWHM is largely determined by several phonon/carrier or impurity/carrier scattering mechanisms during the carrier recombination process.42 Interestingly, the 3-nm-thick CsPbBr3 SQW showed wide FWHMs, similarly to those of the CsPbBr3 powder, whereas the 20-nm-thick CsPbBr3 SQW displayed narrow FWHMs over the entire temperature range. This FWHM broadening of the PL spectrum was regularly observed by decreasing LQW (Figure 2b and 4b). Figure S6 shows top-view SEM images of CsPbBr3 films on TPBi/glass,


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which reveal small, isolated CsPbBr3 domains for a 3-nm-thick quantum well. We assumed that when the quantum-well thickness (LQW) is smaller than a few nanometers, isolated CsPbBr3 domains are formed, which eventually merge by lateral growth as the film thickness increases. This characteristic manifests as spatial fluctuation of LQW during the initial growth stage. Considering this quantum-well morphology, we assume that the quantum-well layer acts as a layer of quantum dots (QDs) with decreasing LQW. Thus, such regular FWHM broadening can be explained by carrier localization in QD-like potential minima owing to the spatial fluctuations of the layer forming the quantum well at very low LQW.43

Figure 4. Temperature-dependent PL (TDPL) of perovskite-based artificial MQWs. (a) Normalized TDPL contour maps of CsPbBr3 powder, TPBi (7 nm)/3D CsPbBr3 (20 nm)/TPBi (7 nm) SQW, and TPBi (7 nm)/3D CsPbBr3 (3 nm)/TPBi (7 nm) SQW. (b) Integrated PL intensities, PL peak positions, and


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FWHMs as functions of temperature. (c) Normalized TDPL contour maps of asymmetric dual QWs with various LQB values. The asymmetric dual QWs (LQW1: 20 nm and LQW2: 5 nm) were constructed using quantum barriers (QBs) with thicknesses of 14 nm, 7 nm, and 3 nm, each sandwiched between two asymmetric QWs.

Asymmetric dual QWs showed a tunneling effect depending on the value of LQB. By changing LQB from 15 nm (QW1) to 3 nm (QW2), the carrier transfer through the quantum barrier between two quantum wells was clearly modulated. Figure 4c and S7 shows the normalized TDPL contour maps in asymmetric dual QWs from 100 K to 300 K and their bandgap structures presenting the radiative recombination processes of QW1 (E1,e-1,h) and QW2 (E2,e-2,h), respectively. For an LQB value of 14 nm, the QW1 (lower energy) showed dominant PL intensity, while the PL from QW2 (higher energy) appeared only as a shoulder peak at both 100 K and 300 K; this finding implies that the most of carriers accumulated mainly in the QW1 over the tested temperature range. When LQB was decreased to 7 nm and 3 nm, dual emission with equivalent PL intensities, corresponding to QW1 and QW2, was observed above 200 K. It was also observed that the dominant PL intensity was shifted to QW2 as the temperature increases. Unusually, these perovskite-based asymmetric dual QWs exhibited the distinct carrier capture efficiency of the narrower QW2 instead of the wider QW1 with a variation in LQB. We assume that dominant carrier transition from the higher energy QW2 signifies the strong carrier localization effect in QD-like potential minima of QW2 owing to the spatial fluctuation and reflects the temperature-dependent nature of the phonon-assisted tunneling process in coupled asymmetric perovskite-based QWs.44, 45

Hot-carrier dynamics of perovskite-based artificial MQWs MQWs are well known for reducing hot-carrier cooling rates due to the enhanced “phonon bottleneck” in MQWs, making them excellent candidates for hot-carrier solar cell absorbers.46-48


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Femtosecond transient absorption (fs-TA) spectroscopy is one of the most convenient and robust approaches to probe and decipher the hot-carrier relaxation processes in photoactive materials.49, 50

Here, we have performed fs-TA measurements following above band-edge excitation at 320 nm

to directly probe the intra-band relaxation and the hot-carrier dynamics in both 6-stacked CsPbBr3 (5 nm)/TPBi (7 nm) MQWs and its CsPbBr3 (30 nm)/TPBi (7 nm) analog. Note, carrier confinement is expected to be much weaker in the latter (bulk) heterostructure. The experimental details of the fs-TA setup are reported elsewhere (see also the experimental section).49 Figure 5a-b shows the fs-TA spectra of CsPbBr3 (30 nm)/TPBi (7 nm) bulk heterostructure and 6-stacked CsPbBr3 (5 nm)/TPBi (7 nm) MQWs, respectively, in response to 320 nm optical excitation. Figure 5c-d depicts the band structure of these two structures. As can be seen, a strong groundstate bleaching (GSB) appeared at 520 nm and 506 nm for the CsPbBr3 (30 nm)/TPBi (7 nm) bulk heterostructure and 6-stacked CsPbBr3 (5 nm)/TPBi (7 nm) MQWs, respectively, which is consistent with the steady-state absorption spectra. Moreover, broad photo-induced absorption was predominantly observed for both samples over the range from 525 nm to 600 nm. Interestingly, we noted an increase in the GSB population and de-excitation of the photo-induced absorption of MQWs within a short time window of 5 ps. This observation could be attributed to the intraband relaxation and hot-carrier dynamics.51 In contrast, this phenomenon was not observed in the CsPbBr3 (30 nm)/TPBi (7 nm) bulk heterostructure when it was excited with the same energy and fluence. In addition to the high-energy excitation, we performed band-edge excitation to understand the nature of carrier relaxation in these structures. Similarly to the results of the above band-edge excitation, we observed GSB of both materials at exactly the same position (Figure S8). At the same time, the broad photo-induced absorption and the slow growth in the GSB signal of 6-stacked CsPbBr3 (5 nm)/TPBi (7 nm) MQWs were undetectable upon band-edge excitation.


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This observation supports the attribution of the appearance of the broad photo-induced absorption signal upon above band-edge excitation to the excited-state electrons, which were clearly ascribed to the quantized band structure. Importantly, due to the quasi-continuous band structure, both the increase in the bleaching signal and the appearance of the photo-induced absorption in response to excess energy were not observed for the CsPbBr3 (30 nm)/TPBi (7 nm) bulk heterostructure. To elucidate the effect of quantization, we compared the kinetic traces at GSB maxima of both samples in response to optical excitation at 320 nm (Figure 5e) and 475 nm (Figure S8c). In contrast to those of the bulk materials, the GSB kinetics of the MQWs showed a gradual rise in the GSB signal, which varied with the delay time in response to the 320 nm optical excitation with time constants of < 100 fs ± 20 (20%) and 2.1 ps ± 0.1 (80%). However, the GSB kinetics of the CsPbBr3 (30 nm)/TPBi (7 nm) bulk heterostructure also exhibited two time components: < 100 fs ± 20 (80%) and 0.5 ps ± 0.05 (20%). Interestingly, the kinetics at GSB maxima of those two materials showed a pulse-width-limited (< 100 fs) rise upon band-edge excitation. The slow rise in the GSB signal (intraband relaxation) of the MQWs upon above band-edge excitation clearly indicates the retardation of hot-electron cooling due to quantization of the electronic states of the QW structure. Figure 5f describes the energetic processes of the carrier density from the hotcarrier distribution (n(E)hot) to the equilibrium carrier distribution (n(E)eq) in CsPbBr3 (30 nm)/TPBi (7 nm) bulk heterostructure and 6-stacked CsPbBr3 (5 nm)/TPBi (7 nm) MQWs. The 6stacked CsPbBr3 (5 nm)/TPBi (7 nm) MQWs showed a modified DOS (gc(E)MQW) compared to that of CsPbBr3 (30 nm)/TPBi (7 nm) bulk heterostructure (gc(E)Bulk). We believe that the mechanism for hot-carrier relaxation is a modification in carrier diffusion due to the confinement effect of quantized electron states in MQWs. These results suggest that perovskite-based artificial MQWs significantly enlarge the window of opportunity for harvesting the excess energy of hot


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carriers, which may help overcome the Shockley–Queisser limit in future hot-carrier solar cells based on MQWs.

Figure 5. Hot-carrier dynamics in perovskite-based artificial MQWs. fs-TA spectra at different time delays in response to above-band-edge laser excitation of (a) CsPbBr3 (30 nm)/TPBi (7 nm) bulk heterostructure and (b) 6-stacked CsPbBr3 (5 nm)/TPBi (7 nm) MQWs. The grey dotted line in panel (b) shows the magnified (× 25) TA spectrum of TPBi at 510 fs in response to 320-nm optical excitation, demonstrating the negligible contribution of TPBi to the observed hot carrier dynamics. Schematic structures and bandgap alignments of (c) CsPbBr3 (30 nm)/TPBi (7 nm) bulk heterostructure and (d) 6stacked CsPbBr3 (5 nm)/TPBi (7 nm) MQWs. (e) Normalized kinetics traces monitored at GSB maxima of 30-nm-thick CsPbBr3 (520 nm) and MQWs (506 nm) in response to above-band-edge excitation (λex = 320 nm) and (f) their E-k diagrams.

In summary, we have introduced the fabrication of perovskite-based artificial MQWs. Unlike III-V compound MQW heterostructures, these perovskite-based MQWs do not require certain substrates as epitaxial thin films, and they can be easily fabricated by conventional thermal evaporation. Our experimental results demonstrated that these perovskite-based artificial MQWs exhibit the desired quantum well structure with the precise tuning of the transition energies. We


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showcased various MQWs with elaborately designed bandgaps that exhibit quantum-confined carrier transitions, multiple-wavelength emission, quantum tunneling, and suppression of hotcarrier cooling rates. The ability to construct asymmetric MQWs, with stacked quantum wells of different thicknesses, makes perovskite-based materials amenable to nanoscale bandgap engineering that generates intentionally unique carrier distributions in the materials. The defect tolerance and fabrication versatility of halide perovskite-based heterostructures can resolve many of the bottlenecks that currently limit the proliferation of traditional semiconductor heterostructure technologies.

Experimental section CsPbBr3 perovskite preparation. CsPbBr3 perovskite was synthesized by the inverse temperature crystallization (ITC) method. A mixture of 10 mmol CsBr and 20 mmol PbBr2 was dissolved in 10 ml DMSO and stirred at room temperature overnight. Then, the mixture was gradually heated from 60 °C to 100 °C, yielding a green precipitate at the bottom of the vial. After reaching 100 °C, the supernatant was gently decanted into another vial and further heated to 120 °C. The orange crystals of CsPbBr3 produced during this process were carefully extracted, rinsed with hot DMSO, and then successively dried using filter paper and a vacuum. The crystals were pulverized into a fine powder before thermal evaporation. Secondary-ion mass spectrometry (SIMS). SIMS experiments were performed on a Dynamic SIMS instrument from Hiden Analytical Company (Warrington, UK) operated under ultra-high vacuum conditions, typically 10-9 Torr. A continuous Ar+ ion beam with an energy of 2.5 keV was employed to sputter the surface while the selected ions ascribed to Br-, Cs+, Pb+, CN, and Si+ were sequentially collected using a MAXIM spectrometer equipped with a quadrupole analyzer. Ions


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were collected from the sample by a shaped extraction field and energy filtered using a parallelplate system, with the energy resolution matched to that of the quadrupole analyzer. After passing through a triple-filter system, detected ions were measured using a pulse-counting detector with a 4 keV post-acceleration potential to further increase the detection efficiency at high masses. To avoid the edge effect for the acquisition of the depth-profiling experiments (profile curves highlighting the ion count vs. the sputtering depth), data were extracted from a small area (typically 75 × 75 µm2) centered in the middle of the sputtered area (estimated to be 750 × 750 µm2) using adequate electronic gating. In contrast, 3D mapping was performed by scanning the ion beam over the sputtered surface area divided into 100 × 100 pixels and collecting ion signals from each pixel as a function of depth. The sputtering time was converted to sputtering depth by measuring the depth of the crater generated at the end of the depth-profiling experiment using a stylus profiler from Veeco. X-ray photoelectron spectroscopy (XPS). XPS measurements were carried out using a Kratos Axis Supra DLD spectrometer with an Al Kα source (λν = 1486.6 eV). The measured binding energies were referenced to the C 1s binding energy of carbon contamination (284.8 eV). The valence-band maximum was calculated by extrapolating the leading edge to zero signal. Three samples were prepared for the XPS measurements: ~50-nm-thick CsPbBr3 on glass, ~50-nm-thick TPBi on glass, and ~5-nm-thick CsPbBr3 on ~50-nm-thick TPBi on glass. Transmission electron microscopy (TEM). TEM lamellas were prepared on a FEI Helios NanoLab 400S FIB/SEM dual-beam system equipped with a Ga+ ion source. Pt layers were deposited on the surface region of interest by an electron and ion beam for sample protection. Samples were thinned to a relative thickness of 80 nm using progressively decreasing ion-beam energies in the FIB down to 5 keV.


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Temperature-dependent photoluminescence (TDPL). The TD-PL spectra were characterized using a Horiba JY LabRAM Aramis spectrometer with an Olympus 50× lens in a Linkam THMS600 stage. A 473 nm laser was used as the excitation source. Femtosecond transient absorption measurement. Femtosecond transient absorption (fs-TA) spectroscopy measurements were performed on timescales of 0.1 ps to 6 ns using a regeneratively amplified Ti:sapphire laser (800 nm laser pulses with a 35 fs pulse width at a 1-kHz repetition rate) in conjunction with a Helios spectrometer. Excitation-pump pulses at 320 and 475 nm were generated after passing through a fraction of an 800 nm beam into the spectrally tunable (240−2600 nm) optical parametric amplifier (Newport Spectra-Physics). The pump fluence of the excitation laser source was adjusted by using a neutral-density (ND) filter to avoid the generation of multiple charge carriers. The probe pulses (UV-visible and NIR wavelength continuum, white light) were generated by passing another fraction of the 800 nm pulses through a 2 mm thick calcium fluoride (CaF2) crystal. Before white-light generation, the 800 nm amplified pulses were passed through a motorized delay stage. Depending on the movement of the delay stage, the transient species were detected following excitation at different time scales. The white light was split into two beams (labeled the signal and reference beams) and focused on two fiber optics to improve the signal-tonoise ratio. The excitation-pump pulses were spatially overlapped with the probe pulses on the samples after passing through a synchronized mechanical chopper (500 Hz), which blocked alternative pump pulses. The absorption change (∆A) was measured with respect to the time delay and wavelength. All spectra were averaged over a period of 2 s for each time delay. X-ray diffraction (XRD). The XRD patterns were collected using a Bruker AXS D8 diffractometer for powder using Cu-Kα radiation with a non-ambient tool suitable for temperature-dependent XRD measurements. XPS studies were carried out in a Kratos Axis Supra spectrometer equipped


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with a monochromatic Al Ka X-ray source (hν = 1486.6 eV) operating at 150 W, a multi-channel plate, and a delay line detector under a vacuum of 10-9 mbar. All spectra were recorded using an aperture slot of 300 μm × 700 μm. Survey spectra were collected using a pass energy of 160 eV and a step size of 1 eV. A pass energy of 20 eV and a step size of 0.1 eV were used for the highresolution spectra. ASSOCIATED CONTENT Supporting Information XPS survey spectra, steady-state PL spectra of CsPbBr3 perovskites, XRD patterns of CsPbBr3, depth profiling of CsPbBr3 perovskites-based artificial multiple quantum wells obtained using SIMS, temperature-dependent PL spectra of CsPbBr3 perovskites, top-view SEM images of CsPbBr3 quantum well, and bandgap structure for asymmetric dual QWs, fs-TA survey spectra AUTHOR INFORMATION Corresponding Author *E-mail: Osman M. Bakr ([email protected]) Author Contributions O.B. supervised the work. O.B., K.J., B.T., N.C., J.S., and S.O. designed and discussed the experiments. A.Z., R.N., and L.S. synthesized and analyzed the perovskite-based materials. K.J. and N.M. fabricated the artificial perovskite-based MQWs. C.H., R.C., and B.O. conducted the optical measurements. O.M., P.M., and I.D. obtained and interpreted the transient absorption results. A.A. performed XRD. N.W performed SIMS. M.H. obtained and analyzed the XPS results. All authors critically evaluated the manuscript.


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ACKNOWLEDGMENT The authors gratefully acknowledge the financial support provided by King Abdullah University of Science and Technology (KAUST). This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korea government (NRF-2017R1C1B5017953).

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Perovskite-Based Artificial Multiple Quantum Wells - Nano Letters

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