Van der Waals epitaxy of type-II band alignment CsPbI₃/TMDC heterostructure for optoelectronic applications

Abstract

Van der Waals epitaxy allows heterostructure formation without considering the lattice match requirement, thus is a promising method to form 2D/2D and 2D/3D heterojunction. Considering the unique optical properties of CsPbI₃ and transition metal dichalcogenides (TMDCs), their heterostructure present potential applications in both photonics and optoelectronics fields. Here, we demonstrate selective growth of cubic phase CsPbI₃ nanofilm with thickness as thin as 4.0 nm and Zigzag/armchair orientated nanowires (NWs) on monolayer WSe₂. Furthermore, we show growth of CsPbI₃ on both transferred WSe₂ on copper grid and WSe₂ based optoelectrical devices, providing a platform for structure analysis and device performance modification. Transmission electron microscopy (TEM) results reveal the epitaxial nature of cubic CsPbI₃ phase. The revealed growth fundamental of CsPbI₃ is universal valid for other two-dimensional substrates, offering a great advantage to fabricate CsPbI₃ based van der Waals heterostructures (vdWHs). X-ray photoelectron spectroscopy (XPS) and optical characterization confirm the type-II band alignment, resulting in a fast charger transfer process and the occurrence of a broad emission peak with lower energy. The formation of WSe₂/CsPbI₃ heterostructure largely enhance the photocurrent from 2.38 nA to 38.59 nA. These findings are vital for bottom-up epitaxy of inorganic semiconductor on atomic thin 2D substrates for optoelectronic applications.

Graphical abstract

Keywords

van der Waals epitaxy / band alignment / growth fundamental / charge transfer / photodetector

1 Introduction

Thanks to the discovery of two-dimensional materials with fruitful unique properties, such as boron nitride (h-BN), TMDCs and graphene, it unlocked new opportunities for next generation of electronic and optoelectronic devices such as photodetectors [1,2], transistors [3,4], light emitting diodes [5], lasers [6-8], and photovoltaic devices [9,10]. Van der Waals heterostructures (vdWHs) allow easy tuning of the optical and electrical properties that are desired for specific applications. First, the formation of vdWH is a promising strategy to overcome the drawback of weak light absorption in monolayer TMDCs. Tsang et al. [11] designed a PtSe₂/CdTe heterostructure, the increased light absorption together with charge transfer improved the responsivity and on/off current ratio to 506.5 mA/W and 7 × 10⁶. Secondly, vertical TMDC heterostructures with twist angle create an interfacial charge transfer pathway and allow formation of interlayer excitons with longer lifetime, thus attracting extensive attentions [12]. Thirdly, vdWH is able to enhance the Förster-type energy transfer efficiency. Kozawa et al. [13] exhibited efficient and ultrafast (~1 ps) fluorescence resonance energy transfer (FRET) in a closely coupled WS₂/MoSe₂ heterojunction, which can be applied for energy harvesting. For the purpose to improve the light absorption, inorganic metal-halide perovskite is an ideal candidate showing high light absorption coefficients, high optoelectronic conversion efficiency, thus have been widely used for low-cost and efficient photonic, photovoltaic and optoelectronic devices [14]. Benefiting from these distinct advantages, inorganic metal-halide perovskites are utilized in the formation of heterojunctions with TMDCs, resulting in a profound alteration of their optical properties as well as enhancements in device performance, such as photoelectric detection [15,16] and light-harvesting [17].

Currently, top-down transferring is widely used to prepare inorganic metal-halide perovskite/TMDCs heterojunctions thanks to its advantages in rapid fabrication of different heterojunctions, such as quantum dots(QDs) [18] and NWs [19] on TMDCs. This method is suitable to study the optical and electronic properties of the heterostructure as well as prototype device fabrication [20‒22]. By utilizing a hybrid structure consisting of metal-halide perovskites and TMDCs, it becomes possible to activate distinct nonradiative exciton relaxation (NRER) channels at the interface. Jiang et al. [18] demonstrated that CsPbBr₃ QDs/monolayer WS₂ forms a type-I band alignment, showing efficient and charge transferring process which can boost the photodetection efficiency. Moreover, MoS₂/CsPbBr₃ heterostructure based device shows efficient charge separation at the heterojunction interface and thus enhanced photoelectric conversion efficiency [23]. Despite the great achievement of top-down transferring approach, their applications are limited by lack of capability in scalable production. On the contrary, van der Waals epitaxy is another method of fabricating 2D material based heterostructure. Differing from traditional epitaxy, this method relies on weak van der Waals force between the 2D substrate and the epitaxial 2D or 3D layer, thus it does not need to satisfy lattice-match requirement during traditional epitaxy method [24]. Moreover, employing van der Waals epitaxy permits adjustment of crystal structure, composition and thickness of the heterostructure, further refining its photoelectric performance. Consequently, These advantages make it highly suitable for heterojunction integration, thus gathering significant attentions [25‒32], such as III‒V semiconductor on graphene/mica [24,33], perovskite NWs/nanoplate on mica [34], ZnO [35], and Ga₂O₃ on mica [36]. During these studies, graphene and mica are mostly applied thanks to their easy availability as well as good heat tolerance. TMDCs are also applied for van der Waals epitaxy of free standing In_xGa_1−xAs NWs [37]. This method can enable large-scale, patterned high-quality 2D material/inorganic semiconductor heterojunction formation, rendering the van der Waals heterojunction field particularly attractive. Indeed, wafer scale high quality GaAs and GaN related heterostructure have been applied to grown on graphene (h-BN) and applied for flexible as well as vertical full-color micro-LED applications [33,38]. Despite of these advantages, the development of van der Waals epitaxy still needs more research. Firstly, more growth fundamental should be revealed to understand its growth mechanism. Secondly, previous results mainly focus on the van der Waals epitaxy itself instead of the properties of the formed heterostructure. Thirdly, even though van der Waals epitaxy of metal-halide perovskite on mica has been reported [39], no attempts on van der Waals epitaxy of CsPbX₃ on TMDCs have been reported.

Considering the importance of van der Waals epitaxy in heterostructure formation and potential application in the photonic and electronic fields between TMDCs and inorganic perovskite semiconductor, we aim to bridge the gap by demonstrating the growth of CsPbI₃ on TMDCs layers via the chemical vapor deposition (CVD) method. We show that CsPbI₃ can selectively grow on WSe₂ within a wide range of growth parameters with morphology of either nanofilm or NWs. The appropriate growth temperature and precursor concentration are critical parameters for controlling the growth mode of CsPbI₃, deepening the van der Waals growth fundamental of CsPbI₃. By tuning the growth parameters, nanofilms with adjustable thicknesses from 4.0 to 11.6 nm and NWs with Zigzag and armchair orientations can be formed, and both of which exhibit epitaxial relationships with the growth substrate. Meantime, the crystal growth prefers to form a cubic phase on WSe₂ substrate and does not damage the underlying fragile monolayer WSe₂. Additionally, high resolution XPS analysis confirms the formation of the CsPbI₃/WSe₂ heterojunction, which exhibits a type-II band alignment characterized by a valence band offset (VBO) of 0.46 ± 0.10 eV and a conduction band offset (CBO) of 0.34 ± 0.10 eV. Optical characterizations reveal that CsPbI₃ and monolayer WSe₂, MoS₂, WS₂, and MoSe₂ form a type-II band alignment, which exhibits efficient and ultrafast charge transport. This hybrid heterostructure induces a lower energy emission peak at the shoulder of pure TMDC emission peak at low temperature. Furthermore, the revealed growth mechanism is also applicable to other two-dimensional materials, including but not limited to MoSe₂, MoS₂, WS₂, and h-BN. Furthermore, we successfully grow CsPbI₃ film directly on monolayer WSe₂ based photodetector devices. The growth of CsPbI₃ notably enhance the photocurrent from 2.38 to 38.59 nA and the on/off ratio from 3.14 to 16.37. Therefore, our growth findings pave the way for integrating inorganic metal-halide perovskites and 2D materials in a controllable and scalable manner through bottom-up approach.

2 Experimental methods

2.1 CsPbI₃/TMDCs heterostructure formation

First, monolayer TMDC was grown on Si/SiO₂ substrate via physical vapor deposition (PVD) method in a single-zone tube furnace (Jiusuo Technology) equipped with 1-inch diameter quartz tube [40]. Then, heterostructure formation is realized by growing CsPbI₃ in a dual-zone tube furnace with 1-inch diameter quartz tubes. A mixture of cesium iodide powder (40 μmol, 99%, Sigma-Aldrich) and lead iodide powder (40 μmol, 98%, Sigma-Aldrich) was used as precursor for CsPbI₃ growth and placed in the heating zone of the furnace. Before heating, the furnace was pumped to base pressure about 0.01 mbar and then a mixed Ar/H₂ with ratio of 19:1 was used as the carrier gas at a flow rate of 200 sccm until the pressure was stabilized at 400-500 mbar. The as-grown TMDC layer on SiO₂/Si substrate was put approximately 4 inches downstream from the precursor. The pressure of the furnace during the growth was controlled by a mechanical pump and measured by vacuum gauge. During the growth, the temperatures of the precursor was set at 660 °C while the growth temperature of the substrate was changed from 170−330 °C. The CVD reaction was carried out for 5 minutes followed by natural cooling down of the furnace. The standard growth temperature, precursor weight and pressure are 230 °C, 80 μmol precursor mixture and 400 mbar, respectively, since it leads to uniform CsPbI₃ film growth.

2.2 Device fabrication and measurement

The obtained WSe₂ films were transferred onto Si/SiO₂ substrates with PDMS thin films. The Au electrodes were fabricated on WSe₂ by photolithography and electron-beam evaporation, leading to a channel width of 2‒10 μm and an electrode width of 3 μm. Subsequently, the fabricated WSe₂ device was loaded into the furnace for CsPbI₃ film growth. The electronic measurements were carried out in a probe station, and data were recorded with a semiconductor device analyzer (Agilent-B1500A).

2.3 Optical characterizations

The morphology of CsPbI₃/TMDCs heterostructures were characterized using optical microscopy, scanning electron microscopy (SEM, Tescan Mira4) and atom force microscopy (AFM, Agilent 5500 AFM/SPM). HRTEM images were recorded using a JEM-F200 microscope operated at an accelerating voltage of 80 kV to avoid damage to the CsPbI₃. For HRTEM characterization studies, WSe₂ was first transferred to a copper grid by a typical KOH-assisted wet-transfer method [41]. Then, CsPbI₃ was grown on the transferred WSe₂ before TEM analysis. XPS analyses (Thermo Scientific K-Alpha) were used to evaluate the chemical compositions and valence states of the as-grown samples. All Raman and photoluminescence (PL) measurements at room temperature were performed using a Renishaw in Via Reflex system where a 532 nm laser was focused onto the sample through a ×50 objective lens, resulting in a spot diameter of smaller than 1 μm. For temperature dependent PL measurements, LINKAM’s cryogenic controller and probe stand was connected to the PL system and the laser spot was focused on the sample through a long working distance ×50 objective lens. Time-resolved photoluminescence (TRPL) were measured in a home-built micro-PL system where 532 nm pulsed laser was focused on the sample via a ×100 objective lens and the decay signal was detected by a single photon detector connected to a Picoharp 300 system [42].

3 Results and discussion

Fig.1 shows the successful growth of CsPbI₃ nanofilm and NWs on monolayer WSe₂. Growth temperature plays an important role in determining the morphology of CsPbI₃. At low temperature, such as 230 °C, CsPbI₃ nanofilms grow selectively on WSe₂ monolayer while no growth is observed on the SiO₂ surface. Increasing temperature to 260 °C leads to the formation of a high density of CsPbI₃ NWs. Further increasing the temperature to 330 °C, the surface adatom concentration of CsI and PbI₂ on WSe₂ becomes lower than that required to overcome the nucleation barrier, thus no CsPbI₃ crystal could be observed. However, due to the presence of dangling bonds, there is a lower nucleation barrier at the edge of WSe₂, allowing for nucleation and formation of CsPbI₃ triangular rings [see Fig.1(c)]. Further increasing the temperature results in no CsPbI₃ nucleation on the whole substrate. It is worth mentioning that during the whole investigated temperature range from 170 °C to 330 °C [see Fig. S1 in the Supporting Information (SI)], CsPbI₃ prefers to selectively grow on the WSe₂ instead of the amorphous SiO₂ substrate.

Precursor concentration is another important parameter determines the growth of CsPbI₃. At standard film growth conditions, either lower or double the precursor weight does not change the morphology of as-grown CsPbI₃ except for the film thickness. Indeed, AFM measurements in Fig.1(h) demonstrate that the film thickness can be tuned from around 5 nm to ~12 nm by changing the precursor weight. Energy Dispersive X-Ray Spectroscopy (EDX) mapping shows the uniform distribution of Pb element on the WSe₂ flake, again confirming the formation of CsPbI₃ film. It is worth mentioning that, the CsPbI₃ thickness can be manipulated by changing the growth time as well (see Fig. S2 in the SI). From both inset SEM images in Fig.1(a)‒(c) and AFM images, there exist CsPbI₃ islands on the WSe₂, indicating that the as-grown CsPbI₃ film is formed by multi nucleation and growth process. Despite of the island formation, our method can selectively grow CsPbI₃ on WSe₂ with tunable thickness at nanometer scale, which is important to study the optical performance of the heterostructure. Surprisingly, further elevating the precursor weight ratio to threefold leads to NW formation against thick film. These formed NWs under two different growth conditions are slightly different. In general, CsPbI₃ NWs prefer to grow along either Zigzag or armchair direction of the WSe₂, as illustrated in Fig.1(i). At high growth temperature, nearly all the NWs grow along the Zigzag direction with a height of around 47 nm. In comparison, the aligned direction of NWs grown at lower temperature but higher precursor weight is different. Statistically, around 68% (32%) of NWs grow along the armchair (Zigzag) direction. The NW height is measured to be around 48.3 nm. In addition to growth temperature and precursor weight, the growth pressure affects the CsPbI₃ growth trend as well (see Fig. S3 in the SI). Maintaining a suitable pressure is a key for successful CsPbI₃ film growth.

The observed growth evolution of CsPbI₃ is explained in Fig.2. Supersaturation of CsPbI₃, which is determined by the growth temperature and the localized concentration on the WSe₂, plays a crucial role in determining the nucleation process. The growth behavior of CsPbI₃ is related to the adsorption, diffusion and desorption of CsI and PbI₂ adatoms on the WSe₂. These three processes compete with each other. According to the literature [43], the low adsorption energy of CsI on TMD leads to a higher precursor concentration on the TMD than the SiO₂ substrate, thus resulting selective growth of CsPbI₃ on WSe₂. At low temperatures, the desorption rate of adatoms on the surface of WSe₂ is low. Consequently, the high precursor concentration on the surface exceeds what is required to overcome the nucleation barrier and leads to a high nucleation density on the whole WSe₂ flake. During growth, these nuclei continue to grow and expand, forming large polycrystal nanofilms on WSe₂. According to the well-known van der Waals growth mechanism, heterostructure with the smallest lattice match is easy to form, since the interface strain caused by lattice mismatch can be effectively reduced [44], thereby minimizing the interfacial energy. During the growth, when lattice mismatch along one certain direction is small but quite large when perpendicular to this direction, crystal thus prefers to grow along one direction to minimize the interfacial energy, thus leading to the formation of NWs [45]. Normally high precursor diffusion rate allows growth of crystal with the smallest lattice mismatch [46]. According to the calculations, lattice mismatch between CsPbI₃ and WSe₂ is quite small along Zigzag and armchair directions, which is 0.2% and 0.5% respectively, as shown in Fig. S4. In comparison, lattice mismatch perpendicular to Zigzag (2.3%) and armchair direction (1.3%) is quite large. At high temperatures, the elevated temperature can increase the diffusion rate of precursor and reduce the adatom concentration on the WSe_2, explaining the reduced nucleation density. Consequently, crystal prefers to grow with minimum lattice mismatch or interfacial energy. Based on our calculation, CsPbI₃ along Zigzag direction shows the smallest lattice match, therefore NWs are formed along Zigzag direction. At high precursor weight, it is surprising to see the formation of NWs but this trend has been double checked. It is probably that the adatom diffusion rate and desorption rate of CsPbI₃ precursor on the WSe₂ surface is chemical potential dependent. Since the NW along armchair and Zigzag direction both show quite small lattice mismatch, the large driving force under high precursor weight allows NWs to be formed either along Zigzag or armchair directions.

To verify the microstructure of perovskite on TMDC, we successfully developed a method to allow growth of CsPbI₃ on WSe₂ transferred to a copper grid. From TEM analysis, both NW like and film like CsPbI₃ can be observed. For the NW like shape [see Fig.3(a)], HRTEM show the lattice distance of (200) plane is 3.15 Å, agreeing with reported value for cubic CsPbI₃ structure [47]. Moreover, the selected area electron diffraction (SAED) presents the same diffraction pattern as cubic phase CsPbI₃ along [001] zone axis. The SAED together with the HRTEM confirms that the as-grown CsPbI₃ NW is cubic structure and grow along [200] direction. Further EDX mapping and EDX spectrum [see Fig.3(d, e)] again confirm the uniform distribution of Cs, Pb and I elements, verifying the successful formation of CsPbI_3. For the CsPbI₃ nanofilm in Fig.3(f), nanofilm is grown along the [111] direction, the corresponding SADP present hexagonal shape. The measured (0\overline{1}1) plane lattice distance is 4.41 Å, agreeing with the previous results. Moreover, from the DP of CsPbI₃ and underline WSe₂, the (0\overline{1}1) plane of CsPbI₃ is parallel to the (1\overline{1}0) direction of WSe₂, demonstrating the epitaxial nature of CsPbI₃ nanofilm. Figure S5(a) shows the XRD patterns of CsPbI₃ nanostructure. Sharp diffraction peaks at 14.1°, 40.7° and 45.5° can be observed, which are ascribed to the (100), (220) and (310) planes of the cubic CsPbI₃ based on the literatures [48]_. The measured XRD results agree with our TEM measurement in Fig.3(c). AFM measurements of the as-grown WSe₂ flake show the thickness of around 1 nm, demonstrating that the used WSe₂ for CsPbI₃ growth is actually monolayer (see Fig. S5 in the SI).

The band alignment of the CsPbI₃/WSe₂ heterostructure was determined using XPS, as shown in Fig.3(i)‒(l). The VBO is primarily derived by evaluating the energy discrepancy between the core levels of W 4f and Pb 4f and their respective energies relative to the valence band maximum (VBM) of CsPbI₃ and WSe₂. Consequently, the VBO of the heterostructures (denoted as ΔE_V) can be accurately calculated by applying the following equation [49]:

(1)\begin{array}{rl}\mathrm{\Delta }{E}_{\mathrm{V}}=& \left(E_{\mathrm{W}4\mathrm{f}}^{{\mathrm{W}\mathrm{S}\mathrm{e}}_2}-E_{\mathrm{V}\mathrm{B}\mathrm{M}}^{{\mathrm{W}\mathrm{S}\mathrm{e}}_2}\right)\\ & +\left(E_{\mathrm{P}\mathrm{b}4\mathrm{f}}^{{\mathrm{W}\mathrm{S}\mathrm{e}}_2/{\mathrm{C}\mathrm{s}\mathrm{P}\mathrm{b}\mathrm{I}}_3}-E_{\mathrm{W}4\mathrm{f}}^{{\mathrm{W}\mathrm{S}\mathrm{e}}_2/{\mathrm{C}\mathrm{s}\mathrm{P}\mathrm{b}\mathrm{I}}_3}\right)\\ & -\left(E_{\mathrm{P}\mathrm{b}4\mathrm{f}}^{{\mathrm{C}\mathrm{s}\mathrm{P}\mathrm{b}\mathrm{I}}_3}{-E}_{\mathrm{V}\mathrm{B}\mathrm{M}}^{{\mathrm{C}\mathrm{s}\mathrm{P}\mathrm{b}\mathrm{I}}_3}\right).\end{array}

The first component of Eq. (1) is obtained by subtracting the energy difference between the W 4f_7/2 core level (32.35 ± 0.10 eV) and the VBM of WSe₂ [Fig.3(i)]. The VBM of WSe₂ is acquired by extrapolating the linear fit to the leading edge of the valence band spectra, reaching the base level with a value of 0.52 eV. Consequently, the first component of the equation is determined to be 31.83 ± 0.10 eV. Correspondingly, the energy difference between the Pb 4f_7/2 core level (138.37 eV) and the VBM of CsPbI₃ (0.76 eV) is measured to be 137.61 ± 0.10 eV [Fig.3(j)], representing the last component of the equation. The second part of the equation is calculated by computing the energy difference between the Pb 4f_7/2 and W 4f_7/2 core levels, with a value of 106.24 ± 0.10 eV [Fig.3(k)]. Based on these results, the VBO of the CsPbI₃/WSe₂ heterojunction is determined to be 0.46 ± 0.15 eV.

The CBO (ΔE_C) can be then determined as follows:

(2)\mathrm{\Delta }{E}_{\mathrm{C}}=E_{\mathrm{g}}^{{\mathrm{W}\mathrm{S}\mathrm{e}}_2}+\mathrm{\Delta }{E}_{\mathrm{V}}-E_{\mathrm{g}}^{{\mathrm{C}\mathrm{s}\mathrm{P}\mathrm{b}\mathrm{I}}_3},

where E_{\mathrm{g}}^{{\mathrm{W}\mathrm{S}\mathrm{e}}_2} and E_{\mathrm{g}}^{{\mathrm{C}\mathrm{s}\mathrm{P}\mathrm{b}\mathrm{I}}_3} are the band gap of WSe₂ and CsPbI₃, respectively. The band gap is 1.72 eV for CsPbI₃ [50] and 1.60 eV for the WSe₂ [19] substrate, respectively. Based on Eq. (2), the CBO of CsPbI₃/WSe₂ heterojunction is found to be 0.34 ± 0.10 eV. The extracted band offset as shown in Fig.3(l), demonstrating the type-II band alignment in the heterostructure.

The formation of both CsPbI₃ nanofilm and crystal orientation tunable NWs provides a good platform to study the optical performance of this hybrid heterostructure. CsPbI₃/WSe₂ forms a type-II band alignment [51], providing an internal electric field to separate the photogenerated electron and holes. Raman scattering comparison of WSe₂ before and after CsPbI₃ growth show unobserved difference in characteristic peak intensity, width nor peak position [see Fig.4(a)], suggesting that van der Waals epitaxy process does not cause damage to the underline monolayer WSe₂ substrate thanks to the low growth temperature. For the CsPbI₃ NW/WSe₂ as well as the CsPbI₃ nanofilm/WSe₂ heterostructure, two emission peaks corresponding to CsPbI₃ (1.78 eV) and WSe₂ (1.60 eV) can be observed with reduced emission intensity [see Fig.4(b)]. Quantitatively, the emission intensity of CsPbI₃ is reduced over ten times. The slight emission energy shift as well as emission intensity difference between CsPbI₃ NW and nanofilm is ascribed to the crystal quality and thickness differences. In addition, after removing of the top CsPbI₃ by water washing, the emission energy as well as emission intensity is nearly unchanged, indicating that CsPbI₃ growth does not cause damage to the WSe₂ layer (see Fig. S7 in the SI). The corresponding TRPL decay curves in Fig.4(c) agree with the observed PL spectra changes. The PL lifetime of isolated CsPbI₃ can be accurately modeled by a single-exponential function, indicating an exciton radiative recombination process [52]. The PL lifetime of pure CsPbI₃ NW 7.90 ns, demonstrating a high crystal quality of as-grown CsPbI₃ NWs [40]. After forming a heterostructure, PL lifetime is reduced to 4.24 ns, which is also ascribed to an exciton radiative recombination process. The charge transfer efficiency ({\varphi }_{\mathrm{E}\mathrm{T}}) and charge transfer rate (K_{\mathrm{E}\mathrm{T}}) from CsPbI₃ to WSe₂ can be calculated using the following equations [53,54]:

(3){\varphi }_{\mathrm{E}\mathrm{T}}=1-{\displaystyle \frac{{\tau }_{\mathrm{H}}}{{\tau }_{\mathrm{N}\mathrm{W}}}},

(4)K_{\mathrm{E}\mathrm{T}}={\displaystyle \frac{1}{{\tau }_{\mathrm{H}}}}-{\displaystyle \frac{1}{{\tau }_{\mathrm{N}\mathrm{W}}}},

where {\tau }_{\mathrm{N}\mathrm{W}} is the PL lifetime of CsPbI₃ NW on the SiO₂ substrate and {\tau }_{\mathrm{H}} is the PL lifetime of CsPbI₃ NW on the monolayer WSe₂. Based on the above equations, K_{\mathrm{E}\mathrm{T}} and {\varphi }_{\mathrm{E}\mathrm{T}} are calculated to be 0.11 × 10⁹ s^–1 and 46%, respectively. The PL lifetime of monolayer WSe₂ can be effectively modeled using a biexponential function, revealing a many-body effect lifetime (τ₁) of 279 ps and an exciton radiative recombination lifetime (τ₂) of 420 ps. After forming a heterostructure, the average PL lifetime of WSe₂ is also dropped from 284 to 94 ps. K_{\mathrm{E}\mathrm{T}} and {\varphi }_{\mathrm{E}\mathrm{T}} are calculated to be 0.71 × 10⁸ s^–1 and 67%, respectively. Therefore, we believe that the observed decreased PL intensity as well as PL lifetime are explained by the fast and efficient charge transfer process due to the formed type-II band alignment.

It is worth mentioning that the observed PL and TRPL trend is valid for both CsPbI₃ NW and nanofilm [54]. We further performed PL intensity mapping at 1.64 eV on a CsPbI₃/WSe₂ heterojunction. Before measurement, the as-grown CsPbI₃ film is partially removed by scratching with a tweezer, exposing the underline WSe₂ [see Fig.4(d)]. The PL intensity of the WSe₂ at the crack location is much brighter while the other region shows weak and uniform emission. This emission distribution confirms that the CsPbI₃ film does not reduce the quality of WSe₂ and the reduction of PL emission is definitely a result of heterostructure formation. The inherent characteristics of the two-dimensional structure, combined with the heightened Coulomb interaction, give rise to profound excitonic effects within the monolayers. Furthermore, the Coulomb interaction facilitates the preferential formation of trions alongside excitons, which frequently exhibit dominance in the PL spectra in the presence of freely mobile charges [55]. Further temperature dependent PL experiments are carried out from 300 K to 93 K. During cooling down process, the emission peak of CsPbI₃ gradually becomes stronger and the emission energy undergoes a redshift (Fig. S8 in the SI), agreeing with previous reports [56‒59]. Below 93 K, the emission intensity of CsPbI₃ is so strong that the emission peak from WSe₂ becomes hard to distinguish. In comparison, the emission evolution of WSe₂ peak at heterojunction and pure WSe₂ flakes are compared in Fig.4(f)‒(h). In general, the emission of WSe₂/CsPbI₃ has a lower emission energy and much broader emission peak. During cooling down, the emission peak undergoes a blueshift with enhanced emission intensity. For the CsPbI₃ film/WSe₂, it presents a broad emission peak with energy linearly shifted from 1.58 to 1.65 eV during 300 to 93 K, which is around 12 meV smaller than the trion emission of WSe₂. For the CsPbI₃ NW/WSe_2, a low energy peak occurs at temperature below 273 K and becomes stronger during cooling down. Since the laser spot diameter is larger than the NW diameter, the emission consists of pure WSe₂ (higher energy peak) and heterostructure (lower energy shoulder). The exact position of the low energy peak is sample dependent and is always lower than the trion emission of WSe₂ [see Fig.4(h)]. Moreover, after removing of the CsPbI₃ by water washing, no such a low energy emission peak is observed (see Fig. S9 in the SI), demonstrating that this low energy peak is not caused by the doping or damage of the TMDCs during CsPbI₃ epitaxy. Consequently, this emission is ascribed to the band alignment of the CsPbI₃/WSe₂ heterostructure.

To test whether the revealed van der Waals growth fundamental of CsPbI₃ is only valid for WSe₂ or universal valid for other 2D materials, we tried to grow CsPbI₃ nanofilm and NWs on monolayer WS_2, MoS₂ and MoSe₂ at the same growth conditions as those on monolayer WSe₂. As displayed by the optical images in Fig.5, CsPbI₃ nanofilms are selectively formed on TMDCs under the same growth conditions. For CsPbI₃ NW/TMDC heterojunction, NWs are formed on TMDCs with different density and length. The general geometric relationship between the NW and the TMDCs layer is the same, most of the NWs grow along the Zigzag direction with the rest along armchair directions. Indeed, we also successfully grow CsPbI₃ on multilayer WSe₂, transferred WSe₂ as well as multilayer h-BN using the standard film growth condition (see Fig. S10 in the SI). These results demonstrate that the revealed growth fundamental of CsPbI₃ is nearly independent on the physical properties of the underline 2D materials but is universal valid. This unique growth behavior of CsPbI₃ provides an advantage of forming different types of CsPbI₃/2D material heterojunction as demanded, which is highly suitable for novel optoelectronic device applications.

Similar to those CsPbI₃/WSe₂ heterojunction, no change could be observed for Raman scattering of monolayer TMDCs before and after CsPbI₃ growth. PL quenching is also observed for all the CsPbI₃/TMDC heterojunction, especially for the monolayer WS₂ related heterojunction. Slightly peak shift is observed, since trion emission is sensitive to the dielectric environment. These results suggest the CsPbI₃ forms a type-II band alignment with all the monolayer TMDCs. Consequently, the internal electric field induces fast carrier transfer in these heterojunctions, thus reducing the recombination chance of photoexcited electrons and holes.

The optical emission of CsPbI₃/TMDC are further studied at low temperature, as shown in Fig.6 Similar to CsPbI₃/WSe₂ heterostructure, CsPbI₃/WS₂ and CsPbI₃/MoSe₂ also show similar emission behaviors. When compared with pure WS₂ and MoSe₂, a broad and low energy emission peak is always observed as a shoulder to TMDCs emission peak. The extracted peak energy evolution with temperature is shown in Fig.6(e, f) together with the evolution of trion and exciton for pure WS₂ and MoSe₂. For pure WS₂ and MoSe₂, trion and exciton emission shows a linear blueshift with decreasing temperature. In the case of the CsPbI₃ NW/WS₂ heterojunction, a low-energy peak is observable during the whole investigated temperature range. When temperature is below 153K, this low energy peak even outweighs the emission of pure WS₂. The lowest emission energy is around 80 meV lower than the trion emission of WS₂. It is worth mentioning that CsPbI₃/WS₂ heterojunction induced emission energy (1.93 eV) is still larger than that of CsPbI₃. Based on the band alignment between CsPbI₃ and TMDC layers in Fig.6(g) [19], even though the heterostructure is supposed to induce a type-II emission with lower energy, this energy should be smaller than the bandgap of both CsPbI₃ and TMDC. Therefore, the observed lower emission energy peak in WS₂/CsPbI₃ is not related to the type-II band alignment. In terms of CsPbI₃ NW/MoSe₂ heterojunction, a low-energy peak at around 1.48 eV is observed when temperature is below 153 K. This peak energy is 72 meV lower than the trion emission of MoSe₂. At low temperature, the emission of CsPbI₃ and MoS₂ overlap with each other, thus is not shown here.

In all, for all the fabricated CsPbI₃/monolayer TMDC, it forms a type-II band alignment, which leads to a drop of PL emission due to fast charge transfer. At low temperature, band misalignment induced PL emission is not observed, probably because photo induced charges can diffuse in CsPbI₃, which reduce the possibility of forming interlayer exciton at the interface as those in pure TMDC heterostructure [56]. Instead, the fabricated heterostructure change the emission behavior of TMDC layer. Firstly, the dominant emission at room temperature is changed from exciton to trion, leading to a redshift of the PL emission spectrum. Secondly, a broad emission peak with energy around several tens meV lower than the trion emission can be observed, which is not aroused by defects as well as band alignment. Instead, it is supposed that this band alignment induced charge injection from CsPbI₃ to TMDCs layer alter the recombination paths in TMDCs. Consequently, this emission energy and intensity is sensitive to the CsPbI₃.

The determined type-II band alignment together with the ability to re-grown CsPbI₃ on transferred TMDC, suggesting that this method is able to enhance the photodetection performance of TMDC based photodetector. Thus, after metal contact to the monolayer WSe₂ (see optical image in the inset of Fig.7), CsPbI₃ film was regrown on the fabricated WSe₂ devices. The I‒V curves before and after CsPbI₃ growth are compared in Fig.7 with a radiation of a 532 nm laser. The growth of CsPbI₃ significantly enhance the dark current, because the grown CsPbI₃ increase the dark current, which is in agreement with previous WSe₂‒CH₃NH₃PbI₃ heterostructure devices [60]. The photocurrent and on/off ratio are also largely enhanced as well under all the applied voltage range. Specifically, photocurrent increases from 2.38 to 38.59 nA and the on/off ratio of CsPbI₃/WSe₂ increases from 3.14 to 16.37 under voltage of −1 V. The enhancement of optoelectronic performance agrees with our expectation, since the CsPbI₃ shows large light absorption coefficients and the fast charge transfer process can increase the photoinduced carrier density and thus the photocurrent. The photo response characteristics indicate that the hybrid CsPbI₃/WSe₂ possesses a great potential in the application of photodetectors.

4 Conclusions

In summary, we demonstrate direct van der Waals epitaxy of CsPbI₃ NWs and thickness tunable CsPbI₃ nanofilms on WSe₂ without causing any damage to the monolayer WSe₂ substrate. The low adsorption energy of CsI and PbI₂ adatoms on TMDC substrate make it possible to selectively grow CsPbI₃ on TMDC layer instead of the amorphous SiO₂ substrate. Moreover, the temperature dependent desorption rate of the precursor leads to possibility of controlling the morphology of CsPbI₃ from nanofilm to Zigzag/armchair orientated NWs. Both the CsPbI₃ nanofilm and NWs display a distinct inclination towards the formation of a cubic phase structure, establishing an epitaxial relationship with the growth substrate. The revealed van der Waals growth fundamental of CsPbI₃ is valid for all the TMDC layers, including but not limited to CVD grown monolayer/multilayer TMDC as well as h-BN, which pose great feasibility to design CsPbI₃ based van der Waals heterostructure for different application purpose. Further XPS and optical characterization confirms the type-II band alignment between CsPbI₃/monolayer TMDC, thus leading to a fast charge transfer process and the resulted reduction of PL emission intensity. Also, the band alignment induced carrier injection seems alters the recombination paths in TMDC monolayer, leading to the occurrence of a broad PL peak at lower energy side of pure TMDC emission. The WSe₂/CsPbI₃ heterostructure photodetectors exhibit enhanced photocurrent and on/off ratio. Therefore, our finding suggests a new approach to alter the emission as well as optoelectrical performance in TMDC layer.

References

[1] M. Long , P. Wang , H. Fang , W. Hu . Progress, challenges, and opportunities for 2D material based photodetectors. Adv. Funct. Mater., 2019, 29(19): 1803807

[2] N. Huo , G. Konstantatos . Recent progress and future prospects of 2D‐based photodetectors. Adv. Mater., 2018, 30(51): 1801164

[3] S. B. Mitta , M. S. Choi , A. Nipane , F. Ali , C. Kim , J. T. Teherani , J. Hone , W. J. Yoo . Electrical characterization of 2D materials-based field-effect transistors. 2D Mater., 2021, 8(1): 012002

[4] Z. Wang , Y. Yang , B. Hua , Q. Ji . Synthetic two-dimensional electronics for transistor scaling. Front. Phys., 2023, 18(6): 63601

[5] J. F. Gonzalez Marin , D. Unuchek , Z. Sun , C. Y. Cheon , F. Tagarelli , K. Watanabe , T. Taniguchi , A. Kis . Room-temperature electrical control of polarization and emission angle in a cavity-integrated 2D pulsed LED. Nat. Commun., 2022, 13(1): 4884

[6] K. F. Mak , J. Shan . Photonics and optoelectronics of 2D semiconductor transition metal dichalcogenides. Nat. Photonics, 2016, 10(4): 216

[7] S. Wu , S. Buckley , J. R. Schaibley , L. Feng , J. Yan , D. G. Mandrus , F. Hatami , W. Yao , J. Vučković , A. Majumdar , X. Xu . Monolayer semiconductor nanocavity lasers with ultralow thresholds. Nature, 2015, 520(7545): 69

[8] Y. Ye , Z. J. Wong , X. Lu , X. Ni , H. Zhu , X. Chen , Y. Wang , X. Zhang . Monolayer excitonic laser. Nat. Photonics, 2015, 9(11): 733

[9] M. Bernardi , M. Palummo , J. C. Grossman . Extraordinary sunlight absorption and one nanometer thick photovoltaics using two-dimensional monolayer materials. Nano Lett., 2013, 13(8): 3664

[10] C. K. Sumesh . Towards efficient photon management in nanostructured solar cells: Role of 2D layered transition metal dichalcogenide semiconductors. Sol. Energy Mater. Sol. Cells, 2019, 192: 16

[11] D. Wu , Y. Wang , L. Zeng , C. Jia , E. Wu , T. Xu , Z. Shi , Y. Tian , X. Li , Y. H. Tsang . Design of 2D layered PtSe₂ heterojunction for the high-performance, room-temperature, broadband, infrared photodetector. ACS Photonics, 2018, 5(9): 3820

[12] L. Zhang , Z. Zhang , F. Wu , D. Wang , R. Gogna , S. Hou , K. Watanabe , T. Taniguchi , K. Kulkarni , T. Kuo , S. R. Forrest , H. Deng . Twist-angle dependence of moiré excitons in WS₂/MoSe₂ heterobilayers. Nat. Commun., 2020, 11(1): 5888

[13] D. Kozawa , A. Carvalho , I. Verzhbitskiy , F. Giustiniano , Y. Miyauchi , S. Mouri , A. Castro Neto , K. Matsuda , G. Eda . Evidence for fast interlayer energy transfer in MoSe₂/WS₂ heterostructures. Nano Lett., 2016, 16(7): 4087

[14] A. Liu , H. Zhu , S. Bai , Y. Reo , T. Zou , M. G. Kim , Y. Y. Noh . High-performance inorganic metal halide perovskite transistors. Nat. Electron., 2022, 5(2): 78

[15] R. Singh , P. Singh , G. Balasubramanian . Effect of heterostructure engineering on electronic structure and transport properties of two-dimensional halide perovskites. Comput. Mater. Sci., 2021, 200: 110823

[16] A. G. Ricciardulli , S. Yang , J. H. Smet , M. Saliba . Emerging perovskite monolayers. Nat. Mater., 2021, 20(10): 1325

[17] S. S. Lo , T. Mirkovic , C. H. Chuang , C. Burda , G. D. Scholes . Emergent properties resulting from type‐II band alignment in semiconductor nanoheterostructures. Adv. Mater., 2011, 23(2): 180

[18] Y. Liu , H. Li , X. Zheng , X. Cheng , T. Jiang . Giant photoluminescence enhancement in monolayer WS₂ by energy transfer from CsPbBr₃ quantum dots. Opt. Mater. Express, 2017, 7(4): 1327

[19] Q. Fang , Q. Shang , L. Zhao , R. Wang , Z. Zhang , P. Yang , X. Sui , X. Qiu , X. Liu , Q. Zhang , Y. Zhang . Ultrafast charge transfer in perovskite nanowire/2D transition metal dichalcogenide heterostructures. J. Phys. Chem. Lett., 2018, 9(7): 1655

[20] M. Zhong , B. Cui , Z. Mo , Y. Yu , Q. Xia , F. Zhang , Z. Zhou , L. Huang , B. Li , J. Yang . Gate controllable band alignment transition in 2D black-arsenic/WSe₂ heterostructure. Appl. Phys. Rev., 2023, 10(2): 021416

[21] F. Zhao , D. Wang , F. Zhang , B. Cui , Q. Xia , M. Zhong . Gate-controlled photoresponse improvement in b-AsP/WSe₂ heterostructures with type-I band alignment. Appl. Phys. Lett., 2023, 122(15): 151105

[22] L. Yuan , Z. Xu , J. Li , F. Zhang , S. Liu , H. Shi , Q. Xia , M. Zhong . Broad-spectrum and ultrasensitive photodetectors based on GeSe/SnS₂ heterostructures with type-III band alignment. Appl. Phys. Lett., 2023, 122(24): 241106

[23] X. Song , X. Liu , D. Yu , C. Huo , J. Ji , X. Li , S. Zhang , Y. Zou , G. Zhu , Y. Wang , M. Wu , A. Xie , H. Zeng . Boosting two-dimensional MoS₂/CsPbBr₃ photodetectors via enhanced light absorbance and interfacial carrier separation. ACS Appl. Mater. Interfaces, 2018, 10(3): 2801

[24] A. G. Saraswathy Vilasam , P. K. Prasanna , X. Yuan , Z. Azimi , F. Kremer , C. Jagadish , S. Chakraborty , H. H. Tan . Epitaxial growth of GaAs nanowires on synthetic mica by metal–organic chemical vapor deposition. ACS Appl. Mater. Interfaces, 2022, 14(2): 3395

[25] L. Britnell , R. M. Ribeiro , A. Eckmann , R. Jalil , B. D. Belle , A. Mishchenko , Y. J. Kim , R. V. Gorbachev , T. Georgiou , S. V. Morozov , A. N. Grigorenko , A. K. Geim , C. Casiraghi , A. H. C. Neto , K. S. Novoselov . Strong light−matter interactions in heterostructures of atomically thin films. Science, 2013, 340(6138): 1311

[26] Z. Liu , L. Song , S. Zhao , J. Huang , L. Ma , J. Zhang , J. Lou , P. M. Ajayan . Direct growth of graphene/hexagonal boron nitride stacked layers. Nano Lett., 2011, 11(5): 2032

[27] W. Yang , G. Chen , Z. Shi , C. C. Liu , L. Zhang , G. Xie , M. Cheng , D. Wang , R. Yang , D. Shi , K. Watanabe , T. Taniguchi , Y. Yao , Y. Zhang , G. Zhang . Epitaxial growth of single-domain graphene on hexagonal boron nitride. Nat. Mater., 2013, 12(9): 792

[28] M.WangS.K. JangW.J. JangM.KimS.Y. ParkS.W. KimS.J. KahngJ.Y. ChoiR.S. RuoffY.J. SongS.Lee, A platform for large‐scale graphene electronics – CVD growth of single-layer graphene on CVD-grown hexagonal boron nitride, Adv. Mater. 25(19), 2746 (2013)

[29] X. Duan , C. Wang , J. C. Shaw , R. Cheng , Y. Chen , H. Li , X. Wu , Y. Tang , Q. Zhang , A. Pan , J. Jiang , R. Yu , Y. Huang , X. Duan . Lateral epitaxial growth of two-dimensional layered semiconductor heterojunctions. Nat. Nanotechnol., 2014, 9(12): 1024

[30] C. Huang , S. Wu , A. M. Sanchez , J. J. P. Peters , R. Beanland , J. S. Ross , P. Rivera , W. Yao , D. H. Cobden , X. Xu . Lateral heterojunctions within monolayer MoSe₂–WSe₂ semiconductors. Nat. Mater., 2014, 13(12): 1096

[31] M. Chen , R. Chang , X. Yang , C. Lu , S. Zhang , Z. Zhang , J. He , X. Yuan . Van der Waals epitaxy of CsPbBr₃/WSe₂ heterostructure and dynamics study of exciton recombination. J. Phys. D Appl. Phys., 2024, 57(23): 235103

[32] Y. H. Mao , H. Shan , J. R. Wu , Z. J. Li , C. Z. Wu , X. F. Zhai , A. D. Zhao , B. Wang . Observation of pseudogap in SnSe₂ atomic layers grown on graphite. Front. Phys., 2020, 15(4): 43501

[33] Y. Kim , S. S. Cruz , K. Lee , B. O. Alawode , C. Choi , Y. Song , J. M. Johnson , C. Heidelberger , W. Kong , S. Choi , K. Qiao , I. Almansouri , E. A. Fitzgerald , J. Kong , A. M. Kolpak , J. Hwang , J. Kim . Remote epitaxy through graphene enables two-dimensional material-based layer transfer. Nature, 2017, 544(7650): 340

[34] J. Chen , Y. Fu , L. Samad , L. Dang , Y. Zhao , S. Shen , L. Guo , S. Jin . Vapor-phase epitaxial growth of aligned nanowire networks of cesium lead halide perovskites (CsPbX₃, X = Cl, Br, I). Nano Lett., 2017, 17(1): 460

[35] S. Ge , F. Huang , J. He , Z. Xu , Z. Sun , X. Han , C. Wang , L. B. Huang , C. Pan . Bidirectional photoresponse in perovskite‐ZnO heterostructure for fully optical-controlled artificial synapse. Adv. Opt. Mater., 2022, 10(11): 2200409

[36] B. R. Tak , V. Gupta , A. K. Kapoor , Y. H. Chu , R. Singh . Wearable gallium oxide solar-blind photodetectors on muscovite mica having ultrahigh photoresponsivity and detectivity with added high-temperature functionalities. ACS Appl. Electron. Mater., 2019, 1(11): 2463

[37] P. K. Mohseni , A. Behnam , J. D. Wood , C. D. English , J. W. Lyding , E. Pop , X. Li . In_xGa_1−xAs nanowire growth on graphene: van der Waals epitaxy induced phase segregation. Nano Lett., 2013, 13(3): 1153

[38] J. Shin , H. Kim , S. Sundaram , J. Jeong , B. I. Park , C. S. Chang , J. Choi , T. Kim , M. Saravanapavanantham , K. Lu , S. Kim , J. M. Suh , K. S. Kim , M. K. Song , Y. Liu , K. Qiao , J. H. Kim , Y. Kim , J. H. Kang , J. Kim , D. Lee , J. Lee , J. S. Kim , H. E. Lee , H. Yeon , H. S. Kum , S. H. Bae , V. Bulovic , K. J. Yu , K. Lee , K. Chung , Y. J. Hong , A. Ougazzaden , J. Kim . Vertical full-colour micro-LEDs via 2D materials-based layer transfer. Nature, 2023, 614(7946): 81

[39] Y. Wang , Y. Shi , G. Xin , J. Lian , J. Shi . Two-dimensional van der Waals epitaxy kinetics in a three-dimensional perovskite halide. Cryst. Growth Des., 2015, 15(10): 4741

[40] Z. Zhang , P. Chen , X. Yang , Y. Liu , H. Ma , J. Li , B. Zhao , J. Luo , X. Duan , X. Duan . Ultrafast growth of large single crystals of monolayer WS₂ and WSe₂. Natl. Sci. Rev., 2020, 7(4): 737

[41] A. Gurarslan , Y. Yu , L. Su , Y. Yu , F. Suarez , S. Yao , Y. Zhu , M. Ozturk , Y. Zhang , L. Cao . Surface-energy-assisted perfect transfer of centimeter-scale monolayer and few-layer MoS₂ films onto arbitrary substrates. ACS Nano, 2014, 8(11): 11522

[42] X. Yuan , H. Liu , S. Liu , R. Zhang , Y. Wang , J. He , H. H. Tan , C. Jagadish . Thermodynamic properties of metastable wurtzite InP nanosheets. J. Phys. D, 2021, 54(50): 505112

[43] K. Zhang , C. Ding , B. Pan , Z. Wu , A. Marga , L. Zhang , H. Zeng , S. Huang . Visualizing van der Waals epitaxial growth of 2D heterostructures. Adv. Mater., 2021, 33(45): 2105079

[44] S. Mukherjee , N. Nateghi , R. M. Jacobberger , E. Bouthillier , M. De La Mata , J. Arbiol , T. Coenen , D. Cardinal , P. Levesque , P. Desjardins , R. Martel , M. S. Arnold , O. Moutanabbir . Growth and luminescence of polytypic InP on epitaxial graphene. Adv. Funct. Mater., 2018, 28(8): 1705592

[45] C. Sheng , Y. Bu , Y. Li , L. Su , Y. Yu , D. Cao , J. Zhou , X. Chen , W. Lu , H. Shu . Phase-controllable growth of air-stable SnS nanostructures for high-performance photodetectors with ultralow dark current. ACS Appl. Mater. Interfaces, 2023, 15(11): 14704

[46] H.RyuH.ParkJ.H. KimF.RenJ.KimG.H. LeeS.J. Pearton, Two-dimensional material templates for van der Waals epitaxy, remote epitaxy, and intercalation growth, Appl. Phys. Rev. 9(3), 031305 (2022)

[47] D. Trots , S. Myagkota . High-temperature structural evolution of caesium and rubidium triiodoplumbates. J. Phys. Chem. Solids, 2008, 69(10): 2520

[48] J. Liang , X. Han , J. H. Yang , B. Zhang , Q. Fang , J. Zhang , Q. Ai , M. M. Ogle , T. Terlier , A. A. Martí , J. Lou . Defect-engineering-enabled high-efficiency all-inorganic perovskite solar cells. Adv. Mater., 2019, 31(51): 1903448

[49] Y. Cheng , P. Tang , P. Liang , X. Liu , D. Cao , X. Chen , H. Shu . Sulfur-driven transition from vertical to lateral growth of 2D SnS–SnS₂ heterostructures and their band alignments. J. Phys. Chem. C, 2020, 124(50): 27820

[50] X. Liu , D. Cao , Y. Yao , P. Tang , M. Zhang , X. Chen , H. Shu . Heteroepitaxial growth and interface band alignment in a large-mismatch CsPbI₃/GaN heterojunction. J. Mater. Chem. C, 2022, 10(6): 1984

[51] H. Yuan , J. Su , S. Zhang , J. Di , Z. Lin , J. Zhang , J. Zhang , J. Chang , Y. Hao . Interfacial transport modulation by intrinsic potential difference of Janus TMDs based on CsPbI₃/J-TMDs heterojunctions. iScience, 2022, 25(3): 103872

[52] D. Guo , D. Bartesaghi , H. Wei , E. M. Hutter , J. Huang , T. J. Savenije . Photoluminescence from radiative surface states and excitons in methylammonium lead bromide perovskites. J. Phys. Chem. Lett., 2017, 8(17): 4258

[53] A. Asaithambi , N. Kazemi Tofighi , N. Curreli , M. De Franco , A. Patra , N. Petrini , D. Baranov , L. Manna , F. D. Stasio , I. Kriegel . Generation of free carriers in MoSe₂ monolayers via energy transfer from CsPbBr₃ nanocrystals. Adv. Opt. Mater., 2022, 10(14): 2200638

[54] X. Yang , S. Zhang , Z. Zhang , J. Lin , X. Liu , Z. Huang , L. Zhang , W. Luo , J. He , X. Yuan . Controlled fabrication of CsPbI₂Br/transition metal dichalcogenide van der Waals heterostructure with fast carrier transfer process and interlayer exciton formation. Physica E, 2023, 153: 115788

[55] M. Drüppel , T. Deilmann , P. Krüger , M. Rohlfing . Diversity of trion states and substrate effects in the optical properties of an MoS₂ monolayer. Nat. Commun., 2017, 8(1): 2117

[56] T. Ye , J. Li , D. Li . Charge-accumulation effect in transition metal dichalcogenide heterobilayers. Small, 2019, 15(42): 1902424

[57] E. M. Alexeev , D. A. Ruiz-Tijerina , M. Danovich , M. J. Hamer , D. J. Terry , P. K. Nayak , S. Ahn , S. Pak , J. Lee , J. I. Sohn , M. R. Molas , M. Koperski , K. Watanabe , T. Taniguchi , K. S. Novoselov , R. V. Gorbachev , H. S. Shin , V. I. Fal’ko , A. I. Tartakovskii . Resonantly hybridized excitons in moiré superlattices in van der Waals heterostructures. Nature, 2019, 567: 81

[58] K. Ludwiczak , A. K. Da̧browska , J. Binder , M. Tokarczyk , J. Iwański , B. Kurowska , J. Turczyński , G. Kowalski , R. Bożek , R. Stȩpniewski , W. Pacuski , A. Wysmołek . Heteroepitaxial growth of high optical quality, wafer-scale van der Waals heterostrucutres. ACS Appl. Mater. Interfaces, 2021, 13(40): 47904

[59] J. Li , X. Yuan , P. Jing , J. Li , M. Wei , J. Hua , J. Zhao , L. Tian . Temperature-dependent photoluminescence of inorganic perovskite nanocrystal films. RSC Advances, 2016, 6(82): 78311

[60] J. Lu , A. Carvalho , H. Liu , S. X. Lim , A. H. Castro Neto , C. H. Sow . Hybrid bilayer WSe₂–CH₃NH₃ PbI₃ organolead halide perovskite as a high-performance photodetector. Angew. Chem., 2016, 128(39): 12124

RIGHTS & PERMISSIONS

Higher Education Press