Van der Waals epitaxy of CsPbBr₃/WSe₂ heterostructure and dynamics study of exciton recombination

In recent years, with the advances in the growth and optical characterization of 2D materials, 2D materials are gradually showing promising applications in both photonics and optoelectronics [1–4]. Monolayer (ML) transition metal dichalcogenides (TMDCs) are typical 2D layered materials that are widely used in photodetectors [5, 6], field effect transistors [7, 8], and light-emitting devices [9, 10] because of their direct bandgap and excellent optoelectronic properties. However, the development of high-performance optoelectronic devices is limited by the intrinsic properties of TMDC atomic thin layers [11]. Hence, researchers have endeavored to overcome this problem by integrating TMDCs with other excellent materials to increase the light absorption and photo detectivity [12–14]. For instance, integration of TMDC with two-dimensional perovskites materials to form a van der Waals heterostructure allow the formation of interlayer exciton, showing another approach for the generation and control of interlayer exciton [15, 16].

Inorganic lead halide perovskites have attracted great attention thanks to their long carrier lifetime, long diffusion length, and excellent photoluminescence quantum efficiency [17–20]. Among them, cesium lead halide (CsPbBr₃) is a typical inorganic perovskite, which is often used for integration with TMDCs to improve the performance of optoelectronic devices thanks to its excellent light-absorbing properties and chemical stability [21–23]. For example, Asaithambi et al reported an order of magnitude increase in the photocurrent of CsPbBr₃/MoSe₂ heterostructures compared to pristine MoSe₂ [24]. Fang et al constructed a CsPbBr₃/TMDC vertical heterostructure that exhibited ultrafast carrier transfer [25]. However, there are few studies on the properties of CsPbBr₃/TMDC heterostructures at low temperatures. Meanwhile, the construction of this type of heterostructures is mostly done by transferring process, which has limitations in terms of interface quality and large-scale production [26–28]. Therefore, a more superior method of constructing heterostructures is needed. Instead, van der Waals epitaxy can take advantage of the absence of dangling bonds on the TMDC surface to design and prepare a variety of van der Waals heterostructures with high interfacial quality, which is highly advantageous in constructing heterostructures.

Here, we successfully demonstrate van der Waals epitaxy of CsPbBr₃ nanowires (NWs), nanoplates (NPs) and nanocones (NCs) on ML WSe₂ by chemical vapor deposition (CVD). Based on the CsPbBr₃ NPs/ML WSe₂, the carrier dynamics of the heterostructure was investigated by temperature-dependent photoluminescence (PL) and time-resolved photoluminescence (TRPL), and it was demonstrated that the as-grown heterostructure constructs a type-II band alignment with efficient and ultrafast carrier transfer. Temperature-dependent PL spectra show that higher energy emission peaks near the characteristic peak of CsPbBr₃ and lower energy emission peaks near ML WSe₂ appear in this heterostructure at low temperatures. Meanwhile, the lattice dynamics of the heterostructure was studied using temperature-dependent Raman spectroscopy. Our results lay the foundation for revealing the photophysical processes and light-matter interactions of the CsPbBr₃/WSe₂ heterostructure, which is crucial for the development of high-performance optoelectronic devices.

2.1. The WSe₂ was synthesized by using physical vapor deposition

2.2. CsPbBr₃/WSe₂ heterostructure was synthesized by chemical vapor deposition

2.3. Morphology and optical characterization

Figure 1 shows the morphology of CsPbBr₃ on ML WSe₂ under different growth conditions (larger magnification SEM see figure S1 in the supporting information). Growth temperature is an important factor affecting the morphology of CsPbBr₃. At 240 °C, rectangular CsPbBr₃ NPs are formed on ML WSe₂ while nano rocks are observed on the rest of the amorphous SiO₂ substrate (figure 1(b)). The grown CsPbBr₃ NPs show clear crystal orientation preference with the WSe₂, as illustrated in figure 1(h), where NPs edges are always parallel to the zigzag direction, which has also been observed in previous reports [29]. AFM results in figures 1(k) and (l) show that the NPs have a very smooth surface with a variation of 0.67 nm. The thickness varies between 15 nm–40 nm and the length generally does not exceed 2 μm. Slightly reducing the growth temperature to 230 °C, the CsPbBr₃ NPs become smaller and tend to accumulate at the edge of WSe₂ (figure 1(c)), since lower temperatures do not provide enough energy for nucleation and the existence of dangling bonds at the WSe₂ edge reduces the nucleation barrier. In contrast, at a higher growth temperature of 250 °C, CsPbBr₃ tends to grow out of plane on WSe₂, resulting in the formation of NCs (figure 1(a)), as illustrated in figure 1(i). The temperature of the source is another important parameter that determines the growth of CsPbBr₃. The source temperature determines the concentration of CsBr and PbBr₂ and the relative content ratio in the vapor, which affects the nucleation and growth of CsPbBr₃ on WSe₂. Decreasing the source temperature from 700 °C to 660 °C, even though a lower concentration of precursor is suspected, the nucleation density of CsPbBr₃ on WSe₂ is surprisingly high. Moreover, there exists a strong source temperature dependent morphology evolution. At 700 °C, only a few NWs and amorphous small NPs are formed. (Figure 1(d)). Slightly reducing the temperature to 680 °C, NWs with a high aspect ratio are successfully grown on WSe₂ (figure 1(e)). As shown in figure 1(g), the CsPbBr₃ NWs are not parallel to either the zigzag or armchair direction of WSe₂, but mostly present an angle of 14° to the zigzag direction. According to our calculations, this is because the NWs have a lower lattice mismatch in the direction 14° from the zigzag direction. In contrast, the lattice mismatch is slightly higher in the zigzag and armchair directions (calculations shown in figure S2). The NW thickness is around 58.8 nm (see figure 1(j)). However, when the temperature is 660 °C, no NWs are formed. Instead, a high density of irregular shaped and nanoplate shaped CsPbBr₃ is formed (see figure 1(f)). This temperature sensitive growth trend of CsPbBr₃ is not similar to PbI₂ based perovskite growth [30], CsPbBr₃ growth behavior is much more difficult and different. First, it is difficult to grow CsPbBr₃ selectively on WSe₂ by changing the growth parameters, including temperature, pressure and flow rate, which is ascribed to the similar adsorption energy of PbBr₂ on both TMDC and SiO₂ substrates. Secondly, the morphology of CsPbBr₃ on WSe₂ is extremely sensitive to the growth parameters, making it hard to control the morphology of as-grown CsPbBr₃.

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The optical properties of the heterostructure are shown in figure 2. Regarding the CsPbBr₃/ML WSe₂ heterostructure, some researchers consider that this heterostructure forms a type I band alignment [31, 32]. However, the applied CsPbBr₃ quantum dot solutions or the non-polar solutions may lead to decoupling of the WSe₂ and enhance the PL intensity, making it difficult to prove that the enhance emission is a result of type I band alignment. Meanwhile, there are also reports claiming a type-II band alignment [25]. For the type-II band alignment, as illustrated in figure 2(a), the photogenerated electrons and holes are physically separated into different layers due to the interlayer electric field during the photo-excitation generation process, which reduces the photon emission intensity and leads to the possibility of interlayer exciton formation. Photon emission intensity mapping of the CsPbBr₃/ML WSe₂ heterostructure at 1.592 eV and 1.587 eV in figures 2(b) and (c) clearly shows that the uniform and strong emission from the exciton of WSe₂ but weak emission intensity at the heterostructure region, regardless of the morphology of CsPbBr₃. This emission trend is in agreement with the type-II band alignment. Since the emission properties of CsPbBr₃/WSe₂ are nearly independent of morphologies, all the following discussions are based on CsPbBr₃ NPs. Quantitative emission comparison of isolated WSe₂, isolated CsPbBr₃ and CsPbBr₃/WSe₂ heterostructure is shown in figure 2(d). As-grown crystal on SiO₂ shows a strong emission at 2.259 eV, demonstrating the successful formation of CsPbBr₃. This emission intensity is dropped over 10 times when forming a heterostructure. Meanwhile, the emission from WSe₂ is reduced by ∼59%. The slightly blue shift of the PL peak position of CsPbBr₃ on WSe₂ compared to the CsPbBr₃ crystals on SiO₂ may be due to the different thickness of the CsPbBr₃ NPs compared to the crystals on SiO₂ [33]. For CsPbBr₃ NWs partly on WSe₂ and partly on SiO₂, as shown in figure S3, the emission peak of CsPbBr₃ is then the same and the emission intensity on WSe₂ is about 10 times lower than that of CsPbBr₃ on SiO₂. In addition, this charge transfer process induced emission is further investigated using TRPL spectroscopy, as shown in figure 2(e). Carrier decay for isolated CsPbBr₃ presents a mono-exponential decay with an extracted PL lifetime of 3.6 ns. This decay behavior agrees with synthetic CsPbBr₃ crystals [34–36], indicating a high crystalline quality. In contrast, after the formation of heterostructure, PL intensity presents a biexponential decay trend. The fast PL lifetime significantly reduces to 58 ps, which is explained as a result of the fast charge transfer from CsPbBr₃ to WSe₂. For the trion/exciton emission of WSe₂, it always presents a biexponential decay trend, where the fast decay process is due to the many-body effect and the slower decay process is attributed to the exciton radiative complexation of WSe₂ [37–39]. The fast decay lifetime slightly drops from 46 to 38 ps, agreeing with the slight intensity reduction. The decrease in PL intensity and PL lifetime proves that WSe₂ and CsPbBr₃ form a type-II heterostructure and show an efficient and ultrafast carrier transfer process.

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The phonon scattering of the heterostructure is examined by temperature-dependent Raman spectra. Figures 3(a) and (b) compare the Raman spectra of ML WSe₂ and CsPbBr₃/ML WSe₂ heterostructure from 93 to 298 K. The temperature-dependent Raman scattering of WSe₂ after heterostructure formation is similar to that of pristine WSe₂. The intensity of the WSe₂ peak (249 cm⁻¹) decreases during cooling down and the peak position is basically unchanged (see figure 3(d)), which indicates that high temperature growth of CsPbBr₃ does not cause damage to the underlying fragile WSe₂. In comparison, the Raman scattering of CsPbBr₃ is more complex. The Raman scattering of CsPbBr₃ is diffuse at room temperature, but increases rapidly and becomes clearer during cooling down. At 213 K, the Raman scattering intensity of CsPbBr₃ outweighs that of WSe₂. At 93 K, the Raman scattering of CsPbBr₃ is so strong that the vibration mode of WSe₂ becomes nearly unobservable. At this temperature, CsPbBr₃ is an orthorhombic phase and all its Raman modes are active, so the heterostructure shows strong Raman characteristic peaks at 48 cm⁻¹, 65 cm⁻¹, 73 cm⁻¹, 123 cm⁻¹, 152 cm⁻¹ and 306 cm⁻¹ [40] where the peaks at 123 cm⁻¹ and 152 cm⁻¹ are ascribed to the stretching vibrations of Pb–Br, which belong to the transverse optical (TO) phonon and longitudinal optical (LO) phonon, respectively [41, 42]. It is interesting to see that the second order LO vibration mode at 306 cm⁻¹ depicts a similar intensity as LO. For the Raman at low wave numbers of 40–90 cm⁻¹ as shown in the enlarged view in figure 3(c), the vibration modes at 65 cm⁻¹ and 73 cm⁻¹ are related to the vibrations of the [PbBr₆]⁴⁻ octahedron [43]. As the temperature increases to room temperature, the diffuse Raman peaks are consistent with the Raman signature of the cubic phase CsPbBr₃ [44]. Meanwhile, the CsPbBr₃ on SiO₂ was also characterized by XRD (see figure 3(d)). Five characteristic diffraction peaks at 15.56°, 21.99°, 30.70°, 38.32° and 44.60° are clearly observed, which can be assigned as the (100), (110), (200), (211) and (220) crystal planes of the cubic phase CsPbBr₃ [45, 46]. The growth of cubic CsPbBr₃ is attributed to the high growth temperature [46]. This suggests that a temperature-dependent phase transition may have occurred in CsPbBr₃ during the cooling process. This phase transition has been observed in single-crystalline CsPbX₃ (X = Cl, Br, I) NPs and CsPbX₃ quantum dots [46, 47]. Figures 3(e) and (f) present the variation of the peak position and full-widths at half-maximum (FWHM) of the phonon modes with temperature. The peaks of the LO and 2LO phonon modes are shifted linearly to the high frequency by about 4.5 cm⁻¹ and 1.5 cm⁻¹, respectively, with the change of temperature from 93 K to 298 K. Meanwhile, the FWHM of LO and 2LO increases linearly with temperature.

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Further temperature-dependent PL experiments on the CsPbBr₃/ML WSe₂ heterostructure are carried out in the range of 298 K–93 K. The intensity of the emission peaks of both CsPbBr₃ and WSe₂ gradually increases during the cooling process (figure S4). The normalized spectra are shown in figures 4(a)–(c). The emission peaks from CsPbBr₃ and WSe₂ undergo a blue shift of 8 meV and 70 meV, respectively, and their FWHMs gradually become narrower during cooling down. The surprising phenomenon of the blue shift of CsPbBr₃ is consistent for all CsPbBr₃ nanocrystals (see figure S5) grown in this work but is in contradiction to the literature reports [48]. Considering the metastable cubic phase as well as smaller emission energy at room temperature, the slightly blue shift of emission during cooling down process may be related to the structural change. For the CsPbBr₃ emission, a high-energy peak appears as a shoulder at 2.316 eV when the temperature is below 133 K. During the power-dependent PL at 93 K in figure S6, this shoulder peak remains essentially constant with increasing excitation power, indicating that the peak is not caused by defects. We speculate that this peak may be related to the phase transition of CsPbBr₃ during the cooling process, which is similar to the phenomenon occurring in the same type of perovskite [49, 50]. The occurrence of this peak is also in agreement with the above Raman analysis. In addition, WSe₂ shows a low-energy peak of about 1.586 eV near the exciton emission peak at temperatures lower than 193 K, which is about 38 meV smaller than the trion emission of the ML WSe₂ alone (figure S7), the quality and the thermal treatment for the WSe₂ and the ML WSe₂/CsPbBr₃ heterojunction area is exactly the same. Continuing to cool down to 93 K, the FWHM of this low-energy peak gradually narrows and exceeds the exciton emission intensity (figure 4(b)). Power-dependent PL in figures 4(d) and (e) shows that the peak positions as well as the relative intensity of this low energy peak to exciton emission is nearly independent of the excitation power for over one orders of magnitude. Thus, the emergence of the low energy peak is not likely due to defects. At the same time, according to previous reports, interlayer exciton energy exhibits a power-dependent blue shifts [51]. However, this heterostructure induced low energy emission peak is nearly unshifted during the investigated excitation power range. In addition, after removing CsPbBr₃ from the heterostructure with deionized water, the low-energy peaks near WSe₂ no longer appeared (figure S8), again confirming that the low-energy emission peak is not due to damage or doping to the WSe₂ during the CsPbBr₃ growth but induced by the heterostructure itself.

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In addition to ML WSe₂, the optical performance of CsPbBr₃/double layer (DL) WSe₂ is compared in figure 4(f). After heterostructure formation, the emission of DL WSe₂ is nearly unaffected, indicating the effect of charge transfer on photon emission in indirect bandgap WSe₂ is not obvious. Reducing the temperature to 93 K, the emission properties between DL WSe₂ and the related heterostructure are still similar, showing similar intensity and peak shifts in the K-K and K-F bands. The above-mentioned low-energy emission peak close to the trion of ML WSe₂ does not appear in the CsPbBr₃/DL WSe₂ heterostructure. Therefore, it further confirms that the formation of CsPbBr₃/ML WSe₂ would alter the carrier recombination paths of pristine WSe₂ and the heterojunction can result in a new emission peak at low temperatures.

Our results demonstrate that van der Waals epitaxy is a powerful approach to construct high quality perovskite and TMDC heterostructure. Using CsPbBr₃/WSe₂ as an example, we show the type-II band alignment, efficient charge transfer process and heterostructure related low energy emission at low temperature. These results demonstrate that this hybrid heterostructure with type-II band alignment shows potential application in self-powered photodetection applications.

In summary, we demonstrate van der Waals epitaxy of CsPbBr₃ on ML WSe₂ with different morphologies, including NPs, NWs, and NCs. The growth of CsPbBr₃ is very sensitive to the growth parameters. Temperature is the main factor driving the change of CsPbBr₃ morphology. We confirm that the CsPbBr₃/ML WSe₂ heterostructure forms a type-II band alignment and exhibits ultrafast carrier transfer, leading to a reduction of emission intensity and an acceleration of carrier decay rate. Heterostructure leads to the appearance of a broad emission peak at the low-energy side of the WSe₂ emission below 193 K, which is ascribed to the band alignment of the CsPbBr₃/ML WSe₂ heterostructure. Meanwhile, a high-energy peak appears near the CsPbBr₃ emission below 133 K, which may be related to the phase transition of CsPbBr₃ from the cubic to the orthorhombic phase. In addition, the lattice dynamics of the heterostructure was further investigated by temperature-dependent Raman spectroscopy. Our findings and results demonstrate that van der Waals construction of CsPbBr₃/WSe₂ heterostructure with type-II band alignment is of potential application in fabricating of self-powered photodetector with superior performance.

This work was supported by the National Natural Science Foundation of China (Grant Nos. 61974166 and 62274184) and the Hunan Provincial Natural Science Foundation of China (Grant No. 2021JJ20080). We thank Prof Lin Zhang for positive discussions and support during the XRD measurements.

The data that support the findings of this study are available upon reasonable request from the authors.