Mass production of self-passivated perovskite microlaser particles by solution-phase processing for gas sensors

Microlaser particles (MLPs) represent the compact laser sources where the active materials serve as the optical gain and the cavity resonator simultaneously.^1–3 These MLPs greatly lower the complexity of a laser and feature a small mode volume or device size. Therefore, the MLPs hold promise for developing plenty of cutting-edge applications, such as on-chip coherent light sources, ultrasensitive sensors, and bio-imaging.^4–6 In particular, the laser-emission-based sensing or imaging offers distinctive benefits in terms of narrow linewidth, strong light–matter interactions, and signal amplification, where the narrow linewidth of the laser emission further renders the multiplexing capability.^7–9 Consequently, tremendous research efforts have been devoted to developing MLPs based on the traditional III–V and II–VI semiconductors.^10,11 However, the fabrication of the III–V and II–VI semiconductors enabled MLPs to typically rely on harsh conditions, such as high temperature, vacuum or inert atmosphere, and enclosed space,^11–13 which may not be feasible for large-scale industrial production.

Recently, the inorganic metal halide perovskites (IMHPs) are emerging as a new star in the semiconductor family by virtue of the high emission efficiency, large absorption coefficient, and superior optical gain performance.^14–16 Compared to the organic–inorganic hybrid perovskites, the IMHPs show much enhanced stability against moisture, thermal, and chemical disturbances, allowing a long-term device operation.^17–20 Since the pioneering work on IMHP lasing in the year of 2015,^21–23 impressive progress has been made in this field. Until now, a number of IMHP-based MLPs have been demonstrated including Fabry–Pérot-based nanowires,^22,24 plasmonic nanolasers,^25,26 and whispering-gallery mode (WGM) microplates.^21,24 In particular, the WGM-MLPs are highly desirable coherent light sources for sensing and imaging, thanks to their high-quality (Q) factor and miniaturized size.²⁷ Regarding the synthesis of WGM-MLPs, the most widely adopted technique is the chemical vapor deposition (CVD).^15,21 This method enables the fabrication of high-quality single-crystals of IMHPs, which show favorable lasing performance.²¹ Nevertheless, the CVD technique requires an enclosed space and does not allow for scalable synthesis of MLPs. Thanks to the intrinsic solution processibility of IMHPs, which is the most attractive character that motivates its utilization in optoelectronic devices, the IMHPs can also be fabricated by the solution-phase course.^28,29 However, suffering from the low solubility of IMHP precursors, the reliable and scalable fabrication of IMHP-MLPs by the solution-phase approach remains a daunting challenge.²⁸ Moreover, the optical/lasing performance of these solution-processed MLPs is typically inferior to those of the CVD counterparts, and the understanding of the growth mechanism by solution processes lags far behind compared to that by the CVD route.^15,21 Therefore, in order to propel the application of IMHP-MLPs, it is imperative to develop a mass-productive solution-processing method toward high-performance IMHP-MLPs, similar to the case in IMHP-based solar cells and light-emitting diodes (LEDs).³⁰ 

In this work, we develop a novel solution-phase technique to realize the CsPbBr₃ MLPs in a scalable manner. We adopt a spin-coating-based two-step growth and regulate the dynamic process of crystallization, where the underlying mechanisms are interrogated on basis of microscopic characterization and crystal growth theory. Importantly, the solution-processed CsPbBr₃ microcrystals exhibit much stronger photoluminescence (PL) than the ones prepared by the CVD method, which is attributed to the low carrier trap density by the formation of a self-passivated and bromine-rich surface. These CsPbBr₃ microcrystals with inverted pyramid morphology are demonstrated to support the WGM lasing, featuring low pump threshold and high Q-factor. Moreover, the technologically important single-mode lasing is achieved from the sub-5 μm-sized MLPs, thanks to their superior optical property. These results represent a significant step toward mass-productive and high-performance MLPs, which could find a variety of applications in optoelectronics, as exemplified by the gas sensors herein.

In order to address the poor solubility issue, PbBr₂ was dissolved in hydrobromic acid rather than the widely employed dimethylformamide (DMF) during CsPbBr₃ powder preparation (see Sec. IV and Fig. S1 for details). The prepared CsPbBr₃ can be well-dissolved in DMSO, which is beneficial for the crystal growth in the next step (Fig. S1). Spin-coating is a facile and inexpensive method to prepare epitaxial films,³¹ especially suitable for mass production, as is adopted here. The 0.5M CsPbBr₃ solution was dispensed onto a glass substrate, followed by spin-coating at 3000 r/m. After that, the sample was heated on a 60 °C plate. Figure 1(a) shows the optical absorption and PL spectra of the CsPbBr₃ sample. The absorption edge (red line) is estimated to be at ∼520 nm, and the full width at half maximum (FWHM) of the PL spectrum is ∼20 nm. To further investigate the in-plane order of these crystals, optical microscopy and scanning electron microscope (SEM) measurements were carried out. It shows that the CsPbBr₃ crystals were deposited as islands on the substrate [Figs. 1(b) and 1(c)]. The distributed large crystals possess a rectangular shape with a side length of about 15 µm. This is in line with the SEM results shown in Figs. 1(d) and 1(e). The enlarged image evidenced the inverted pyramid structure of the CsPbBr₃ crystals. This morphology is clearly seen in the side view of Fig. 1(f). The Raman spectrum of the CsPbBr₃ crystal is shown in Fig. S2. The peak at 74 cm⁻¹ is attributed to Pb–Br stretching modes of the PbBr₆ octahedron, the peaks at 128 and 310 cm⁻¹ represent the head-to-head Cs motion coupled with proximal Br face expansion, and a second-order longitudinal optical mode is associated with the vibration of octahedron.³² Figure 2(a) shows the XRD peaks of the sample, which can be assigned to the (020), (101), (121), (202), (040), and (242) crystal faces of CsPbBr₃. The planes equivalent to {002} and {110} tend to form hexahedron, while introducing {112} planes promotes the formation of orthogonal octahedron and dodecahedron crystals.³³ In our experiment, the inverted pyramid-shaped crystals wrapped by {002}, {110}, and {112} planes were obtained by regulating the dynamic process of crystallization. Figure 2(b) shows the element distribution of the (002) cross section in the crystal. Cesium ions are distributed between the [PbBr₆]⁴⁻ octahedral frameworks. The upper slope is the {112} crystal plane family, and the (112), (11-2), (1–12), and (−112) planes together form the slope of the pyramid. For the XRD characterization, a plenty of the CsPbBr₃ microcrystals are interrogated simultaneously by virtue of the large detection diameter (>3 mm). We also measured different positions of the sample and consistent results were obtained, indicating the reproducibility of the XRD result.

According to the above SEM and XRD results and the crystallization characteristics of the spin-coated sample, we proposed the perovskite growth mechanism as illustrated in Figs. 2(c) and 2(d). The solvent in the spin-coated layer evaporated when the sample was heated on a 60 °C plate. Accompanied by the loss of the solvent, the solution reached supersaturation and nucleation occurred [step 1 in Fig. 2(c)]. Different from the nucleation in the bulk solution (homogeneous nucleation) that would occur in conventional solvent evaporation methods, it can be assumed that the spin-coated thin layer promotes CsPbBr₃ nucleation on the glass substrate.³⁴ As the evaporation process proceeded, the solvent on the substrate was further reduced and resulted in the formation of many small droplets [step 2 in Fig. 2(c)]. For a single droplet, the rapid solvent evaporation in the surface region and the growth of crystal on the substrate induce a diminishing concentration ranging from the surface to the center bottom of the droplet. The concentration gradient serves as the driving force for solute diffusion. The solute near the growth front of the crystal located at a droplet center can reach the crystal face for crystallization. In the boundary region of the droplet, however, the unbalanced solute diffusion and solvent evaporation resulted in nucleation and crystallization during the droplet shrink. Moreover, the dramatically faster solvent evaporation would trap the crystallization at an earlier stage, thus forming small crystals in the boundary region. With further solvent removal, the volume of the droplet was reduced to the critical value at which most solute could contribute to the growth of the crystal at a droplet center [step 3 in Fig. 2(c)]. This is evidenced by the indistinguishable small crystal region around the big crystal, as shown in Fig. 1(b). The crystal at the droplet center continues to grow until the complete removal of the solvent. The formation of the inverted pyramid structure in this process could be mainly attributed to the concentration gradient, the temperature gradient, and the thin spin-coated layer. Due to the diminishing concentration from the surface to the center bottom of the droplet, the growth rate at the edge of the CsPbBr₃ crystal is faster. Moreover, heating by the hot plate, the temperature of the droplet gradually decreases from the bottom to the top, forming a longitudinal temperature gradient. The solubility of the CsPbBr₃ crystals decreases as the temperature decreases, leading to the longitudinal growth of the crystals. Accordingly, the synergistic effects of the concentration gradient and the temperature gradient contribute to the inverted pyramid morphology of the CsPbBr₃ crystals. In addition, since the spin-coated layer is very thin, the growth of the crystal center cannot be supplemented by the solute in the late stage. In this case, the growth at the edge of the CsPbBr₃ crystal promotes inverted pyramid structure formation.

Figures 3(a) and 3(b) show the photographs of CsPbBr₃ microcrystals fabricated by the solution-phase process under room light and UV light irradiation, respectively. As a comparison, the photographs of the CVD-fabricated CsPbBr₃ microcrystals are also displayed in Figs. 3(c) and 3(d), where the same UV light irradiation was adopted. Surprisingly, it is found that the sample fabricated by the solution method exhibits much better light-emission properties compared to the CVD one as manifested by the more intense green color under UV irradiation. This phenomenon is further confirmed by quantitative PL and PL lifetime measurements [Figs. 3(e) and 3(f)], where much higher PL intensity and longer PL lifetime of solution-prepared microcrystals are observed than those fabricated by the CVD under the same excitation condition. In particular, the average lifetime (∼7.12 ns) of the CsPbBr₃ microcrystals prepared by the solution method is in sharp contrast to that of the sample prepared by the CVD method (∼2.18 ns). The much stronger PL intensity and longer lifetime indicate the less carrier trapping defects in our solution-processed sample. The PLQY of the two samples are also measured using an integrating sphere system, where the solution-phase and CVD grown samples exhibit PLQY of ∼58% and ∼11%, respectively. To quantitatively derive the trap density, the intensity-dependent PL measurements were performed.^35,36 Upon the photoexcitation, the carrier density n_c(t) can be described by the following differential equation:³⁶ 

where a is the carrier capture coefficient, n_tp(t) is the trap state density, and τ is the lifetime of carriers. Accordingly, the carrier density n_c(0) immediately after the photoexcitation is given by

where I_PL represents the light-emission intensity and k is a constant. See the detailed derivation in the supplementary material. By fitting the experimental results based on Eq. (3) [Fig. 3(g)], the average carrier trapping density is derived to be ∼4.0 × 10¹⁶ cm⁻³ for the solution-processed sample, whereas the trap density for the CVD sample is much higher (∼6 × 10¹⁷ cm⁻³). To shed light on the superior optical properties of the CsPbBr₃ microcrystals by the solution method, the surface composition is interrogated by the surface-sensitive x-ray photoelectron spectroscopy (XPS). Figures 3(h) and 3(i) display the XPS spectra of Pb 4f and Cs 4d, Br 3d for the solution-processed CsPbBr₃ crystals. In contrast to the CVD sample whose surface always lacks bromine,^18,37 it is found that the solution-processed CsPbBr₃ microcrystals exhibit a Br-rich surface with the Cs:Pb:Br ratios of ∼1:1.1:3.4. The Br-rich surface is known to be able to passivate the surface defects and to reduce the nonradiative recombination.¹⁸ Moreover, the much higher composition ratio of Pb and Br than its stoichiometry of 1:1:3 indicates the presence of a PbBr_x layer with a large bandgap (∼4 eV) on the CsPbBr₃ surface. This PbBr_x-rich layer could effectively passivate and protect the CsPbBr₃ microcrystals and render the high PL performance. Similar phenomenon has been observed in the solution-synthesized CsPbBr₃ nanocrystals.³⁸ 

To explore the tunability of the solution-processed CsPbBr₃ microcrystals, we investigated the influence of spin-coating speed, solution concentration, and the annealing temperature on perovskite crystallization. A detailed discussion is shown in Fig. S3. With an increase in the spin-coating speed, the crystal size decreased from a dozen micrometers to about 1 μm. The lateral size distribution of the MLPs is presented in Fig. S3. It is found that there is a wide distribution of MPLs. Nevertheless, the dominant size range can be controlled by the spin-coating speed. For annealing temperature control, the crystals annealed at 80 °C are smaller than those obtained at 60 °C, and the crystals tend to form a dense thin film due to unbalanced crystallization between solute diffusion and solvent evaporation at higher temperature. Therefore, we used a 0.5M precursor solution, a spin-coating speed of 3000 r/min, and an annealing temperature of 60 °C to prepare the CsPbBr₃ crystals with the inverted pyramid structure and size of ∼20 μm.

Inspired by the superior PL property and the regular shape, which may naturally serve as the cavity resonator, we move forward to explore the lasing characteristics of the solution-processed sample. A home-built micro-PL system was employed, which allows us to investigate the photo-response from the single CsPbBr₃ microcrystal. A femtosecond amplified laser was used as the pump source (see the detailed description of the setup in Sec. IV). Figure 4(a) shows the representative pump fluence-dependent PL spectra from the sample. Under low pump fluences, the PL spectra show the relatively broad spontaneous emission (see the 2D plot of the PL spectra as a function of pump fluence in Fig. S4). With the gradual increase in excitation fluence, the evenly spaced sharp spikes appear, indicating the achievement of lasing action.^7,37 The variation in the PL linewidth below and above the threshold is presented in Fig. S5. The fluorescent image above the pump threshold is displayed in the inset of Fig. 4(a). The bright interference pattern along the periphery of the CsPbBr₃ microcrystal agrees with the WGM lasing behavior since the square-shaped boundary could confine the light by total internal reflection.³⁷ In the case of the WGM cavity, the free spectral range (FSR) is described by²⁶ 

where λ is the lasing wavelength and L is the width of the microcrystal or the cavity length. Accordingly, the group refractive index, n_g, is derived to be ∼4.37, which matches well with previous reports, further confirming the WGM lasing.³⁷ In order to get more insight into the lasing spectra, the resonant wavelength is derived and given by²⁷ 

where n₁ is the refractive index of the microcrystal, n₂ is the refractive index of the environment, $nr=n1n2$⁠, q = n_r for the transverse electric (TE) mode and q = 1/n_r for the transverse magnetic (TM) mode, m is the mode number, A_r is the root of the Airy function, and r is the radial mode number. Consequently, the lasing peaks can be well consigned to mode numbers from 381 to 386 for the first-order TM mode, as shown in Fig. 4(d).

The plot of the integrated PL intensity as a function of pump fluence is presented in Fig. 4(b), which exhibits a nonlinear behavior with an increase in pump fluence. This indicates the achievement of lasing action with a threshold (P_th) of 24.1 µJ/cm², which is comparable to those of the previous CsPbBr₃ WGM lasers (tens to hundreds of μJ/cm²).^37,39 Considering the cost-effective solution-processing method, such a low lasing threshold is highly desirable for the practical laser applications.^40,41 Moreover, the Q-factor is also extracted according to the equation Q = λ/Δλ,²¹ where λ and Δλ are the wavelength and the linewidth of the lasing peak, respectively. In our case, the linewidth of the resonant peak is determined to be ∼0.24 nm [Fig. 4(c)], and the Q-factor of ∼2246 is derived. Again, the Q value is similar to those of the CsPbBr₃ WGM lasers reported before and much larger than those of the CsPbBr₃ Fabry–Pérot lasers.^39,42 Furthermore, the lasing spectra from different-sized CsPbBr₃ microcrystals are investigated. Figure S6 shows the plot of mode spacing as a function of crystal size, manifesting that the mode spacing increases monotonously with the decrease in the crystal size.

Notably, the practically important single-mode lasing can be achieved in our solution-processed CsPbBr₃ microcrystal. Figure 4(e) displays the pump fluence-dependent PL spectra of the CsPbBr₃ microcrystal with a width of ∼4.0 µm. With the increase in pump fluence, the PL spectra transform from relatively broad spontaneous emission to a single sharp peak. The integrated PL intensity vs pump fluence exhibits the typical lasing behavior [Fig. 4(f)]. These results evidently confirm the development of single-mode lasing from the CsPbBr₃ microcrystal with a relatively low pump threshold of ∼28.1 µJ/cm². The Q-factor is derived to be ∼1900 (Fig. S7). It is noted that it is not straightforward to achieve single-mode lasing by simply reducing the crystal size because the optical loss would be exclusively deteriorated in this case. This is why the single-mode lasing is usually realized by more advanced and complicated techniques, such as the coupled cavities.⁴³ The success here could be attributed to the sufficient gain in the CsPbBr₃ microcrystal and the minor nonradiative recombination in the solution-processed samples.

It is noted that our solution-phase method allows the mass production; for example, thousands of MLPs can be obtained under one process. In fact, the scale of the fabricated IMHP-MLPs is only limited by the substrate size. As shown in Fig. S8, a 2-inch wafer is used as the substrate, and millions of MLPs are acquired at one batch. In addition to the glass, the CsPbBr₃ MLPs can also be grown on the conductive substrate, such as the indium tin oxide (ITO). Figure S9 shows the representative image of the CsPbBr₃ MLPs on the ITO. The finding could facilitate the light emission and lasing applications of the CsPbBr₃ MLPs by the current injection in the future. It is noted that our solution-phase method fails to produce the reproducible lasers due to the variation in the size and morphology of the CsPbBr₃ microcrystals. In the future, this issue may be addressed by applying the template-based approach, where the crystal size can be controlled by a predesigned template. Such scalable and high-performance MLPs could find a wide range of applications in the optoelectronic field. As a proof-of-concept, the gas sensing application is demonstrated (see Sec. IV for details). Figure 5(a) shows the lasing spectra of the CsPbBr₃ microcrystal before and after the acetone vapor is fluxed. The amplification fraction is shown in Fig. 5(b). The curve of 0 s represents the onset of acetone vapor injection in the system. After the injection, the acetone vapor could freely diffuse out of the beaker, and the spectrum is collected continuously at the intervals of 1 s. It can be seen that as the time increases, the vapor gradually diffuses out and the lasing peaks show a gradual blue shift. This phenomenon can be explained by the decrease in the refractive index of the surrounding medium with time. Since the acetone vapor continuously diffuses into the external atmosphere through the open window in the beaker, the concentration of the acetone vapor around the microcrystal continuously decreases, resulting in a decrease in the refractive index of the surrounding medium. This process continues until the surrounding environment of the microcrystal is pure air. By comparing the lasing spectra before vapor addition and after complete vapor diffusion, we found that the lasing peak nearly returned to its original position before the acetone vapor was added [Fig. 5(b)], thus verifying that the blue shift of the lasing peak position was indeed related to the decrease in the refractive index in the beaker due to the decrease in the acetone vapor concentration.

In summary, we report the mass production of CsPbBr₃ microcrystals with adjustable size by the solution-phase route. Importantly, the solution-prepared microcrystals exhibit much better optical properties compared to that prepared by the CVD method, which is attributed to the self-passivation by the bromine-rich surface. The CsPbBr₃ microcrystals with inverted pyramid morphology manifest the WGM lasing with low pump threshold and high Q-factor. Moreover, the practically desirable single-mode laser was achieved. Finally, the laser-emission-based gas sensor is demonstrated. Our work would promote the development of perovskite-based optoelectronic applications.

Fabrication of CsPbBr₃ microcrystals: The CsPbBr₃ powder was fabricated by dissolving PbBr₂ in hydrobromic acid and CsBr in deionized water. In this process, 10 mmol PbBr₂ was dissolved in hydrobromic acid (8 ml) and 10 mmol CsBr was dissolved in deionized water (3 ml) at room temperature. The dissolved CsBr was then added to the PbBr₂ solution to produce an orange color precipitate. The precipitate was filtered and cleaned 4–5 times with ethanol and then transferred to a drying oven at 60 °C for 12 h. After that, the CsPbBr₃ powder was obtained. The prepared CsPbBr₃ powder was dissolved in DMSO (dimethyl sulfoxide) for spin-coating. Glass sheets (0.5 × 0.5 cm) were used as the substrate for the growth of the CsPbBr₃ microcrystals. They were ultrasonic cleaned with acetone, ethanol, and deionized water successively for 10 min each time. Then, the glass sheets were dried in N₂ atmosphere and exposed to ultraviolet light for 25 min, which could improve the affinity of the glass surface. In the spin-coating process, the CsPbBr₃ precursor solution was heated to 60 °C and then dropped (20 μl) onto the glass substrate. The spin-coated sample was transferred to a thermostatic plate for annealing. The detailed experimental parameters are demonstrated in the results of each experiment.

Scanning electron microscopy (SEM) and x-ray diffraction (XRD) characterization: SEM images were taken on a FEI Quanta 250F field emission environment scanning electron microscope. X-ray diffraction patterns of the CsPbBr₃ microcrystals were measured using a Bruker D8 advanced x-ray diffractometer equipped with graphite-monochromatized Cu Kα radiation (λ = 1.5418 Å).

Optical characterization: The time-resolved fluorescence decay curves were measured by the 375 nm pulsed laser (model: EPL-375, Edinburgh Instruments) (pulse duration: 60 ps and frequency: 5 MHz) in the same test system and detected by a single photon counting photomultiplier module (model: id 100, Edinburgh Instruments). For the lasing characterization, a femtosecond amplified Ti:sapphire laser source (model: Solstice Ace, Spectra-Physics) (pulse duration: ∼100 fs, central wavelength: 800 nm, frequency: 1 kHz, and maximum output power: 7 W) is adopted. Then, a laser beam with the wavelength of 400 nm is generated through the double-frequency method. The laser beam (pump wavelength: 400 nm) is focused onto the individual perovskite microcrystal by an objective lens. The size of the focused laser beam was controlled by the objective lens combined with an aperture in the optical path. The emission signal was collected by a monochromator coupled with a charge coupled device (CCD) (model: iDus 420, Andor) via the same objective lens, and the images were taken by a high-resolution camera (Nikon).

Gas sensor demonstration: The solution-processed single CsPbBr₃ microcrystal was mounted inside a beaker. There was a hole in the beaker to allow the injection and out diffusion of the acetone gas. By monitoring the lasing peak shift, the acetone gas concentration can be monitored as a function of time.

See the supplementary material for the photographs of the precursor solution, Raman spectrum, emission linewidth as a function of pump fluence, length-dependent free spectral range, and crystals grown on the ITO.

This work was supported by the National Natural Science Foundation of China (Grant No. 11904172), the Natural Science Foundation of Jiangsu Province (Grant Nos. BK20190446 and BK20210592), and the NUPTSF (Grant No. NY221030). Y.W. thanks the support of the start-up funding from the Nanjing University of Science and Technology.

The authors have no conflicts to disclose.

Y.R. and C.M. contributed equally to this work.

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