Interface-directed synthesis of CsPbBr₃-based particles with water-ODE antisolvent systems

Abstract

All-inorganic lead halide perovskites, specifically CsPbX₃ (X = I, Cl, Br), have attracted significant attention over the past few years due to their immense potential in optoelectronic devices, such as liquid crystal displays (LCDs), light emitting diodes, photodetectors, and solar cells. Currently, most available synthesis techniques rely on high temperatures, inert environments, and toxic non-polar solvents, with limiting scalability and environmental compatibility. Herein, we present a facile, room-temperature, and partially green approach for synthesizing CsPbBr₃-based particles via a biphasic solvent system comprising non-polar octadecene (ODE) layered over deionized water. By simply tuning the volume ratio of DI water to ODE, we used interfacial engineering to modulate the formation and growth of CsPbBr₃-based particles. Pure water system without ODE yielded highly crystalline, phase-pure CsPbBr₃ of ∼0.5 µm in size (BPPT0) with strong green photoluminescence at ∼534 nm. A system with 1 : 1 of water to ODE favored the formation of small CsPbBr₃ particles of ∼0.1 µm in size with trace amounts of Cs₄PbBr₆ impurity (SPPT1). The as-synthesized SPPT1 particles exhibited strong photoluminescence at ∼524 nm and broader color gamut coverage of ∼123% of NTSC 1953 and ∼174% of sRGB when being incorporated in white emitting LCD backlight structures. Further, both BPPT0 and SPPT1 displayed excellent ambient stability. This work demonstrates interfacial solvent engineering as a powerful strategy for the synthesis of size-controlled perovskite under scalable and environmentally responsible conditions.

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1 Introduction

All-inorganic lead halide perovskites of CsPbX₃ (X = Cl, Br, I) have emerged as a key player in developing photovoltaic devices such as LEDs (light emitting diodes), solar cells and photodetectors owing to their excellent opto-electronic properties namely high defect tolerance, remarkable thermal stability, fine emission bandwidth, high color purity etc.^1–7 However, realizing these materials in practical devices critically depends on the ability to fabricate high-quality, phase-pure, and morphologically uniform materials with minimal defects.^8–12

Over the past few years, several synthesis techniques have been implemented to fabricate high-quality green emitting CsPbBr₃ particles with high crystallinity and notable opto-electronic performance. For instance, the standard hot-injection synthesis method enables high-quality CsPbBr₃ nanocrystal formation with tunable morphology and excellent optoelectronic properties, but typically requires elevated temperatures (140–300 °C) and strict inert conditions with controlled precursor injection, which complicate large-scale production.^13–16 Another widely adopted room-temperature synthesis route for CsPbBr₃ nanocrystals is ligand-assisted reprecipitation (LARP) method which offers simplicity and operation at ambient temperatures (<70 °C) with minimal equipment requirements, however, it still relies heavily on organic and toxic solvents (such as toluene and hexane) and lengthy ligand purification steps, and achieving reproducible quality at high yields remains challenging.^17–19

The strong ionic behaviour of CsPbBr₃ perovskites further aggravates these limitations, making them highly unstable in the presence of moisture and polar solvents such as water, acetone, and ethanol etc.^20–23 Such exposure often leads to rapid phase transformation from CsPbBr₃ to non-radiative Cs₄PbBr₆ and/or the decomposition of CsPbBr₃.^24,25 Other synthesis approaches for CsPbBr₃ particles, such as solvothermal and microwave-assisted methods, have been shown to enhance crystallinity and accelerate reaction kinetics, however, these strategies often demand elevated temperatures, autogenous pressures, and specialized reactor setups, which often complicate the process control and pose challenges for large-area or high-throughput scale-up strategies.^26–28

Tremendous efforts have been invested into understanding the degradation mechanism of CsPbBr₃ at atomistic scale in polar environments.^24,25,29 On contrary, several reports highlighted the correlation of moisture with enhanced opto-electronic behavior of perovskite materials including CsPbBr₃ with better crystallinity and lower surface defects.^30–33 A few water-assisted synthesis methods have recently emerged as promising alternatives to improve crystallinity, reduce toxicity, and modify band gap or consequently the luminescent behavior of the prepared perovskites.^34–40 For instance, Zhang, et al.³⁶ introduced small traces of water in toluene when synthesizing CsPbBr₃ particles using LARP method and reported the change in size, crystallinity and phase purity with different amounts of water introduced leading to red shift of PL (photoluminescence) peak and enhanced opto-electronic properties such as higher PLQY (quantum yield) and better stability.³⁶ However, such reports not only demand non-polar solvent environments during synthesis but also lack the mechanisms controlling nucleation and growth, phase purity, and size selectivity.

There are various approaches implemented to control the nucleation and growth kinetics of CsPbBr₃ such as solvothermal, sonication, mechanochemical and microwave irradiation.^41–44 These techniques utilize non-polar solvents such as benzyl ether (BE) and 1-octadecene (ODE) during the synthesis to achieve better control of crystallinity and morphology of perovskite particles including CsPbBr₃, owing to their non-polar behavior and high boiling point. Zamani, et al.⁴⁵ studied the effect of non-polar solvents on the growth of CsPbBr₃ particles and concluded that using BE in one-pot synthesis and microwave irradiation resulted in formation of highly crystalline nanoplates while using ODE resulted in formation of smaller crystallites owing to effective solvation of PbB₂ in BE compared to ODE. However, such methods still rely on high temperatures, high energy radiations, inert atmosphere, complex synthesis etc.⁴⁵ Therefore, achieving a synthesis technique that balances phase selectivity, crystallinity, morphology, optical performance, and ambient stability while being simple, cost effective and environment friendly up to a certain degree, under benign and scalable conditions, remain elusive.

In this work, we report for the first time a facile and scalable synthesis route, as shown in Fig. 1, under ambient conditions for CsPbBr₃-based particles via a biphasic system with non-polar ODE over polar deionized (DI) water, achieving a unique interfacial environment to control the particle size, phase purity of as-obtained CsPbBr₃-based particles, without the need for inert environments, high-temperature injection, or ligand-intensive stabilization. As a result, this approach provides a fundamentally distinct and scalable processing framework that is more compatible with low-cost, environmentally benign, and potentially continuous manufacturing of CsPbBr₃-based optoelectronic materials. By tuning this polar-non polar solvent interface, we obtained large precipitates (BPPT0) of ∼0.5 µm in size and small precipitates (SPPT1) of ∼0.1 µm in size. The as-obtained CsPbBr₃-based particles exhibit distinct phase composition, impurity levels, PL behavior, ambient stability, and display color profiles.

Fig. 1 Schematic of the synthesis route of CsPbBr₃-based particles.

2 Experimental details

2.1 Synthesis of CsPbBr₃-based particles

For the precursor preparation, CsBr (0.042 mmol) and PbBr₂ (0.042 mmol) were dissolved in 5 ml of N,N-dimethylformamide (DMF) containing 50 µl of oleylamine (OAm) and 187.5 µl of oleic acid (OA) at room temperature (∼16 °C) under vigorous stirring for ∼8 hours until a clear transparent solution is obtained as the “precursor” solution. Four different glass vials, pre-filled with 1 ml DI water each, were filled with different amounts of ODE (specifically, 1 ml, 0.5 ml, 0.25 ml, and 0 ml) to achieve volume ratios of 1 : x (x = 1, 0.5, 0.25, 0) for DI water to ODE, respectively. After two different immiscible layers were formed, the prepared precursor solution of 5 ml was poured into all four glass vials. Immediately, green emitting particles were formed at liquid–liquid and liquid–glass interfaces while most particles were observed on the bottom of the vials. These as-formed CsPbBr₃-based particles were extracted carefully and redispersed into 1 ml of toluene to separate bigger particles from the “colloidal” suspension.

2.2 Preparation of CsPbBr₃-based ‘films’

Three different kinds of ‘films’ were produced on glass substrate from the CsPbBr₃-based particles, which were formed with different ratios of DI water to ODE. Overall, 12 different films were fabricated and characterized.

2.2.1 Big precipitates (BPPTs) thin films. After extracting BPPTs from colloidal solution through natural sediment, they were deposited on clean glass substrate using a drop-caste method and were dried in open-air conditions for nearly 2 hours before characterization.

2.2.2 Small precipitates (SPPTs) thin films. The colloidal solutions after removing BPPTs were centrifuged at 8000 rpm for 5 min resulting in sediment of bigger particles on bottom. These particles were extracted carefully and were deposited on clean glass substrate via the drop-caste method and were dried in open-air for nearly 5–6 hours before characterization.

2.2.3 Supernatant (SNT) thin films. After centrifugating colloidal QDs and extracting SPPTs, 100 µl of the supernatant was deposited on clean glass substrate via the drop caste method and were dried in open-environmental conditions for 3–4 hours before characterization.

2.2.4 Red emissive phosphor film. A red EXPO marker pen was used to extract red ink. 100 µl of the ink was drop-casted on glass substrate, which was followed by open-air drying.

2.3 Materials characterization

The morphological and chemical compositional analyses were conducted on a scanning electron microscope (SEM) (JEOL JSM-5900lLV) equipped with an energy-dispersive X-ray spectroscope (EDS). The structural characterization of as-prepared CsPbBr₃ films was performed on an X-ray diffractometer (XRD) (Siemens D500) with CuKα radiation (λ = 1.5406 Å). The PL studies of the as-prepared CsPbBr₃ films were carried out using a spectrometer (Ocean optics, FLAME-S-VIS-NIR-ES, Ocean Optics, Orlando, FL, USA) with a UV excitation wavelength of 365 nm. The lifetime decay studies were conducted using a DeltaHub™ high throughput time correlated single photon counting (TCSPC) controller.

3 Results

Fig. 2a and b shows typical SEM images of BPPT0 (BPPT prepared using pure DI water) and SPPT1 (SPPT prepared using 1 : 1 ratio of DI water to ODE), respectively. Both the BPPT0 and SPPT1 exhibit cuboid-shaped morphologies in accord with the orthorhombic phase of CsPbBr₃. Fig. 2c and d shows the size distribution of the BPPT0 and SPPT1, respectively, which are derived from Fig. 2a and b. The average particle size of the BPPT0 is ∼500 nm significantly bigger than ∼100 nm of the SPPT1. The EDS spectral analysis, shown in Fig. S1a–d, demonstrated that the Cs : Pb : Br atomic ratios in both BPPT0 and SPPT1 are close to the 1 : 1 : 3 stoichiometry, consistent with CsPbBr₃ being the dominant phase in these samples.

Fig. 2 (a and b) SEM images of BPPT0 and SPPT1, and (c and d) histograms depicting the size distribution of the BPPT0 and SPPT1 shown in a and b.

Note that both the BPPT0 and SPPT1 exhibit high crystallinity compared to some previous works.⁴⁰ For instance, Xiang, et al.⁴⁶ reported lower crystallinity (shown via SEM images) of the prepared CsPbBr₃ films with average particle size of ∼0.47 µm via vacuum thermal evaporation.

Fig. 3a and b shows XRD spectra of the BPPTs and SPPTs synthesized with 4 different volume ratios of DI water to ODE. As evident from Fig. 3a, the BPPT sample prepared via pure DI water (DI water : ODE ratio of 1 : 0) exhibits only the characteristic reflections of orthorhombic CsPbBr₃ phase, with major peaks centered at 2θ ∼15.2°, 21.4°, 30.3°, 30.6°, 34°, 37.7°, and 43.7° corresponding to (110), (112), (004), (220), (130), (132), and (224) planes (PDF#97-009-7851), indicating phase-pure CsPbBr₃ within the detection limit of XRD.

Fig. 3 (a) XRD pattern of CsPbBr₃-Cs₄PbBr₆ mixed-phase BPPTs, and (b) XRD spectra of CsPbBr₃-Cs₄PbBr₆ mixed-phase SPPTs with different molar ratios of water to ODE.

On contrary, all other BPPT samples prepared with 1 : 1, 1 : 0.5, and 1 : 0.25 ratios of DI water to ODE, showed additional diffraction peaks at 2θ ∼12.6°, 22.4°, 25.4°, 27.5°, 28.6°, 38.9°, 45.8°, which are indexed to the (012), (300), (024), (131), (214), (324), (137) planes of hexagonal Cs₄PbBr₆ phase (PDF#97-002-5124), confirming the formation of CsPbBr₃-Cs₄PbBr₆ mixed-phase composites in these samples. Similar coexistence of dual phases is well reported in the literature.^47–49 Among these BPPTs (1 : 1, 1 : 0.5 and 1.0.25), the sample prepared with the ratio of 1 : 0.5 (DI water: ODE) exhibited the most intense peaks for secondary Cs₄PbBr₆ phase, suggesting preferential trap of Pb²⁺ ions at the water-ODE interface due to partially nonpolar environment and being consistent with the reports that partial nonpolar environments and excess CsBr favour the formation of Cs₄PbBr₆ phase.⁵⁰

Similar CsPbBr₃-Cs₄PbBr₆ phase coexistence was observed in all SPPT samples prepared with 1 : 1, 1 : 0.5, 1 : 0.25, and 1 : 0 ratios of DI-water to ODE (Fig. 3b). Moreover, in all SPPT samples, orthorhombic CsPbBr₃ was identified as the dominant crystalline phase, while hexagonal Cs₄PbBr₆ was present only as a minor secondary phase, indicated by its low-angle reflections at ∼20° and depicted in Fig. 3b. Using Bragg's law⁵¹ and the XRD patterns (Fig. 3a and b), we calculated the interplanar spacing (d) of the orthorhombic CsPbBr₃ phase for all BPPT and SPPT samples using the characteristic peak centred around ∼30.6°, corresponding to (220) crystallographic planes of orthorhombic CsPbBr₃. The mean interplanar spacing for the orthorhombic CsPbBr₃ reflections was found to be ∼2.903 ± 0.012 Å for BPPT samples and 2.911 ± 0.003 Å for SPPT samples, indicating that variations in processing conditions did not induce any measurable distortion in the orthorhombic CsPbBr₃ lattice structure.

Fig. S2a–c in the SI presents optical images of 12 films deposited on glass substrates using BPPTs, SPPTs, and SNT synthesized using four different ratios of DI water to ODE. The drop-casting concentrations used for producing BPPTs, SPPTs and SNT were determined quantitatively and are reported in Tables S1 and S2. For BPPTs, the highest concentration (69 µg µl⁻¹) was obtained for the 1 : 0 ratio (DI water: ODE), while the lowest concentration (27.2 µg µl⁻¹) was observed for the 1 : 1 ratio, with the 1 : 0.5 and 1 : 0.25 ratios yielding comparable intermediate concentrations (61.8 µg µl⁻¹ and 58.1 µg µl⁻¹). These concentration variations directly impact the resulting film morphology. Accordingly, BPPTs-derived films (Fig. S2a) exhibited rough and heterogeneous surfaces, with the dense dark-yellow domains observed for the 1 : 0 ratio (pure DI water) arising from the higher BPPTs loading and the presence of CsPbBr₃ as the dominant phase, whereas lower concentration samples (1 : 1, 1 : 0.5 and 1 : 0.25) displayed sparse deposition and visible white precipitates, which can be correlated with the increased detectability of the Cs₄PbBr₆ secondary phase in the corresponding XRD patterns (Fig. 3b).

In contrast, the SPPTs exhibited the highest concentration for the 1 : 1 ratio of DI water to ODE (70.9 µg µl⁻¹), followed by the 1 : 0.5 (38.1 µg µl⁻¹) and 1 : 0.25 ratios (30.9 µg µl⁻¹), while the 1 : 0 ratio showed the lowest concentration (16.3 µg µl⁻¹), as summarized in Tables S1 and S2. As a result, SPPTs-derived films (Fig. S2b) prepared with 1 : 1 ratio of DI water to ODE depicted the most uniform coverage, whereas the films prepared from the lower drop-casting concentrations displayed localized particle accumulation and discontinuous regions (Fig. S2b). A similar qualitative trend is observed for SNT-derived films (Fig. S2c), where the films prepared via 1 : 1 ratio of DI water to ODE showed visibly higher particle coverage, while progressively reduced coverage is observed at lower ODE amounts. These observations for SNT are based on visual inspections rather than quantitative concentration measurements owing to the low particle yield in the supernatant fractions. Overall, these results demonstrate that synthesis-dependent variations in particle concentration play a dominant role in governing the resulting film's uniformity.

The films in this study were prepared by simple drop-casting and were not optimized for thickness or lateral uniformity, as evident from Fig. S2, and the resulting coatings are discontinuous and rough, reflecting both synthesis-dependent concentration variations (Tables S1 and S2) and the inherent limitations of drop-casting technique.⁵² Consequently, the reported optical results represent ensemble-averaged emission from deposited particle (BPPTs, SPPTs and SNT) layers, rather than from device-grade continuous thin films. Further optimization of synthesis routes and deposition strategies will be required to achieve continuous films suitable for practical device applications.⁵³

Fig. 4a–d shows PL characteristics of the prepared films (Fig. 4a and c), colloidal solutions (Fig. 4b), and supernatants (SNT) (Fig. 4d), which were prepared with different ratios of DI water to ODE. From Fig. 4a–d, the PL characteristics of the prepared films, colloidal solutions, and supernatants are analysed and presented in Fig. 4e–f for the PL peak intensity, peak wavelength, and FWHM (Full Width at Half Maximum), respectively.

Fig. 4 PL spectra of (a) BPPTs, (b) colloidal solutions, (c) SPPTs, (d) SNT; variations of PL peak intensity (e), peak wavelength (f) and FWHM with the water : ODE ratio.

For the films made from the BPPTs, increasing the ODE fraction causes the decrease of the PL peak intensity with the one prepared with pure DI water (1 : 0 ratio of DI water to ODE) delivering the most intense PL peak intensity, as shown in Fig. 4e. This can be attributed to the combined effects of the highest drop-casting concentration (69 µg µl⁻¹) in the 1 : 0 (pure DI water) BPPT sample and the formation of a predominantly single-phase orthorhombic CsPbBr₃ (Fig. 3a) with minimal surface-bound ligands. The absence of non-luminescent Cs₄PbBr₆ secondary phase also minimizes the effective non-radiative recombination pathways, resulting in enhanced radiative emission.⁵⁴

For the films made from SPPTs, the PL peak intensity increases with increasing the ODE fraction, as shown in Fig. 4e. The film with the SPPTs prepared from the 1 : 1 ratio (DI water : ODE) delivers the highest PL peak intensity, which can be attributed to the combined effects of the highest drop-casting concentration (70.9 µg µl⁻¹) in this sample and the dominance of the CsPbBr₃ phase (Fig. 3b). Although a trace amount of the non-luminescent Cs₄PbBr₆ secondary phase is detected by XRD (Fig. 3b), its minimal presence does not significantly impact the overall radiative emission, allowing the PL intensity to remain maximized (Fig. 4e). A similar pattern has been observed for SNT-derived films (Fig. 4e), where the film prepared using a 1 : 1 ratio of DI water to ODE exhibited the highest PL peak intensity (Fig. 4e); however, their corresponding drop-casting concentrations were below the detection limit of the available instrumentation.

Fig. 4f and g summarize the dependence of the PL peak wavelength and FWHM on the DI water : ODE ratio used to synthesize BPPTs, SPPTs, and SNT particles. For BPPT-derived films, the PL maximum initially exhibited a pronounced blue-shift (from ∼534 nm to ∼526 nm) as the DI water : ODE ratio is varied from 1 : 0 to 1 : 0.5, shown in Fig. 4f, indicating a progressive suppression of aggregation-induced red-shift effects under biphasic growth conditions. A slight red-shift (to ∼528 nm) observed for the 1 : 1 ratio of DI water to ODE (Fig. 4f) suggests the presence of competing solid-state optical processes, such as enhanced reabsorption or local particle accumulation under highly biphasic environments. However, the SPPT-derived films exhibited a relatively modest monotonic blue-shift (from ∼528 nm to ∼524 nm) as the DI water : ODE ratio is varied from 1 : 0 to 1 : 1 (Fig. 4f), consistent with their smaller particle size and weaker inter-particle optical coupling. Further, the SNT-derived films showed minor variations in PL peak positions (from ∼521 nm to ∼524 nm) with no clear monotonic trend as the DI water : ODE ratio was varied (Fig. 4f), which can be attributed to their low drop-casting concentration and limited solid-state interactions. Moreover, the linewidth values (Fig. 4g) displayed no pronounced systematic dependence on the ODE ratios, although BPPT0 exhibited the broadest emission, consistent with stronger aggregation-induced spectral broadening. In comparison, SPPTs and SNTs show relatively narrow and nearly invariant FWHM (variations <1–2 nm), indicating more homogeneous optical environments and reduced ensemble broadening effects.

Notably, the BPPT-derived films synthesized using pure water (1 : 0 ratio of DI water to ODE) exhibited a PL peak maximum at ∼534 nm, which is red-shifted relative to PL emission wavelengths typically reported for well-dispersed CsPbBr₃ nanocrystals (∼515–520 nm).^55,56 However, red-shifted PL emission (≥530 nm) has often been reported for solid-state CsPbBr₃ films.^57,58 For instance, Zalrhi, et al.⁵⁷ reported a PL emission peak at ∼2.3 eV (∼539 nm) for pristine CsPbBr₃ films deposited via spin coating at 5000 rpm for 30 s. Yeh, et al.⁵⁸ fabricated green-emitting CsPbBr₃ films via thermal evaporation, exhibiting PL emission at ∼533 nm with narrow linewidth of <18 nm. Moreover, close packing in CsPbBr₃ films can induce a red-shit in PL emission wavelengths compared to less-coupled or dispersed counterparts, consistent with inter-particle optical coupling and radiative transfer in dense films.⁵⁹ Such red-shifts are further supported by photon-recycling behavior (self-absorption and re-emission), which has been explicitly shown to produce PL red shifts that increase with increasing emitter concentration and optical path length.⁶⁰ Therefore, the PL emission at ∼534 nm observed for BPPT-derived films (1 : 0) can be attributed to the dense and heterogeneous drop-cast regions that enhance the reabsorption-re-emission processes and modify the ensemble emission profiles.

Based on their highest drop-casting concentration, strongest PL intensity (Fig. 4e), and dominant orthorhombic CsPbBr₃ phase (Fig. 3a), the 1 : 0 (pure DI water) BPPT and 1 : 1 (DI water: ODE) SPPT-derived films were selected as representative samples for all subsequent analyses and are hereafter referred to as “BPPT0” and “SPPT1”, respectively. Notably, BPPT0 exhibits phase-pure orthorhombic CsPbBr₃ character (with no detectable amount of Cs₄PbBr₆ phase), whereas SPPT1 is dominated by the orthorhombic CsPbBr₃ phase with only trace amounts of the secondary Cs₄PbBr₆ phase, as identified by their corresponding XRD spectra (Fig. 3a and b).

Fig. 5a and b shows the PL spectra over five weeks for the films from the BPTT0 prepared with pure water and the SPPT1 prepared with the 1 : 1 ratio, respectively. The insets in individual figures present optical images of corresponding fresh films under room light (RL) and UV light of 365 nm in wavelength. In general, the PL emission of the films remains relatively stable. Fig. 5c and d shows the temporal evolution of PL intensities for both the films, as derived from Fig. 5a and b The PL intensities of both the films exhibit a nearly linear decrease in PL intensity over the test period, indicating gradual photodegradation under ambient conditions (T ∼18 °C, relative humidity (RH) ∼60%, and normal indoor laboratory lighting). Under such exposure, metal-halide perovskites (including CsPbBr₃) are well known to undergo moisture-assisted degradation involving halide migration, surface defect formation, lattice distortion, and oxidation, which collectively enhance non-radiative recombination and suppress PL emission.^61–65

Fig. 5 PL spectra at different weeks for the films from (a) the BPTT0 prepared with pure water (inset: optical images of fresh films under room light (RL) and UV light (365 nm in wavelength)), (b) the SPPT1 prepared with the 1 : 1 ratio (inset: optical images of fresh films under RL and UV light); temporal evolution of PL intensity for the films from (c) the BPPT0 and (d) the SPPT1; temporal dependence of PL wavelength for the films from (e) the BPPT0 and (f) the SPPT1; and (g) optical images of the films from the SPPT1 prepared with the 1 : 1 ratio (left) and BPPT0 with pure water (right), after 5 weeks.

Using the linear regression to fit the results in Fig. 5c and d, we obtain the degradation rates of 1494 and 775 in the unit of a.u. per week for the films from the SPPT1 and BPPT0, respectively. The films from the SPPT1 degrade nearly twice as fast as the ones from the BPPT0. Such faster degradation for smaller particles (SPPT1) is consistent with previous reports dictating the critical role of surface-to-volume ratio in determining the stability of perovskite materials.^66,67 Smaller particles can experience faster moisture and oxygen-induced degradation due to their higher surface defect densities and larger environmental exposure.

Importantly, CsPbBr₃ perovskites are highly susceptible to structural degradation when exposed to moisture, with water molecules disrupting the perovskite lattice and initiating decomposition pathways, ultimately degrading their optoelectronic properties.⁶⁸ For instance, Di Girolamo, et al.⁶⁹ reported a partial degradation of CsPbBr₃ perovskite phase under prolonged exposure to humidity, accompanied by the formation of secondary crystalline phases in these films, consistent with moisture-assisted phase evolution during aging.^68–72 These previously established degradation pathways^68–72 align well with the monotonic PL decay observed in Fig. 5c and d.

Interestingly, the PL peak wavelength for the film made from the BPPT0 exhibits different evolution from the one for the film made from the SPPT1, as shown in Fig. 5e and f. Precisely, the PL peak wavelength for the film made from the BPPT0 shows an exponential decay over time, indicating a continuous blue-shift of PL peak over time. The temporal evolution of the PL peak wavelength is fitted with a single exponential decay model, as shown in Fig. 5e. This blue-shift might arise due to a variety of mechanisms such as loss of surface halide, partial degradation of larger particles, or ion migration, defects formation and lattice distortions.^73–76 These factors are among unique properties of halide perovskites, which in this case result in the widening of band gap due to reduction in effective-negative antibonding overlap between 6 s Pb and 4p Br orbitals of CsPbBr₃ ultimately causing blue shift.^73–77 For instance, Phung, et al.⁷⁷ suggested that lattice distortions can cause the widening of band gap due to introduction of defects in MaPbI₃, leading to the blue shift of its PL peak.

In contrast, the PL peak wavelength for the film made from the SPPT1 depicts a non-monotonic PL peak shift, initially exhibiting red-shift followed by blue-shift over time. Such behavior resembles a Gaussian profile, which is well-fitted by a Gaussian function (Fig. 5f), and can be attributed to occurrence of multiple competing processes such as initial surface healing, trap-state passivation or possibly agglomeration. All these possibly cause the red shift. After the read shift, degradation-driven blue-shift might occur due to enhanced antibonding interactions between Pb and Br orbitals, leading to wider band gap and consequently.^73–77

Overall, such results highlighted the relative photostability advantages of larger CsPbBr₃ particles (BPPT0) synthesized with pure water only, which offers slower degradation pathways and more predictable spectral evolution over time. In contrast, although, the films from the SPPT1 prepared with 1 : 1 ratio initially exhibit a higher PL intensity, they undergo rapid and more complex degradation dynamics due to their smaller size (high surface-to-volume ratio), higher defect density, and greater environmental reactivity.^66,67

Fig. 5g shows optical images of BPPT0 and SPPT1-derived films after storage under ambient conditions (T ∼18 °C and RH ∼60%) for five weeks. The BPPT0-derived films retained relatively larger and more connected particle domains over time (Fig. S2a and 5g), consistent with the higher BPPT loading and larger particle sizes. In contrast, the SPPT1-derived films, which initially exhibits higher surface coverage (Fig. S2b), evolves into spatially separated particle domains distributed across the substrate after five weeks (Fig. 5g). This behaviour suggests progressive film fragmentation or island formation, which can be attributed to weaker interparticle cohesion, reduced effective film thickness, and the higher surface-area-to-volume ratio of the smaller SPPTs. These factors enhance the susceptibility to moisture- and oxygen-induced redistribution, leading to morphological instability over time. This interpretation is consistent with the comparatively faster PL degradation (Fig. 5d) and the associated Gaussian-like spectral evolution observed in Fig. 5f.

Fig. 6 shows the time-resolved PL decay curves for the films derived from the BPPT0 and SPPT1, and these curves can be fitted using tri-exponential decay model as given below:^78,79

I_TRPL = A₁e^−t/τ₁ + A₂e^−t/τ₂ + A₃e^−t/τ₃ (1)

Here, A_i (i = 1, 2 and 3) and τ_i are the amplitude and lifetime of the i-th component, respectively.

Fig. 6 Time-resolved PL decay curves for the films from the BPPT0 and SPPT1 synthesized with 1 : 0 (pure water) and 1 : 1 ratios, respectively.

For the BPPT0-derived film, three decay components were obtained using eqn (1) namely τ₁ = 0.29 ns (6.75%), τ₂ = 2.26 ns (15.97%), and τ₃ = 39.85 ns (77.28%), yielding an amplitude-weighted average lifetime of <τ> = 39.61 ns with a reduced χ² of ∼1.55. In contrast, the SPPT1-derived film exhibited τ₁ = 0.25 ns (10.70%), τ₂ = 2.22 ns (70.50%), and τ₃ = 17.48 ns (18.80%), corresponding to <τ> = 3.84 ns with a reduced χ² of ∼1.34 (Table S3). Based on both the lifetimes and their amplitude fractions, the longest decay component (τ₃) is assigned to band-edge radiative recombination, the intermediate component (τ₂) to surface-related recombination, and the fastest component (τ₁) to deep trap-assisted non-radiative recombination.^79,80 The faster radiative rate for the SPPT1 (τ₃ = 17.48 ns) compared to the larger BPPT0 (τ₃ = 39.85 ns) is consistent with previous reports given by Abbas, et al.,⁸¹ who concluded that radiative decay lifetime in weak confinement region decreases with decreasing particle size due to thermal accessibility to “dark” states for CsPbBr₃ materials with different phases including orthorhombic phase. Another factor for the higher radiative decay lifetime for the BPPT0 might be smaller surface area-to-volume ratio for the BPPT0 compared to the SPPT0 leading to lesser surface defect density and comparatively lesser non-radiative recombination routes.^82–84 The faster radiative lifetime for the film from the BPPT0 (τ₂ = 39.85 ns), which were synthesized at room temperature (T ∼16 °C), is significantly higher compared to other reports in literature for halide perovskite and other luminescent films. For instances, Bai, et al.⁷ achieved a maximum radiative lifetime of ∼6 ns via the annealing of CsPbB₃ films at a temperature of 460 °C. Tang, et al.⁵⁴ reported a radiative decay lifetime of ∼17 ns for the CsPbB₃ thin films synthesized via antisolvent method with toluene being the antisolvent, which is noticeably lower than 39.85 ns for the film from the BPPT0 synthesized via pure water, noting the superiority of water as an antisolvent to produce CsPbBr₃ films of higher quality.

Fig. 7a and b depicts the PL curves at different temperature for the films from the BPPT0 and SPPT1 synthesised with 1 : 0 (pure water) and 1 : 1 ratios, respectively. Using the PL spectra, the variation of the corresponding PL peak intensity with temperature is presented in Fig. 7c and d, respectively, for the films. The decline in the PL peak intensity with rising temperature can be attributed to the increase of the dissociation of excitons and enhanced phonon scattering at higher temperatures, altering the carrier's effective masses leading to slower radiative recombination rates. Such phenomena of thermal quenching in PL intensity are common for most luminescent materials.^85–87 The temperature dependence of the PL peak intensity can be described by the following equation:^85–88

where I(T) and I(0) are the PL peak intensities at temperatures of T and 0 K, respectively, B is a constant, k_b is the Boltzmann constant, and E_a is the activation energy.

Fig. 7 PL spectra at different temperatures for the films prepared with the (a) BPPT0, and (b) SPPT1; variation of PL peak intensity with temperature for the films prepared with the (c) BPPT0, and (d) SPPT1; variation of PL wavelength with temperature for the films prepared with the (e) BPPT0, and (f) SPPT1.

The temperature dependence of the PL intensity for the films (Fig. 7c and d) are curve-fitted with eqn (2), yielding thermal activation energies of 100.21 meV and 54.11 meV for the films made from the BPPT0 and SPPT1, respectively. These fitted activation energies can be interpreted as the effective thermal barriers for PL quenching (energy required to activate competing non-radiative recombination pathways), rather than the intrinsic exciton binding energy of CsPbBr₃ perovskite structures.^89–93 For bulk and high-quality CsPbBr₃ perovskites, the intrinsic exciton binding energies are typically reported in the range of ∼18–44 meV, as determined from optical and magneto-optical spectroscopy studies of perovskite single crystals and thin films.^91–93 In contrast, the larger values obtained for BPPT0 (∼100.21 meV) and SPPT1 (∼54.11 meV) may reflect the combined influence of thermally activated carrier trapping, exciton delocalization, and phonon-assisted non-radiative recombination processes in these micron-sized CsPbBr₃ based particles.⁸⁹ Further, the higher activation energy observed in BPPT0 (∼100.21 meV) indicates that non-radiative pathways are less readily thermally activated in the larger particles in contrast to SPPT1 where a lower value of activation energy (∼54.11 meV) can be attributed to the higher surface-to-volume ratio of the smaller particles, increasing the density of surface-associated trap states and facilitating thermally induced PL quenching at lower energies.⁹⁴ This is consistent with longer radiative lifetime observed in BPPT0 (τ₃ = 39.85 ns) compared to SPPT1 (τ₃ = 17.48 ns) (Fig. 6), indicating suppressed non-radiative recombination pathways in the larger BPPT0 particles.

Fig. 7e and f shows the linear temperature dependence of the PL peak wavelength for the films from the BPPT0 and SPPT1 synthesised with 1 : 0 (pure water) and 1 : 1 ratios, respectively, which depicts blue-shift with increasing temperature and can be described by the following equation:⁸⁵

λ(T) = λ(T₀) + β_TΔT (3)

for |ΔT/T₀| ≪ 1. Here, λ(T) and λ(T₀) are the PL peak wavelengths at temperature T and a reference temperature T₀, and β_T is a factor determining the temperature dependence of the PL peak wavelength. Using eqn (3) to fit the data in Fig. 7e and f, we obtain the numerical values of β_T as −0.01 nmK⁻¹ and −0.008 nmK⁻¹ for the films from the BPPT0 and SPPT1 synthesised with 1 : 0 (pure water) and 1 : 1 ratios, respectively.

The blue shift of the PL wavelength can be attributed to various factors, such as enhanced electron–phonon interactions at higher temperatures, thermally induced defects, lattice distortions etc., which hinder the negative antibonding overlapping of 6s Pb and 4p Br orbitals of orthorhombic CsPbBr₃ and increases the band gap.^73–77,95 A smaller numerical value of β_T for the film from the BPPT0 compared to the one from the SPPT1 points towards better PL stability of the BPPT0 compared to the SPPT1 likely due to reduced electron–phonon interactions,⁹⁵ which is consistent with a longer radiative decay lifetime (τ₃ = 39.85 ns) and a lower rate of degradation (−775) for the BPPT0 synthesized via pure water.

Fig. 8 presents a qualitative proof-of-concept demonstration of white light generation by combining the green emission from CsPbBr₃-based particles (BPPT0 and SPPT1) with complementary red and blue fluorescent components. A commercial red EXPO fluorescent marker was employed as a readily available broadband red emitter to verify spectral overlap and color-mixing capability rather than to serve as a device-optimized red phosphor. Using the films made from the BPPT0 (λ ∼532 nm) and SPPT1 (λ ∼524 nm) with a red emissive phosphor film (λ ∼610 nm) and following the approach reported by Tang, et al.,⁸⁵ we constructed a qualitative ‘white’ emitting LCD backlight configuration (Fig. 8a). The LCD backlight structure was activated by a blue LED (∼450 nm in wavelength). Fig. 8b and c shows the corresponding PL emission spectra of the resulting ‘white’ backlights based on BPPT0 and SPPT1. The insets depict optical images of the corresponding backlight structures emitting bright white light, similar to the one demonstrated by Tang, et al.,⁸⁵ who mimicked a white-emitting LCD backlight structure using a green-emitting PMMA-PbBr₃ film with PL emission peak at ∼537 nm.

Fig. 8 (a) Schematic illustration of the experimental setup for the RGB backlight configuration (R: red phosphor film, G: green emitting perovskite film (SPPT1/BPPT0), B: blue LED); PL spectra for white emitting backlight using (b) BPPT0 and (c) SPPT1 green films (insets: optical images of white emitting backlight structures).

4 Discussions

The layered water/ODE system offers a biphasic environment with a water-ODE interface, in which ODE forms a non-polar layer over polar water layer. Such a system likely separates the nucleation and growth of orthorhombic CsPbBr₃ particles spatially. Upon injecting the DMF precursor with CsBr, PbBr₂ and small amounts of ligands, nucleation occurs at the ODE-water interface, as supported by the presence of green emitting particles. Polar water quickly solvates precursor ions, promoting nucleation and growth of CsPbBr₃ particles, while non-polar ODE layer modulates the diffusion of ions across the interface. The amount of ODE plays a pivotal role in dictating ionic diffusion, nucleation zone, crystal growth, and consequently, the particle size, phase as well as PL behaviour of final obtained particles. A flowchart summarizing the chemical kinetics involved in this water/ODE layered system is presented in Fig. 9, respectively.

Fig. 9 A flowchart for the formation and growth of CsPbBr₃ particles in a water/ODE layered system.

For the BPPTs, formed predominantly via rapid growth kinetics, the chemical reaction generally follows the stoichiometric perovskite formation as:

CsBr + PbBr₂ → CsPbBr₃ (4)

In a pure water system, the absence of ODE leads to fast and homogeneous mixing of precursor ions in this aqueous medium, resulting in uncontrolled nucleation and formation of large bulk precipitates (BPPT0) which were settled at the bottom and exhibited pure orthorhombic CsPbBr₃ phase with no detectable trace of Cs₄PbBr₆ phase, as verified by XRD (Fig. 3a). However, increasing the ODE amount (ratios of 1 : 0.25, 1 : 0.5, and 1 : 1 for water: ODE systems) led to excess CsBr availability from altered solvation and ion diffusion dynamics which led to partial formation of the Cs-rich Cs₄PbBr₆ phase with reaction involved below:

CsPbBr₃ + 3CsBr ⇄ Cs₄PbBr₆ (5)

The coexistence of double phases is not only evidenced by the XRD peaks depicted in Fig. 3a but also is well reported in literature.^47–49

For SPPTs, which consist of smaller particles with higher surface-area-to-volume ratios, the detectability of secondary phases (such as Cs₄PbBr₆) was observed to be more sensitive to local chemical fluctuations and drop-casting concentrations. In the pure DI water system, the limited solubility of PbBr₂ and reduced ionic transport across the reaction environment may promote local compositional inhomogeneity, leading to the emergence of detectable Cs₄PbBr₆ secondary reflections alongside the dominant orthorhombic CsPbBr₃ phase. Introduction of ODE creates a biphasic environment that moderates nucleation and growth dynamics at the water-ODE interface, thereby influencing the effective particle concentration of the resulting dispersions. Accordingly, the SPPT1 sample prepared with 1 : 1 ratio of DI water to ODE, which also exhibits the highest drop-casting concentration (Tables S1 and S2) among other SPPTs, shows the strongest green PL emission (Fig. 4e), while CsPbBr₃ remains the dominant phase with only trace amounts of Cs₄PbBr₆ secondary phase.

All subsequent studies have been focused on the films from the BPPT0 prepared with pure water and the SPPT1 prepared with the 1 : 1 ratio according to the combinational factors of particle size, phase purity, and optical performance. In the pure water system (BPPT0), the absence of ODE allows for rapid nucleation and uncontrolled growth of crystals due to the lack of interfacial stabilization. This results in the formation of larger CsPbBr₃ precipitates of ∼500 nm (Fig. 2a) with no detectable amounts of non-luminescent Cs₄PbBr₆ phases (Fig. 3a). In contrast, the biphasic system of the 1 : 1 ratio (SPPT1) offers a more balanced environment for the formation and growth of CsPbBr₃ due to presence of the water-ODE interface, which effectively moderates the nucleation and growth process. This results in formation of smaller CsPbBr₃ precipitates of ∼100 nm (Fig. 2b) with enhanced colloidal stability and stronger green emission, alongside formation of hexagonal Cs₄PbBr₆, as confirmed by the XRD analysis (Fig. 3b).

Fig. S3 in SI shows the absorbance and PL spectra for the films made from the BPPT0 and SPPT1 films. It is evident that there exists size dependence of the absorption and PL emission under UV light of ∼365 nm in wavelength. The SPPT1 exhibit an absorption peak at ∼521 nm and a PL emission peak at ∼524 nm, leading to a Stokes shift of ∼20 meV; the BPPT0 depict an absorption peak at ∼524 nm and a PL emission peak at ∼534 nm, leading to a Stokes shift of ∼50 meV. This further suggests the inverse dependence of Stokes shift on particle size, which are well reported in multiple experimental and theoretical studies for many perovskite materials including CsPbB₃ particles.^96–98

The difference between the band gap (E_g) of 2.36 eV for the SPPT1 and 2.31 eV for the BPPT0 at ∼16 °C reveals the size dependence of optical properties of orthorhombic CsPbBr₃. Such a small difference in band gap with the size difference of ∼400 nm between the SPPT1 and the BPPT0 might be attributed to the weaker confinement effect due to large difference between the particle sizes (both the SPPT1 and BPPT0) and the Bohr radius of CsPbBr₃.^99,100 Thereafter, the temperature dependence of band gap for both the BPPT0 (Fig. 10a) and the SPPT1 (Fig. 10b) are fitted with the Bose–Einstein relation as:^101–103

where E₀ (T) represents the band gap at absolute zero, a describes the strength of electron–phonon interaction, and b is the characteristic temperature linked to average photon energy.

Fig. 10 Temperature dependence of photon energy for the (a) BPPT0 and (b) SPPT1.

Eqn (6) is used to fit the temperature dependence of the band gap, as shown in Fig. 10a and b. It is evident that the parameter ‘a’ is negative for both the BPPT0 and SPPT1, revealing different behaviour of these CsPbBr₃ particles from the trend revealed by Huang, et al.,¹⁰³ who recently reported an decreasing trend of the band gap with increasing temperature (positive ‘a’) for CsPbBr₃ films with grain size of ∼0.52 µm and reduced electron–phonon interaction (lower ‘a’) for larger particles. However, the increasing trend of the band gap with increasing temperature is similar to PbS and PbSe quantum dots,⁹⁵ and such behaviour has also been observed for all inorganic perovskite materials. For instance, Li, et al.⁹⁵ reported the increase in the band gap of CsPbBr₃ particles (E_g = 2.35 eV to 2.39 eV) with increasing temperature from 80 K to 300 K) with the increase rate of the band gap depends inversely on the particle size.

The temperature dependence of FWHM (Fig. 11a and b) is analysed for both the films using the Bose model as:^104,105

where Γ₀ represents inhomogeneous broadening at 0 K, α refers to the degree of phonon scattering, β denotes the coupling strength between excitons and longitudinal optical (LO) phonons, and ħω corresponds to LO phonon energy respectively. The fitting results with eqn (7) are also included in Fig. 11a and b, respectively.

Fig. 11 Temperature dependence of FWHM with the fitting of Bose model (red dotted curve) for the films made from (a) BPPT0 and (b) SPPT1.

According to the fitting results given in Fig. 11a and b, the LO phonon coupling energy of the BPPT0 is ∼43 meV higher than ∼25 meV of the SPPT1) and further higher the one given by Huang, et al.,¹⁰³ who reported an LO phonon coupling energy of ∼20 meV for PMMA-CsPbBr₃ films with CsPbBr₃ of micron sizes (∼0.5 µm) synthesized via a mechanochemical method. Such a higher value of the coupling energy might attribute to the combination of internal stresses, reduced structural disorder, and modified chemical environments introduced via the polar-nonpolar water: ODE systems. This result suggests the sensitivity of the LO phonon dynamics altered with subtle variations in the nucleation and growth pathways and highlights the importance of the synthesis conditions in tailoring opto-electronic behaviour of emissive particles.

According to Fig. 8b and c, the white emitting LCD backlight structures depict different results owing to the use of different particle sizes in the green-emissive films. For instance, the backlight structures constructed from the films with the BPPT0 exhibited ‘colder’ white emission, as evident by the higher colour coordinated temperature (CCT) of ∼5317 K, in comparison to ‘warmer’ white light emitted by the backlight structures constructed from the films with the SPPT1 (CCT ∼4317 K). Generally, a warmer white light (lower CTT) is usually preferred due to lesser amount of biological damages caused when perceived by human eyes such as reduced cell viability, cell cycle arrest, severe DNA damage etc.¹⁰⁶ All these give an edge to the SPPT1 over the BPPT0.

The colour gamut coverage areas of white light emitted by the backlight structures were calculated using Commission Internationale de l’Eclairage (CIE) colour coordinates. The backlight light structure constructed from the films with the SPPT1 spans relatively a higher colour gamut area of ∼123% of NTSC 1953 and ∼174% of sRGB (Fig. 12b), in comparison to a colour gamut coverage of ∼112% of NTSC 1953 and ∼158% of sRGB (Fig. 12a) for the backlight light structure constructed from the films with the BPPT0. Such a result not only highlights the dependence of gamut coverage area with the particle size of films employed to construct LCD backlight structures but also marks the potential of the SPPT1 and BPPT0 synthesized via the water: ODE system for applications in LCD display.

Fig. 12 CIE colour gamut coverage of the white emitting backlight structures constructed from the films with (a) BPPT0 and (b) SPPT1.

Table 1 compares the optical properties of BPPT0 and SPPT1-derived films with representative literature reports on CsPbBr₃ quantum dots, nanocrystals, composite systems, and thin films fabricated using spin-coating, vacuum deposition, thermal evaporation, and single-source vapor deposition methods. Notably, the present aqueous interface-directed synthesis yields films exhibiting narrow PL peaks at ∼534 nm (BPPT0) and ∼524 nm (SPPT1), with longer radiative decay lifetimes of up to ∼100.21 ns (BPPT0) and ∼54.11 ns (SPPT1) and significant thermal activation energies (∼100.21 meV for BPPT0 and ∼54.11 meV for SPPT1), which are comparable to those reported for conventionally processed CsPbBr₃ films. Furthermore, a comparison among different green emitting films,^107–117 utilized in fabrication of white LCD backlight structures, is made in terms of the CIE colour gamut coverage area with respect to NTSC standard and is shown correspondingly in Fig. 13. It is clearly evident that the films prepared with SPPT1 in this work outperform other reported works while the films prepared with BPPT0 remain competitive relatively, making them excellent candidates for next generation display applications.

Table 1 Comparison of optical properties of CsPbBr₃-based emitters reported in literature and in this work

Study	Material/Processing	PL peak (nm)	FWHM (nm)	Radiative decay lifetime (τ3) (ns)	Thermal activation energy (Ea) (meV)	Applications
Qaid, et al.^118	CsPbBr3 quantum dots (QDs) and thin films/spin coating	515	15	14.9	—	Lasers with low energy density threshold of amplified spontaneous emission (22.25 µJ cm^−2)
Zalrhi, et al.^57	CsPbBr3 films/spin coating	∼539	20.66	—	60.50	—
He, et al.^119	CsPbBr3 nanocrystals (NCs) and CsPbBr3-Cs4PbBr6 composite NCs/one-pot method	519 and 515	36 and 33	—	—	LEDs with luminous efficiency of 45 lmW^−1 at a 10 mA driving current
Bai, et al.^7	CsPbBr3 films/thermal co-evaporation	517	18	0.79 (average)	—	LEDs with maximum luminance of ∼252 cdm^−2
Yeh, et al.^58	CsPbBr3 films/single-source vapor deposition process	533	<18	—	—	LEDs with maximum luminance of ∼14 079 cd m^−2
Xiang, et al.^46	CsPbBr3 films/vacuum deposition	525	17	—	25.97 ± 2.99	LEDs with maximum luminance of 51 933 cd m^−2 and wide gamut coverage area of ∼114% of NTSC standard
Huang, et al.^103	PMMA-CsPbBr3 films/mechanochemical-ultrasonic method	538	∼29	46.99	217.48	—
Our work	BPPT0-derived film	534	26.5	39.85	100.21	LCD backlight structures with wide gamut coverage of ∼112% of NTSC 1953 standards
Our work	SPPT1-derived film	524	20.9	17.48	54.11	LCD backlight structures with wide gamut coverage of ∼123% of NTSC 1953 standards

---

Fig. 13 Relative NTSC gamut coverage areas of LCD backlight structures with different green emitting films.

5 Conclusions

In summary, we have synthesized multi-sized CsPbBr₃-based particles using the biphasic water: ODE system. The ODE serves not just as a non-polar passive solvent, but also as a kinetic barrier that enables interface-confined nucleation and formation of smaller, relatively phase-purer, and more luminescent CsPbBr₃-based particles. Such a system allows tunability across phase composition (CsPbBr₃vs. Cs₄PbBr₆), particle size (BPPTs vs. SPPTs), and opto-electronic behavior (PL, thermal stability, decay behavior, CCT and gamut coverage area of LCD backlights etc.), by simply varying the amount of ODE in the water: ODE system. It offers a versatile mechanism for designing next-generation perovskite materials with controlled opto-electronic behavior.

Conflicts of interest

The authors declare no competing financial interest.

Data availability

All possible experimental and analysed results have been included in this manuscript explicitly. No other new data has been generated by any further experiments/analyzation.

Supplementary information (SI) is available. See DOI: https://doi.org/10.1039/d5ta07491f.

Acknowledgements

FY is grateful for the support by the NSF through the CBET-2018411 monitored by Dr Nora F Savage.

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