Low temperature green synthesis of red emitting Pb-free CsMnBr₃ perovskite films

Abstract

Lead halide perovskites (APbX₃) have demonstrated exceptional opto-electronic properties, but their inherent toxicity and environmental hazards hindered their practical deployment in display technologies such as liquid crystal display (LCD) backlights. Herein, we report for the first time a facile, water-mediated synthesis of red-emitting CsMnBr₃ thin films from an aqueous solution of CsBr and MnBr₂ precursors at a low temperature of ∼50 °C. Unlike traditional synthesis routes reported for synthesis of CsMnBr₃ powders or nanocrystals, relying heavily on toxic solvents, high temperatures, or inert atmospheres, the green approach utilizes water as a benign medium to facilitate the [MnBr₆] octahedral coordination assembly, yielding continuous red films with strong photoluminescence (λ: ∼644 nm, FWHM: ∼75 nm). The as-synthesized CsMnBr₃ films exhibit remarkable optical quality with an ultra-wide color gamut coverage (∼132% of NTSC 1953 and ∼186% of sRGB color standards), making them a promising alternative for traditional red phosphors in LCD backlight applications. The characterization of electrical and photo-responses reveals a negative photoconductivity under UV irradiation, attributed to the powdered microstructure and hygroscopic nature of MnBr₂ under ambient air conditions. The photo-response of the red-emissive CsMnBr₃ films exhibits a power-law dependence on high-energy irradiation under ambient conditions at ∼18 °C and a relative humidity of ∼65%, along with faster self-recovery behavior, highlighting the complex role of defect-mediated charge transport. This green, low-cost, and scalable synthesis route offers a promising pathway toward sustainable and lead-free phosphor materials for next-generation wide-color-gamut display technologies.

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1. Introduction

Metal halide perovskites have shown great promise over the past decade to revolutionize next-generation energy and optoelectronic devices such as light-emitting diodes (LEDs), and display technologies, due to their tunable band gaps, high defect tolerance, high thermal resilience, strong light absorption, high photoluminescence quantum yields (PLQY), and facile solution processability.^1–3 Among a wide family of perovskites, all-inorganic halide perovskite systems (such as CsBX₃, B = Pb, Sn, Mn, and X = Cl, Br, I) overpower hybrid organic–inorganic systems owing to their superior thermal and environmental stability while maintaining their excellent opto-electronic properties including higher color purity and higher defect tolerance, making them particularly attractive for device integration and optimization.^4,5

There are serious concerns, such as lead toxicity and their long-term environmental hazards, which have hindered further growth and inclusivity of lead-based halide perovskites (APbX₃) in high-performance photovoltaic and display technologies. Also, common synthesis approaches for producing lead halide perovskites (APbX₃), such as hot injection,⁶ LARP,⁷ ultrasonic,⁸ solvothermal⁹ and microwave assisted,¹⁰ further pose greater challenges, hindering their large-scale commercialization, due to the involvement of high temperature, high vacuum and inert environments, high cost, toxic polar and non-polar solvents of DMF, DMSO, hexane and toluene, low reproducibility, excessive external energies, and complex synthesis processes.^9,11–13 Therefore, the need to explore alternative cations such as Sn, Mn, Bi, Sb, and double perovskites, and develop eco-friendly processes has become imperative, although their stability and toxicity trade-offs remain significant.¹⁴

Manganese-based halide perovskites (e.g., CsMnCl₃ and CsMnBr₃) have attracted a significant amount of attention, given that Mn²⁺ is relatively less toxic than Pb²⁺ cations and exhibits a bright and efficient photoluminescence (PL) red emission arising from the spin- and parity-forbidden d–d (⁴T₁ to ⁶A₁) transitions within the [MnBr₆] octahedra of the CsMnBr₃ lattice.^15–17 This red PL emission is particularly valuable for display technologies, where deep-red emitters are required for wide gamut color reproduction for backlight units (BLUs) in liquid crystal displays (LCDs).¹⁸ However, conventional red phosphors such as rare-earth nitride phosphors, Mn⁴⁺ fluorides, and cadmium quantum dots, often suffer from high synthesis costs,¹⁹ limited tunability,²⁰ synthesis complexity,¹⁹ environmental impact,²¹ low color purity,²⁰ and poor thermal or moisture stability,²² whereas Mn-based perovskites such as CsMnBr₃ offer low-cost, lead-free, and tunable alternatives for these traditional phosphor materials.^15,23,24

Despite these lucrative advantages, synthesizing high-quality red-emitting CsMnBr₃ films remains challenging to date as most synthesis methods typically involve solid-state synthesis, microwave-assisted synthesis, or colloidal hot-injection routes, which predominantly yield red-emitting powders or nanocrystals rather than continuous red-emitting thin films suitable for device integration.^17,25–27 Moreover, these approaches typically involve toxic solvents, surfactants, elevated temperatures, and inert environments, which limit the scalability and compatibility with low-cost device fabrication.^17,25–27 For instance, Almutlaq et al.¹⁵ reported a hot-injection method to produce red-emitting CsMnBr₃ nanocrystals, requiring high temperature (∼170 °C), an inert environment, and toxic solvents such as octadecene (ODE) and hexane, thereby raising significant environmental concerns. Interestingly, Sahu et al.²⁷ employed a microwave-assisted approach to synthesize red-emitting CsMnBr₃ crystals. However, this approach is limited by demands of high vacuum and inert environments for the preparation of a Cs-oleate precursor, complicated synthesis conditions, and a non-stoichiometric excess of MnBr₂ (an optimized molar ratio of Cs₂CO₃ : MnBr₂ was 1 : 7). Gao et al.¹⁶ reported a solid-state synthesis of CsMnBr₃ powders exhibiting bright red emission at ∼650 nm, but is challenged by non-uniform film formation and poor substrate adhesion issues.

Recent advances in halide-based luminescent materials further highlight the importance of film or composite morphologies for functional opto-electronic applications.^28,29 For instance, Cai et al.²⁹ demonstrated the potential of CsPbBr₃/PS composites for ultrawide-range wearable temperature sensors, while Cai et al.²⁸ incorporated CsPbBr₃/PDMS nanospheres for multi-dimensional sensing and interactive displays. However, continuous emissive films of pure CsMnBr₃ remain unexplored, leaving a clear gap for device-compatible, red-emitting lead-free halide films. Therefore, fabricating a continuous, red emissive CsMnBr₃ films is crucial as unlike powders or nanocrystals, continuous CsMnBr₃ films would allow direct device integration, ensure uniform optical properties, and enable charge transport and photoconductivity characterization, which are inaccessible in non-film morphologies. Additionally, solution-processable film deposition approaches can improve scalability, reduce reliance on toxic solvents, and expand compatibility with standard thin-film optoelectronic fabrication techniques.³⁰

In this work, we report for the first time a simple, low-cost, and facile green synthesis of red-emitting CsMnBr₃ thin films at a low temperature of ∼50 °C via thermal evaporation of an aqueous solution containing CsBr and MnBr₂ precursors in a molar ratio of 1 : 2. Unlike conventional toxic solvents (e.g., ODE, hexane, and DMF/DMSO), water acts as green and low-cost medium enhancing ionic transport and facilitating the assembly of [MnBr₆] octahedral coordination, yielding red-emissive CsMnBr₃ films (λ: ∼644 nm, FWHM: ∼75 nm) directly in aqueous solutions. The as-synthesized CsMnBr₃ films exhibited an ultra-wide gamut coverage area of ∼132% of NTSC 1953 and ∼186% of sRGB color standards, thereby demonstrating their strong potential as an alternative to conventional red phosphors used in wide color gamut display applications (such as LCD backlights). This simple water-mediated synthesis route to produce red-emissive CsMnBr₃ films opens up new avenues for scalable, low-toxicity, easily processable, reproducible, and device-compatible phosphor materials, addressing both the sustainability and performance needs of next-generation display technologies.

2. Experimental details

2.1. Synthesis of CsMnBr₃ films

CsBr (99.9%, Beantown Chemical) and MnBr₂ (≥97%, Strem Chemicals Inc.) powders were mixed in a 1 : 2 molar ratio in 2 mL of DI water to obtain a precursor solution. 100 µL of this precursor solution was drop-casted on a cleaned Cu substrate, which was vacuum-dried at ∼50 °C to obtain crystallized films, as shown in Fig. 1, for further characterization.

Fig. 1 Schematic depicting the synthesis route of CsMnBr₃ films via an aqueous solution of CsBr and MnBr₂ in a 1 : 2 molar ratio.

2.2. Synthesis of CsPbBr₃-PMMA films for LCD backlights

CsPbBr₃ films used as green convertor film in white LCD backlights were prepared via the standard LARP method, where CsBr (0.4 mmol) (99.9%, Beantown Chemical) and PbBr₂ (0.4 mmol) (>98%, Strem Chemicals Inc.) were mixed in 10 mL of DMF (VWR) with 1 mL of oleic acid (Ward's Science) and 0.5 mL of oleylamine (>50%, TCI America). The resulting mixture was stirred at 20 °C overnight to form a precursor solution. 1 mL of this prepared precursor solution was quickly dripped in 10 mL of toluene (VWR) under vigorous stirring at 20 °C to form CsPbBr₃ nanocrystals. The precursor solution was added to a PMMA solution (formed via dissolving PMMA powders (Mw = 35 000) (Fisher Scientific) in toluene in 0.02 g mL⁻¹ concentration). The formed PMMA solution was coated on glass slides of 0.80 × 6.66 × 72.62 mm³ in dimensions, ultimately forming CsPbBr₃-PMMA films

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 the as-prepared CsMnBr₃ films was performed on an X-ray diffractometer (XRD) (Siemens D500) with CuKα radiation (λ = 1.5406 Å). X-ray photoelectron spectroscopy (XPS) (Thermo Scientific K-alpha X-ray photoelectron spectrometer) was utilized to analyze the chemical composition of the as-prepared CsMnBr₃ films. The Photoluminescence studies of the as-prepared CsMnBr₃ films were carried out using a spectrometer (Ocean optics, FLAME-S-VIS-NIR-ES, Ocean Optics) under a UV excitation wavelength of 365 nm. Absolute photoluminescence quantum yield (PLQY) measurements of CsMnBr₃ films were conducted on a Horiba Scientific Fluoromax Plus-C fluorometer in which a solid sample was mounted on a sample stage and placed in an integrating sphere. An excitation wavelength of 365 nm was used for PLQY measurements with a 1 nm slit width and a 0.1 s integration time. The Raleigh scattering peak of the excitation and the PL from these CsMnBr₃ films were extracted via scanning the emission in a range of 345–710 nm. A blank copper substrate without specimen was also scanned in the same scanning range of 345–710 nm for reference. Furthermore, PLQY calculations were performed using Horiba Scientific FluorEssence™ software with a factory acquired integrating sphere correction defined as PLQY = ((PB − PA)/(LA − LB)) × 100%, where (PB − PA) is the number of emitted photons, (LA − LB) is the number of absorbed photons, respectively. The lifetime decay studies were conducted using a DeltaHub™ high throughput time correlated single photon counting (TCSPC) controller. The electrical responses (I–V) of the as-prepared CsMnBr₃ films were investigated using a Keithley 2400 source-meter (Keithley, Solon, OH, USA) under white and UV irradiation.

3. Results

Fig. 2a–c depicts schematically the crystal structure of CsMnBr₃ along the a, b, and c crystallographic axes, respectively, highlighting strongly coupled one-dimensional connectivity of the [MnBr₆] octahedra for red photoluminescence of the material.^15,17,26,31Fig. 2d presents the surface morphology (SEM micrograph) of the as-prepared CsMnBr₃ films with a densely packed granular structure, suggesting successful formation of CsMnBr₃ microcrystals at ∼50 °C. The microcrystals are irregularly shaped but interconnected, which is favorable for charge transport in optoelectronic applications. The analysis of the size distribution (Fig. 2e) demonstrates a near-normal distribution with a mean size of 0.62 ± 0.17 µm. The EDS analysis, as shown in Fig. S1 (SI), confirms a 1 : 1 : 3 atomic ratio of Cs : Mn : Br in the as-prepared films, consistent with the stoichiometric CsMnBr₃.

Fig. 2 Crystal structure of CsMnBr₃ projected along (a) c-axis, (b) a-axis, and (c) b-axis, (d) a SEM image of the as-prepared CsMnBr₃ film, (e) histograms depicting the distribution of particle sizes, (f) XRD spectra of the as-prepared CsMnBr₃ film, and (g) PL spectra of the as-prepared CsMnBr₃ film with the inset showing optical images of the film under white and UV light (∼365 nm), respectively.

The XRD pattern of the as-deposited CsMnBr₃ films is presented in Fig. 2f. A set of diffraction peaks are observed around ∼26.58°, 29.95°, 43.1° and 43.36°, corresponding to the (200), (201), (301), and (103) planes of the hexagonal phase of CsMnBr₃ with a space group of P6₃/mmc (PDF #97-000-9703). Using Bragg's law (eqn (S1)³² in the SI) and the hexagonal system's d-spacing equation (eqn (S2)³³ in the SI), the lattice constants are calculated to be a ∼7.63 Å and c ∼6.97 Å, which are in good accordance with the reported results on the CsMnBr₃ structure.^16,31 However, all the observed Bragg peaks exhibit a left shift toward lower diffraction angles compared to the standard pattern, indicating lattice expansion as explained by Bragg's Law³² (eqn (S1)). Such shifts are well reported for perovskite materials.^34–39 The leftward shift with some peak broadening likely can be attributed to transient hydration or intercalation of water during crystallization^37,40 and defect-induced micro-strain. Note that moderate secondary peaks related to CuBr₂ (∼17 wt%) and MnO₂ (∼10 wt%) were also detected (Fig. 2f) in addition to the dominant phase CsMnBr₃ (∼73 wt%), likely arising from the residues of precursors. These values were estimated from the relative intensities of representative phase reflections, likely representing an upper bound for impurity contents. Although, these impurity phases may introduce localized non-radiative trap states or scattering centers at grain boundaries, potentially reducing photoluminescence quantum yield or charge-carrier mobility,^41–43 a relatively moderate fraction of such impurities suggests that the film's opto-electronic performance is likely governed by the dominant red-emissive CsMnBr₃ phase, with only slight quenching from these impurity phases, respectively.

Fig. 2g shows the PL spectrum of the CsMnBr₃ film deposited on a Cu substrate under UV excitation of ∼365 nm, depicting a broad peak (FWHM ∼75 nm) centered at ∼644 nm and exhibiting photoluminescence quantum yield (PLQY) of ∼30%, respectively. This emission arises from the characteristic d–d transitions of octahedrally coordinated Mn²⁺ ions, successfully excited under a high-energy UV source (λ ∼ 365 nm) and is consistent with the reports in the literature.^15,44,45 For instance, Almutlaq et al.¹⁵ reported a broad PL spectrum (FWHM: ∼78 nm) for colloidal CsMnBr₃ centered at ∼643 nm arising due to d-d transitions of Mn²⁺ ions in octahedral coordination from a low energy excited state (⁴T₁) to the corresponding ground state (⁶A₁). The insets in Fig. 2g display the optical images of the as-prepared CsMnBr₃ films, exhibiting bright red emission under UV irradiation (λ: ∼365 nm) and validating the effectiveness of the high energy UV light in exciting the Mn²⁺ transitions responsible for red emission.^46,47 Furthermore, the Cu substrate was explicitly chosen for deposition of CsMnBr₃ films owing to its high electrical conductivity and compatibility with our experimental setup. Although copper is known to be chemically active and may tend to react with halide species to form copper halides or other by-products,^48,49 several factors might limit such interactions under our synthesis conditions namely rapid solvent (water) evaporation during deposition, brief halide–substrate contact, and the absence of any acidic or strongly oxidative species to promote Cu halide formation. Moreover, X-ray diffraction spectra (Fig. 2f) also revealed minor peaks corresponding to CuBr₂ and MnO₂, indicating trace impurity formation, while preserving the dominant red emitting CsMnBr₃ phase. Importantly, photoluminescence measurements conducted on both copper (Cu) and glass substrates (Fig. S3) exhibited similar PL emission spectra, confirming that the observed optical properties (Fig. 2g) are intrinsic to the CsMnBr₃ phase and are not significantly affected by the utilization of the Cu substrate.

Thereafter, the as-prepared CsMnBr₃ films were also exposed to blue LED light of ∼450 nm in wavelength. The PL spectrum is presented in Fig. 3a together with the one under UV light of ∼365 nm in wavelength. The films exhibit the nearly same red PL emission with a peak centered around ∼644 nm, supporting that the observed characteristic luminescence originates from the intrinsic d–d transition of Mn²⁺ ions, i.e., the spin- and parity-forbidden transitions [excited 4T¹ state to ground 6A¹ state] within the octahedral [MnBr₆] coordination environment.^15,26,50,51 The nearly same red emission under different excitation wavelengths highlights the robust optical response of the as-prepared CsMnBr₃ films via the green route and further validates their suitability as stable red emitters in optoelectronic applications.

Fig. 3 (a) PL spectra of CsMnBr₃ films under different excitation sources, (b) CIE chromaticity diagram of red-emitting CsMnBr₃ films, (c), (e) and (g) long-term stability curves of CsMnBr₃ films stored under vacuum conditions, and (d), (f) and (h) long-term stability curves of CsMnBr₃ films exposed to ambient conditions (T: ∼17 °C, RH: ∼65%).

The photo-emission characteristics of CsMnBr₃ films were evaluated using the CIE chromaticity diagram (Fig. 3b). The CsMnBr₃ films exhibit CIE color coordinates of (∼0.60, ∼0.30), placing them within the red region of the standard CIE 1931 color space. This is well in accord with the PL emission peak at ∼644 nm. Such chromaticity positions are highly desirable for opto-electronic applications as they align closely with the red color gamut requirements of different standard color spaces such as NTSC and sRGB.^52–55

The long-term stability of CsMnBr₃ films was examined under both vacuum and ambient conditions (T: ∼17 °C, relative humidity (RH): ∼65%), as shown in Fig. 3c–h, respectively. Fig. 3c depicts the time-dependent PL spectra of the CsMnBr₃ films in a vacuum environment, monitored over a period of 12 days. A gradual decrease in PL intensity with a small blue shift in PL wavelength was observed with respect to increasing time. This is in contrast to the PL spectra observed for the CsMnBr₃ films under ambient conditions (T: ∼17 °C, RH: ∼65%), as shown in Fig. 3d, which experienced a faster degradation in PL intensity and a significant blue shift of PL wavelength and limited stability to several hours in ambient environments.

The time dependent PL intensities and PL wavelengths are extracted from Fig. 3c and d and plotted independently for both the vacuum and ambient conditions, as shown in Fig. 3e–h. The PL intensity of CsMnBr₃ films exhibits slower decay under the vacuum conditions (Fig. 3e) in contrast to rapid decay under ambient conditions, as shown in Fig. 3f, respectively. Both the decay behavior follows exponential dependence on time, as illustrated by the fitting curves. The enhanced PL stability of the as-prepared CsMnBr₃ films in the vacuum environment might be attributed to the absence of extrinsic degradation pathways, such as oxygen and moisture-induced halide oxidation, and/or hydrolysis reactions, which generally accelerate the non-radiative recombination processes for faster PL quenching over time.^56–61

The peak emission wavelength of CsMnBr₃ films experiences a small linear blue shift from ∼644 nm to ∼641 nm within 12 days under the vacuum conditions (Fig. 3g), in contrast to a rapid exponential blue shift under ambient conditions (T: ∼17 °C, RH: ∼65%) from ∼644 nm to ∼612 nm within 3 hours, as shown in Fig. 3h. This further highlights the slower structural and/or surface-related relaxation processes, such as minor lattice contraction and/or defect reconfiguration, experienced by the as-synthesized CsMnBr₃ films in the vacuum environment.^56–61

The observed unstable PL characteristics of CsMnBr₃ films under ambient conditions (Fig. 3f and h) can be attributed to an array of factors such as moisture-induced halide migration and structural lattice distortions^56–61 and inherent hygroscopic and deliquescent nature of MnBr₂ (utilized in the precursor formation), which is highly soluble in water^62–65 and hence readily absorbs atmospheric moisture, especially in low temperature and high humid environments. This accelerates the hydrolysis process of Mn–Br bonds and ultimately destabilizes the [MnBr₆] octahedra present inside the CsMnBr₃ lattice, facilitating some lattice distortions and creating corresponding non-radiative recombination centers resulting in faster PL decay profiles (Fig. 3f) observed in such ambient environments.^62–65

These experimental findings (Fig. 3c–h) highlight the strong susceptibility of CsMnBr₃ films to ambient conditions (involving moisture and oxygen) and underscores the necessity of encapsulation or passivation strategies to enable long-term device functionality of these films.^66–70 We also emphasize that these PL measurements were conducted without incorporating any surface coating or encapsulation strategies, and hence, they only reflect the unprotected, intrinsic stability of the as-fabricated CsMnBr₃ films under ambient conditions (T: ∼17 °C, RH: ∼65%). Similar degradation pathways have been extensively reported for different halide perovskites, such as MAPbI₃ and CsPbBr₃, where water ingress and subsequent hydrate formation induce structural collapse and severe luminescence quenching behavior.^71,72 However, several strategies have been previously reported for similar halide perovskite systems to enhance their long-term ambient durability.^73,74 For instance, silica (SiO₂) coatings have been widely utilized to shield perovskite nanocrystals from moisture- and oxygen-induced degradations,⁷⁵ while amidinium-based molecular passivation layers were recently employed for perovskite systems to dramatically enhance the environmental robustness, showing ∼10-fold reduction in ligand deprotonation and ∼2-fold increase in retention of the photoluminescence quantum yield under ambient conditions.⁷⁶ Overall, such encapsulation strategies can offer realistic pathways to enhance the long-term stability of the as-synthesized CsMnBr₃ films for practical optoelectronic and display applications.^73,74

The charge-carrier recombination dynamics of the as-prepared CsMnBr₃ films was probed using time-resolved photoluminescence spectroscopy (TRPL). Fig. 4 presents the temporal profile of TRPL for the CsMnBr₃ films with the PL emission peak centered around ∼644 nm. The temporal profile of TRPL is fitted using a tri-exponential decay equation given as:^77,78

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

Fig. 4 Transient PL decay curves of the as-prepared CsMnBr₃ films fitted with the tri-exponential decay equation.

here, ‘A_i’ (i = 1, 2 and 3) represents the relative amplitudes, and ‘τ_i’ represents the characteristic lifetime of the ith component, respectively.

The curve fitting reveals three distinct characteristic lifetimes of 78.58 ns, 1.58 ns and 0.41 ns, corresponding to radiative recombination, surface-mediated recombination and trap-mediated recombination processes, respectively.⁷⁸ The longer radiative carrier lifetime component of ∼78 ns for the CsMnBr₃ films, compared to previously reported decay components for different perovskite systems such as CsPbBr₃ and MAPbI₃,^4,15,79,80 indicates strong suppression of non-radiative recombination channels that are typically introduced via surface defects or trap states,⁷⁸ and points towards an effective defect passivation within these films. Additionally, the longer radiative lifetime components in halide perovskite systems, such as CsMnX₃, CsPbX₃ and MAPbX₃, are generally associated with higher internal quantum yield (QY) and improved device performance in opto-electronic applications, such as LEDs and other photovoltaics systems.^81,82 The mechanism behind such prolonged radiative decays is usually ascribed to an array of factors, including low defect density, reduced electron–phonon coupling, and possible charge carrier localization.^83,84

The temperature dependence of photoluminescence was investigated for the as-prepared CsMnBr₃ films, using the PL spectra shown in Fig. 5a, under ambient conditions and ∼365 nm UV irradiation. The temperature was slowly increased from ∼20 °C to ∼120 °C. A decrease in PL intensity with increasing temperature was observed (Fig. 5b), which can be ascribed to enhanced exciton–phonon dissociation⁸⁵ from a higher degree of scattering at elevated temperatures, resulting in alteration of charge carrier's effective masses and ultimately reducing the radiative recombination rates.⁸⁵ Such thermal quenching of photoluminescence is a common phenomenon for most luminescent materials, including perovskites of CsPbX₃^4,86,87 and CsMnX₃.^44,85,88,89 The temperature-dependent PL behavior can be approximately described using the following equation,^85,88,89

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. 5 (a) PL spectra of the as-prepared CsMnBr₃ films at different temperatures, (b) temperature dependence of PL peak intensity, and (c) temperature dependence of PL peak wavelength.

Using eqn (2) to curve-fit the data in Fig. 5b yields a thermal activation energy of ∼141 meV, which is significantly higher than some previously reported values for halide perovskite systems, including CsPbBr₃ (∼14–75 meV),^90,91 MAPbX₃ (∼6–53 meV),⁹² Cs₃BiBr₆ (∼41 meV),⁹³ and MAPb(Br_xI_1−x)₃ (∼20–38 meV).⁹⁴ A higher energy barrier reduces an overall thermal accessibility of non-radiative decay pathways, suggesting stronger confinement of the emitting centers and fewer trap-mediated recombination losses. This behavior aligns with the localized nature of Mn²⁺ d–d characteristic transitions in [MnBr₆] octahedra present inside the CsMnBr₃ lattice, making them less susceptible to phonon-mediated thermal scattering compared to the delocalized band-edge excitons in Pb-based perovskite systems.⁹⁵

Having a Br-rich environment during the synthesis of CsMnBr₃ films may facilitate surface passivation via improving surface coordination by saturating under-coordinated metal sites, increasing the thermal energy barrier for non-radiative recombination channels.^96,97 For instance, Ghimire et al.⁹⁶ demonstrated enhancement in luminescence behavior (such as PLQY of ∼98%) of CsPbBr₃ perovskite by filling Br-vacancies with NaBr, KBr and CsBr halides, resulting in a reduction of non-radiative decay rates. Similarly, Jing et al.⁹⁷ developed a Br-rich passivation layer via selective etching in mixed halide CsPb(Br_xI_1−x)₃ perovskite, enhancing the PL intensity by ∼10³ times and suppressing non-radiative decay pathways through this passivation layer.

The temperature-dependent PL spectrum of the CsMnBr₃ films shows a linear monotonous blue shift in the PL emission peak with increasing temperature in accordance with eqn (S3) in the SI⁸⁵. The decreasing rate of the PL peak wavelength with respect to temperature is ∼0.024 nm K⁻¹ from the linear regression fitting. This blue shift is usually attributed to an interplay of lattice expansion, exciton–phonon coupling, and possible modifications in self-trapped exciton (STE) states, which effectively shift the dominant PL emission to higher energy states as temperature increases.^64–67 A similar behavior has already been observed for many halide perovskite materials, including CsMnBr₃ and CsPbBr₃.

The electrical response of the as-prepared CsMnBr₃ films was evaluated under different illumination conditions, including dark, white light, and UV light (λ: ∼365 nm) at varying excitation intensities, as shown in Fig. 6a–d, respectively. Fig. 6a shows the current–voltage (I–V) characteristics of these films under white, dark and UV excitations at a scan rate of 0.55 V s⁻¹. The linear I–V response points towards ohmic contacts between electrodes and the film, enabling the extraction of conductance using eqn (S4) shown in the SI.^98,99 The conductance values for these films decrease progressively under dark (G: ∼2.06 × 10⁻⁹ S) to white light (G: ∼1.69 × 10⁻⁹ S) and to UV light (G: ∼9.6 × 10⁻¹⁰ S) illumination conditions, with the lowest conductance value (G ∼9.6 × 10⁻¹⁰ S) under the UV irradiation. This indicates the presence of negative photoconductivity (NPC)^100–102 inside the as-synthesized CsMnBr₃ films and is consistent with some recent reports on other perovskite materials, including CsPbX₃, Cs₃Bi₂Br₉, and Cs₃Bi₂Cl₉.^100–102

Fig. 6 (a) Current–voltage (I–V) characteristics of the as-prepared CsMnBr₃ films under dark, white light, and UV light (λ: ∼365 nm) at a scan rate of 0.55 V·s⁻¹, (b) I–V curves under UV light with different intensities (0.57 W cm⁻², 0.14 W cm⁻², 0.03 W cm⁻², and 0.01 W cm⁻²) at a scan rate of 0.55 V·s⁻¹, (c) I–t curves of the as-prepared CsMnBr₃ films under UV light (λ: ∼365 nm) of different intensities at a scan rate of 0.55 V·s⁻¹, and (d) I–t response under UV-assisted ON–OFF cycles at a scan rate of 0.25 V·s⁻¹, demonstrating negative photoconductivity (NPC).

For instance, Tailor et al.¹⁰¹ observed a linear I–V response for the Cs₃Bi₂Br₉ crystal, whose conductivity decreases under light exposure in comparison to the ‘dark’ state, exhibiting negative photoconductivity (NPC) due to trap of photo-generated charge carriers at charged defect states and an internal electric field opposite to the externally applied field, leading to a decrease of the net current under light illumination. Similarly, the time-resolved current–time (I–t) measurements conducted on these films show a linear I–t behavior with the change rate of current (r) (extracted via eqn (S5) in the SI) exhibiting a similar decreasing behavior from the dark (r: ∼1.21 × 10⁻⁹ A s⁻¹) to white light (r: ∼9.95 × 10⁻¹⁰ A s⁻¹) to UV light (r: ∼5.6 × 10⁻¹⁰ A s⁻¹) (Fig. S2a in the SI). The lowest change rate is ∼5.6 × 10⁻¹⁰ A s⁻¹ under UV light, consistent with the suppressed carrier transport observed in the I–V curves shown in Fig. 6a.

To further investigate the behavior of negative photoconductivity (NPC), the electrical response of the as-prepared films was observed under varying UV intensities of 0.01 W cm⁻² (UV₁), 0.03 W cm⁻² (UV₂), 0.14 W cm⁻² (UV₃), and 0.57 W cm⁻² (UV₄), as shown in Fig. 6b. Interestingly, the electrical conductance exhibits an intensity-dependent suppression (from ∼2.06 × 10⁻⁹ S under dark to ∼5.82 × 10⁻¹⁰ S under UV), with the highest UV illumination intensity yielding the lowest conductance (∼5.82 × 10⁻¹⁰ S). Correspondingly, the I–t curves under the corresponding UV illumination intensities (UV₁ to UV₄) (Fig. 6c) reveal a decreasing trend in the change rate of current (r) with increasing UV intensity, consistent with the I–V results. The observed trends are consistent with trap-mediated mechanisms,^100–103 where higher excitation intensities (such as UV irradiation) can promote carrier capture at surface defects, grain boundaries, and other photoactive sites within the micron-sized, powder-like CsMnBr₃ films, thereby reducing free carrier density and transport efficiency, leading to NPC behavior. However, we note that direct dynamic measurements, such as time-resolved photoluminescence (TRPL), temperature-dependent carrier transport, or power-dependent lifetime studies, would be required to unambiguously confirm this NPC mechanism, however, similar NPC responses, attributed to trap activation, space-charge buildup, and light-induced defect occupations, have been widely reported in other halide perovskite systems including CsPbX₃, Cs₃Bi₂Br₉, and Cs₃Bi₂Cl₉, respectively.^100–103

Fig. S2b in the SI shows the I–V response of the as-prepared CsMnBr₃ films during ON–OFF alternate cycling under UV illumination (λ: ∼365 nm) at a scan rate of 0.25 V·s⁻¹. Electric current exhibits an approximately exponential rise when transitioning to the dark (‘OFF’) state and a corresponding exponential decay under UV (‘ON’) illumination. These effects become more pronounced at higher applied voltages. Correspondingly, Fig. 6d displays the I–t response for the ON–OFF cycles, showing well-defined transient behavior. Fitting these I–t curves with a single-exponential equation (eqn (S6)¹⁰⁴ in the SI) yields a rise characteristic time (τ_rise) of ∼4.07 s and a decay characteristic time (τ_decay) of ∼2.54 s. Interestingly, despite the larger value of τ_rise than the τ_decay, the overall ‘dark’ current after each ‘OFF’ cycle recovered to progressively higher values than the previous baseline, indicating an effective and relatively faster self-recovery process, as shown in Fig. 6d. This is in contrast to other NPC materials, for instance, Tailor et al.¹⁰¹ reported slower I–t response curves with an overall decaying current for Cs₃Bi₂Br₉ during ON–OFF cycling under highly intense light exposure (100 mW cm⁻²).

The practical applicability of the as-prepared CsMnBr₃ films was demonstrated by constructing a white-emitting backlight structure from the red-emitting CsMnBr₃ and green-emitting CsPbBr₃ films with a commercial blue LED source (λ: ∼450 nm), as schematically illustrated in Fig. S4 in the SI. Upon excitation, the backlight structure exhibits distinct and well-separated PL emission peaks corresponding to blue (∼450 nm), green (∼524 nm), and red (∼644 nm) components, as shown in Fig. 7. The inset in Fig. 7 depicts an optical image of the constructed structure under dark, confirming the production of bright white light through color mixing of RGB components. Such clear spectral resolution of RGB components is highly beneficial for achieving wide color gamut coverage and superior color rendering in display applications (discussed later in this paper), and aligns with the recent drive toward perovskite-based backlighting technologies for next-generation optoelectronic devices.^105,106

Fig. 7 PL response of the as-constructed white-emitting backlight structure showing distinct blue (∼450 nm), green (∼524 nm), and red (∼644 nm) emission peaks with the inset depicting an optical image of the backlight structures under dark.

4. Discussion

In this work, CsMnBr₃ films were deposited on the Cu substrate at ∼50 °C from an aqueous solution of CsBr and MnBr₂. The XRD pattern (Fig. 2f) reveals small traces of impurities of CuBr₂ and MnO₂, suggesting the following reaction

Cu + CsBr + 2MnBr₂ + O₂ → CsMnBr₃ + CuBr₂ + MnO₂ (3)

In principle, water plays a crucial and active role in the synthesis of CsMnBr₃ instead of just serving as a passive solvent. The presence of even small amounts of water significantly enhances ionic mobility and cation–halide (Mn²⁺–Br⁻) interaction, facilitating the formation of well-defined [MnBr₆] octahedral coordination, and accelerating perovskite formation. This effect has been demonstrated by Tsvetkov et al.¹⁰⁷ in the mechano-chemical synthesis of similar halide perovskite materials, such as CsPbI₃, where water was demonstrated to strongly promote the reaction kinetics of CsPbI₃ from CsI and PbI₂ mixtures via increased mobility of the constituting ionic species. Additionally, Xiao et al.¹⁰⁸ demonstrated a nuanced role of water in crystallization of the MAPbI₃ perovskite. They reported that a suitable water content can facilitate high-quality perovskite film growth by enabling intermediate hydrate formation such as MAPbI₃·H₂O, facilitating the conversion process from precursors to perovskite structures.

In the CsMnBr₃ lattice, the Mn²⁺ ion is octahedrally coordinated with Br⁻ anions to form [MnBr₆] octahedra (Fig. 8), which is responsible for the observed red PL emission (∼644 nm) (Fig. 2g), owing to the d–d transition of Mn²⁺ cations from ⁴T₁ (⁴G) to ⁶A₁ energy states (Fig. 8) and this is further verified by exciting the lattice with two different high-energy sources (∼365 nm and ∼450 nm) (Fig. 3a) resulting in similar red PL emission (∼644 nm with FWHM of ∼7 nm), consistent with previous reports on CsMnBr₃.^46,47 Almutlaq et al.¹⁵ concluded that the red PL emission (∼643 nm) of CsMnBr₃ is indeed due to ⁴T₁ to ⁶A₁ (d–d) transitions of Mn²⁺ in [MnBr₆] octahedra after extracting a similar red PL emission (∼643 nm with FWHM of ∼78 nm) under three different excitation wavelengths (∼380 nm, ∼45 nm, and ∼540 nm). Furthermore, a relatively modest photoluminescence quantum yield (PLQY ∼30%) in our CsMnBr₃ films (particle size ∼0.6 µm), compared to CsMnBr₃ nanocrystals (PLQY ∼54%) reported by Almutlaq et al.¹⁵ may be attributed to an array of factors including strong light scattering and re-absorption arising from the micron-sized powder morphology compared to the nanocrystalline systems reported in past literature,^15,17 residual impurity phases and surface/bulk defect states, acting as nonradiative recombination centers, and partial oxidation of Mn²⁺ cations under ambient air, overall leading to a quenched photoluminescence behavior in Mn-based halide materials.¹⁰⁹

Fig. 8 Energy level diagram of octahedrally coordinated Mn²⁺ ions in CsMnBr₃, depicting d–d transitions and the photoluminescence pathway for red emission (∼644 nm).

Fig. 9 shows the temperature dependence of the band gap of the as-prepared CsMnBr₃ films, which follows the Bose–Einstein model¹¹⁰ given as:

where E₀ (T) represents the band gap at 0 K, a describes the strength of the electron–phonon interaction, and b is the characteristic temperature linked to average photon energy, respectively.

Fig. 9 Temperature dependence of photon energy for CsMnBr₃ films, fitted via the Bose–Einstein model.

The blue shift of the band gap with increasing temperature indicates higher electron–phonon activity and thermal lattice expansion at elevated temperatures, likely reducing the antibonding overlap between 3d Mn and 4p Br orbitals and ultimately widening the band gap. A similar behavior is reported for different halide perovskites, such as CsPbBr₃. Higher temperatures lead to enhanced electron–phonon scattering, resulting in lower antibonding overlap between 6s Pb and 4p Br, thereby shifting the valence band maxima (VBM) to lower energy values, and ultimately leading to a shift of PL emission peaks to higher energy values.^{86,95,111–114}

The correlations between electronic conductance (G) and irradiance (I) and between photo-resistance (ΔR_ph) and UV-light intensity can be expressed as:

G = G₀I^a and ΔR_ph = R₀I^b (5)

where ‘G₀’ and ‘R₀’ represent the amplitude factors, and ‘a’ and ‘b’ represent the power law exponents, respectively. Fig. 10a and b presents the variations of electronic conductance and photo-resistance with irradiance as well as the fitting curves obtained using the power law relations in eqn (5).

Fig. 10 Electrical and photo-response characteristics of CsMnBr₃ films under UV (∼365 nm) irradiation: (a) conductance as a function of UV irradiance, (b) photoresistance as a function of UV irradiance, (c) photocurrent–voltage characteristics, (d) change rate of current as a function of UV irradiance, (e) change rate of photocurrent as a function of irradiance, and (f) photocurrent at V_NL (onset of the non-linear region) as a function of irradiance.

The power-law dependence of electronic conductance and photo-resistance indicates that a strong UV irradiation enhances the trapping and/or recombination of charge carriers, limiting photocarrier migration, consistent with the observed NPC phenomena (Fig. 6a–d). Although direct defect characterization (such as EPR or XPS) or humidity-dependent photoconductivity studies were not conducted in present work, the observed NPC phenomena can plausibly be attributed to an array of defect-mediated charge-trapping mechanisms including the hygroscopic nature of MnBr₂ facilitating rapid moisture uptake¹¹⁵ (via formation of hydrates or other bromide complexes) and activating new deep trap states or ionic defects which can trap the charge carriers,^57,116 higher surface area or powdered morphology exhibited by our as-synthesized CsMnBr₃ films, which may enhance the recombination ‘hot spots’ for charge carriers, and UV activation of meta-stable trap states that may capture the photogenerated carriers, thereby impeding the overall carrier transport.^{100–102,117–119}

Some related studies showed that in trap-rich systems, photoconductivity often transitions from linear to sub-linear behavior depending on trap-state density. For instance, Yi et al.¹²⁰ demonstrated that the photoconductivity of hybrid and all inorganic perovskite crystals exhibited a crossover in power law dependence with respect to excitation intensities between power exponents 1 (linear) and 1/2 (sublinear). This behavior was observed in all compounds of cation type (organic or inorganic) or crystallographic phases.

The photocurrent–voltage (ΔI_ph–V) curves exhibit an initial linear region followed by a nonlinear threshold (V_NL) that shifts to higher voltage values with increasing irradiance, as shown in Fig. 10c. The slopes of the linear region of these ΔI_ph–V curves show a decaying trend with increasing excitation irradiance, as fitted via the following equation:¹²¹

ΔI_ph = cV + d (6)

where ‘ΔI_ph’ represents the photo-current, ‘c’ represents the slope of ΔI_ph–V curves, and ‘d’ represents the intercept of ΔI_ph–V curves referring to the initial photocurrent values, respectively. Such a behavior further highlights the decaying behavior of photoconductance with increasing excitation intensities, consistent with the NPC behavior of the films (Fig. 10b). Correspondingly, the time-resolved photocurrent (ΔI_ph–t) traces also mirror similar results with diminishing slopes over time with increasing excitation intensities, as shown in Fig. S2c in the SI. These trends strongly reflect trap-mediated conduction behavior for the CsMnBr₃ films, where higher biases and longer times are needed to overcome the immobilization of carriers by deep trap states.

The analysis of the change rates of current (r) and photocurrent (Δr_ph), as shown in Fig. 10d and e, demonstrates an irregular decaying trend of the change rate of current with excitation irradiance. The change rate of photocurrent (Δr_ph) follows a power-law decaying model (Δr_ph = r₀I^α)¹²⁰ with α ∼ 0.2, implying diminishing gains at higher photon flux levels. The photocurrent at V_NL (onset of non-linear region) continues to decline with excitation intensity following a power-law model (ΔI_ph = s₀I^β),¹²⁰ as shown in Fig. 10f, with β ∼ 0.19, consistent with other perovskite materials of CH₃NH₃PbBr₃ and CsPbBr₃, where photocurrent exhibited a power dependence with incident excitation intensities with a crossover between power exponents of 1 and 1/2, respectively.¹²⁰ These results highlight the trapping of photo-generated charge carriers via different trap states, leading to NPC behavior of the CsMnBr₃ films.

The as-constructed white emitting backlight structures (Fig. 7) exhibit a ‘colder’ white light with a higher color coordinated temperature (CCT) of ∼8455 K and a wider color gamut coverage area (extracted through Commission Internationale de l’Eclairage (CIE) color coordinates) spanning 132% of NTSC 1953 and ∼186% of sRGB color standards, as calculated via eqn (S7) and (S8)⁸⁵ and depicted in Fig. 11, respectively. Although the spectral overlaps between emission bands were not separately integrated, it was inherently represented in the measured combined PL spectrum for these stacked LCD backlight structures (Fig. 7). Furthermore, these wide gamut coverage values exceed those commonly reported for red-emitting phosphor films^18,85,122 (commonly used in white backlight structures), underscoring not only the superior spectral purity but also the color rendering capability of the CsMnBr₃ films. When coupled with the environmentally benign aqueous synthesis route, it further highlights the promise of these green-fabricated CsMnBr₃ films for next-generation display technologies.

Fig. 11 CIE color gamut coverage of white emitting backlight structures from the as-prepared CsMnBr₃ films (λ_P ∼ 644 nm).

Comparison among different red emitting phosphors,^123–136 utilized in white LCD backlight structures, is made in terms of the CIE color gamut coverage area with respect to the NTSC standard, as shown in Fig. 12. It is evident that the red-emitting CsMnBr₃ films outperform other red-emitting phosphors reported, making them excellent candidates for next generation display applications.

Fig. 12 Relative NTSC gamut coverage areas of white-emitting backlight structures with different green emitting films.

4. Conclusions

In summary, we have demonstrated a sustainable, facile, and low-cost approach to synthesize red-emitting CsMnBr₃ films of high quality by using water as the only solvent. This low temperature green route overcomes the limitations of conventional methods that rely on toxic solvents, high temperatures, and/or high inert atmospheres, and directly yields emissive thin films rather than powders or colloidal nanocrystals. The as-fabricated CsMnBr₃ films exhibit efficient deep-red photo-luminescence (λ ∼ 644 nm) and an exceptionally wide color gamut (spanning ∼132% of NTSC 1953 and ∼186% of sRGB color standards), underscoring their suitability as next-generation backlight materials for wide-color displays, which majorly rely on red phosphors to date. These films exhibit an intriguing NPC behavior and power-law dependent photo-responses, which provide new insights into the charge transport mechanism of Pb-free Mn-based halide perovskites. Overall, this work offers fresh insights for scalable, environment-friendly, and high-performance phosphor films for optoelectronic and display technologies.

Conflicts of interest

The authors declare no competing financial interest.

Data availability

All possible experimental and analysed results have been included in this manuscript and supplementary information (SI) explicitly. No other new data have been generated by any further experiments/analyses. Supplementary information is available. See DOI: https://doi.org/10.1039/d5ma01256b.

Acknowledgements

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

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