Promoting photoluminescence quantum yields of glass-stabilized CsPbX₃ (X = Cl, Br, I) perovskite quantum dots through fluorine doping†

Daqin Chen *^abc, Yue Liu ^d, Changbin Yang ^abc, Jiasong Zhong *^d, Su Zhou ^d, Jiangkun Chen ^abc and Hai Huang ^abc
aCollege of Physics and Energy, Fujian Normal University, Fujian Provincial Key Laboratory of Quantum Manipulation and New Energy Materials, Fuzhou, China. E-mail: dqchen@fjnu.edu.cn
^bFujian Provincial Collaborative Innovation Center for Optoelectronic Semiconductors and Efficient Devices, Xiamen, 361005, China
^cFujian Provincial Engineering Technology Research Center of Solar Energy Conversion and Energy Storage, Fuzhou, China
^dCollege of Materials & Environmental Engineering, Hangzhou Dianzi University, Hangzhou, 310018, China. E-mail: jiasongzhong@hdu.edu.cn

First published on 5th September 2019

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All-inorganic CsPbX₃ (X = Cl, Br, I) perovskite quantum dots (PQDs) have recently emerged as hot light-emitting materials because of their superior optical performances, such as a high photoluminescence quantum yield (PLQY, up to 90%), narrow full width at half maximum (FWHM, down to 12 nm), and wide color gamut.^1–6 In the past few years, they have shown unprecedented radical progress, ranging from their synthesis and structure/property optimization to practical applications in QD-LEDs, lasers, photodetectors and scintillators.^7–26 Unfortunately, colloidal CsPbX₃ PQDs generally suffer from poor long-term stability upon the impact of moisture, heat and light irradiation due to their low formation energy and ionic crystal features.^27,28 Recently, embedding CsPbX₃ PQDs in inorganic oxide glasses (PQDs@glass) via in situ nucleation/growth (crystallization) has been demonstrated to be a feasible strategy to improve their stability for the effective protecting role of the robust glass matrix.^29–31 SiO₂-, P₂O₅-, GeO₂- and B₂O₃-based oxyhalide glasses containing Cs⁺, Pb²⁺ and X⁻ elements have been prepared, with subsequent heat-treatment-induced CsPbX₃ crystallization in a glass matrix.^32–36 However, unlike wet-chemical synthesis, the nucleation and growth of PQDs in glass are generally limited by a tight glass network, which hinders the diffusion of Cs⁺, Pb²⁺ and X⁻ ions and impedes the highly efficient whole-family precipitation of CsPbX₃ PQDs. Moreover, the current PQDs@glass nanocomposites still have several shortcomings. First, the design and optimization of the glass compositions involving trial and error experimentation are complex and random. Second, it is difficult to obtain the whole-family precipitation of CsPbX₃ PQDs in glasses, especially CsPbCl₃ and CsPbI₃, leading to insufficient coverage of the color gamut. Third, PLQYs of the CsPbX₃ PQDs@glass products are still low compared with that of the corresponding colloidal PQD counterparts, especially CsPbCl₃, CsPb(Br/I)₃ and CsPbI₃, which make their practical applications in optoelectronic devices difficult. Recently, CsPbCl₃ PQDs@glass has been successfully prepared by our group, but its PLQY is too low to be measurable.³¹ The highest PLQY value for CsPbBr₃ PQDs@glass is 81%, which has been realized in boron-germanium glass; however, the PLQY value for CsPb(Br/I)₃ red PQDs in glass is only 20% and the values for CsPbCl₃ and CsPbI₃ PQDs in glass have not been reported.³⁵ Pure CsPbI₃ PQDs@glass has been fabricated by Xiang et al. and its PLQY is only 4.2%.³⁶

Herein, F⁻ dopants were introduced into oxyhalide glass to modify the network structure, which was demonstrated to be beneficial for the controllable growth of CsPbX₃ PQDs. As schematically illustrated in Fig. 1a, taking SiO₂ glass as a typical example, [SiO₄] tetrahedra are tightly connected to each other by bridging oxygen (BO) ions and the introduced F⁻ ions will partially enter the network structure by breaking the Si–O bonds to produce non-bridging oxygen (NBO, Fig. 1b), which is expected to provide enough space for ionic diffusion and thus promote the precipitation of CsPbX₃ PQDs in glass. To reduce the glass-melting temperature, B₂O₃ and ZnO were introduced into the SiO₂ glass network in the present work, together with the perovskite components of Cs₂CO₃, PbX₂ and NaX. NH₄F was selected as the F⁻ source. All these raw materials were well ground and melted at 1200 °C for 15 min to produce bulky precursor glass (PG). After heat treatment at 460–580 °C for 2 h, CsPbX₃ PQDs were expected to be crystallized inside the glass (Fig. 1c). Without the addition of the F⁻ ions, the XRD patterns of the samples prepared by heat treatments at various temperatures showed no crystalline diffraction signal (Fig. 1d). As a comparison, obvious CsPbBr₃ diffraction peaks were detected and became intensified and narrow with the elevation of the crystallization temperature owing to the growth of PQDs (Fig. 1e). The percentage of PQDs in glass was evaluated to be 10–15% based on the ratio of the integrated area of the crystalline diffraction peaks and the total XRD pattern. All the results confirm the promoting role of the F⁻ additives for CsPbX₃in situ crystallization in glass. Importantly, this strategy enabled the precipitation of the whole-family CsPbX₃ (X = Cl, Cl/Br, Br, Br/I, I) PQDs in glass (Fig. 1f), leading to bright and colorful luminescence covering the entire visible spectral region (Fig. 1g).

The high-angle annular dark-field (HAADF) scanning transmission electron microscopy (STEM) observation for three typical CsPbX₃ (X = Cl, Br, I) PQDs@glass (Fig. S1, S2,†Fig. 2a) materials showed the homogeneous distribution of PQDs in the glass matrix. The obvious contrast between CsPbX₃ PQDs (bright) and the glass matrix (dark) is distinctly discernible due to the large difference in the atomic numbers between Cs/Pb (Z = 55/82) and Si/B (Z = 14/5). The selected area electron diffraction (SAED) pattern (Fig. S3†) shows discrete polycrystalline diffraction rings assigned to cubic CsPbX₃. The high-resolution TEM (HRTEM) micrograph (Fig. S4†) confirmed the single-crystalline nature of PQDs with high crystallinity and distinctly resolved lattice fringes. As a comparison, PQDs were difficult to precipitate from the glass without F⁻ doping (Fig. S5†), verifying the F additive-promoted CsPbX₃ nucleation/growth.

A series of structural characterizations were carried out to obtain information about the glass network structure. The Fourier transform infrared (FTIR) spectra (Fig. 2b) show Si–O–Si rocking and asymmetrical vibrations at ∼435 cm⁻¹ and ∼1030 cm⁻¹, respectively, B–O–B linkage at ∼697 cm⁻¹, [BO₃] vibrational structural units at ∼1380 cm⁻¹ and B–O stretching vibrations in the [BO₃] triangles at ∼1285 cm⁻¹.^37,38 The Raman spectra (Fig. 2c) evidence the existence of the [ZnO₆] structural units at ∼265 cm⁻¹, di-borate groups at 460 cm⁻¹ and Si–O–Si bending and stretching units at 765 cm⁻¹ and 1050 cm⁻¹, respectively.^38,39 All these results indicate that the glass network consists of the [SiO₄], [BO₄] and [BO₃] units. The ¹¹B magic-angle spinning (MAS) nuclear magnetic resonance (NMR) spectra (Fig. 2d) exhibit two resonance bands at −3 ppm and 9 ppm assigned to the B³⁺ ions in the [BO₃] and [BO₄] units, respectively.⁴⁰ The ²⁹Si MAS-NMR spectrum for the glass without F doping showed a relatively intense and narrow resonance band at −98 ppm, which substantially weakened and broadened upon the introduction of the F⁻ ions (Fig. 2e). This phenomenon was attributed to the incorporation of the F⁻ ions into [SiO₄] and the destruction of the glass network structure by breaking the Si–O bonds.^41,42 As a supplement, the ¹⁹F MAS-NMR spectra (Fig. 2f) provide information on Si–F bonding, with a resonance signal at −140 ppm^41,42 for the F-doped PG and PQDs@glass nanocomposite. An extra weak resonance band at −180 ppm originated from the Zn–F bond.⁴³ Therefore, it can be concluded that the added F⁻ ions break the network structure to provide space for ionic diffusion and promote the nucleation/growth of CsPbX₃ PQDs in the glass. To demonstrate the versatility of the proposed doping mechanism, we extended the experimental study to doping with other fluorides. Similar to the case of NH₄F, the introduction of fluorides, namely, LiF, NaF, CaF₂, PbF₂, YF₃ and LuF₃ into glass could indeed promote the precipitation of CsPbX₃ NCs (Fig. S6†).

The PL spectra (Fig. 3a) evidence the tunable luminescence of CsPbX₃ PQDs from violet to deep red via the modification of the halogen types and ratios. The FWHM values were in the range of 15–52 nm and the time-resolved decays indicated their radiative lifetimes of 2–70 ns with faster emission from wider-bandgap PQDs (Fig. 3b). All these results are comparable to the cases of colloidal CsPbX₃ PQDs,^1,2 confirming the successful growth of PQDs in glass. Indeed, the X-ray diffraction (XRD) patterns verify that the precursor glass is amorphous and typical cubic CsPbX₃ (X = Cl, Br, I) diffraction peaks appear after glass crystallization (Fig. 3c). The CsPbX₃ PQDs@glass nanocomposites possessed a wide gamut of pure colors, as shown in the Commission International de L'Eclairage (CIE) chromaticity diagram (Fig. 3d); a selected triangle of bright blue (B), green (G) and red (R)-emitting PQDs (inset of Fig. 3d) covers up to ∼200% of sRGB and ∼140% of the National Television Systems Committee (NTSC) TV color standard.

In a further experiment, the influence of the crystallization temperature on the optical properties of PQDs in glass was investigated. The emission band of CsPbCl₃ PQDs shifted from 407 nm to 412 nm as the heat treatment temperature increased from 460 °C to 540 °C (Fig. S7a†). The precipitation of PQDs in glass was a typical diffusion-controlled process and elevating the crystallization temperature contributed to the growth and increase in the particle size, leading to the reduction in the bandgap energy of PQDs due to the quantum confinement effect and the subsequent red-shift in PL bands. As evidenced in Fig. S8–S10,† the diffraction peaks of PQDs become intensified and narrow with the increase in the crystallization temperature. Importantly, taking CsPb(Br/I)₃ as a typical example, increasing the crystallization temperature will not induce a shift in the diffraction peaks (Fig. S10†), indicating that the Br-to-I ratio in the precipitated CsPb(Br/I)₃ PQDs is stable. Certainly, for CsPb(Cl/Br)₃, CsPbBr₃, CsPb(Br/I)₃ and CsPbI₃ PQDs in glass, the PL bands exhibited similar variations with the elevation of the crystallization temperatures, i.e., redshifts from 463 nm to 471 nm, 503 nm to 519 nm, 563 nm to 590 nm, and 638 nm to 688 nm, respectively (Fig. S7b–S7e†). These results confirm the ability of elaborately tuning the bandgap energies of CsPbX₃ PQDs in glass via heat treatment. Notably, the crystallization temperatures (460–540 °C) for precipitating I-containing PQDs in glass should be higher than those of Cl or Br-containing ones (500–580 °C) owing to the requirement of large activation energy for the diffusion of heavy I⁻ ions in glass. Furthermore, it is worth mentioning that the present glass-protected cubic CsPbI₃ PQDs were highly stable and could not be converted into other non-luminescent phases, which is different from the case of the colloidal cubic CsPbI₃ counterpart.⁴⁴

The time-resolved decay curves show the gradual elongation of the radiative lifetime for exciton recombination with the elevation of the crystallization temperature (Fig. S7f–S7j†), indicating the significant reduction in non-radiation de-excitation for charge carriers in PQDs. High heating temperatures are beneficial for increasing the PQD size and improving their crystallinity, which will reduce the surface defects of PQDs and improve PLQYs. Fig. 3e shows the determined PLQY values for several typical CsPbX₃ PQD@glass samples. For CsPbCl₃ and CsPb(Cl/Br)₃ PQDs, the PLQYs are in the range of 3%–20%; for CsPbBr₃ PQDs, the PLQY reaches as high as 80%; for CsPb(Br/I)₃ PQDs, the PLQYs are in the range of 50%–60%. Indeed, taking CsPbBr₃ and CsPb(Br/I)₃ as the typical examples, PLQY monotonously increases with the elevation of the crystallization temperature. As far as we know, this is the first report for PLQY of CsPbCl₃ PQDs@glass and the PLQY values of CsPb(Br/I)₃ and CsPbI₃ PQDs@glass are the highest reported so far (Table S1†).

Furthermore, the long-term stabilities of the as-prepared blue, green and red PQDs@glass samples were investigated by directly immersing them in aqueous solutions for different durations. The PL spectra show that there is no obvious change in PL intensity (Fig. 4a) and PL above 90% can be retained after immersing the nanocomposite in water for 30 days (Fig. S11†). The time-resolved spectra obtained by monitoring exciton recombination indicate that their decay kinetics are not remarkably affected by the elongation of the storage time in water (Fig. S12†). As evidenced in Fig. 4b, the intense RGB emissions from the three typical CsPbX₃ (X = Cl/Br, Br, Br/I) PQDs@glass nanocomposites in water are retained over a period of 30 days. Therefore, it can be concluded that the inorganic glass host is indeed beneficial for efficiently protecting PQDs from decomposition by water. Furthermore, we demonstrated that the detrimental halogen anion exchange among different PQDs can be completely prohibited (Fig. 4c and d). Blue-emitting (460 nm) CsPbCl_2.5Br_0.5, green-emitting (520 nm) CsPbBr₃, orange-emitting (580 nm) CsPbBr₂I and red-emitting (660 nm) CsPbBr_0.5I_2.5 PQDs@glass were ground into powders and appropriately mixed in various ratios. The PL spectra show invariable emission profiles for these four kinds of PQDs with the elongation of storage times (Fig. 4c), yielding stable multi-color luminescence under UV irradiation (Fig. 4d). This result confirms that anion exchange among different glass-stabilized CsPbX₃ PQDs can be completely inhibited, which is important for their practical applications in the optoelectronic field.

As a proof-of-concept experiment, the as-prepared CsPbX₃ PQDs were demonstrated to be applicable in phosphor-converted light-emitting diodes due to their high PLQYs and superior stability. As evidenced in Fig. S13,† green CsPbBr₃ (520 nm), orange CsPbBr₂I (580 nm) and red CsPbBr_1.5I_1.5 (630 nm) PQDs@glass phosphors can be effectively excited after coupling with the commercial InGaN blue chip. Herein, two kinds of prototype lighting devices, i.e., blue-chip/CsPbBr₃/CsPbBr_1.5I_1.5 and blue-chip/CsPbBr₃/CsPbBr₂I/CsPbBr_1.5I_1.5, were constructed. Benefiting from the inhibited anion exchange, stable blue/green/red and blue/green/orange/red emissions can be detected in the electroluminescence (EL) spectra (Fig. 4e and f), yielding white-light luminescence with a tunable correlated color temperature (CCT, 2000–5800 K), color rendering index (R_a, 50–90) and luminous efficiency (LE, 60–93 lm W⁻¹). The extra introduction of orange CsPbBr₂I PQDs contributed to the optimization of R_a and CCT of devices and correspondingly, the emissive color of LED could be tuned from cold white to warm white (insets of Fig. 4e and f). Additionally, controlling the amount of mixed green/orange/red PQD phosphors in the devices enabled the emitting light to move along the black-body radiation locus (Fig. S14 and S15†). It is worth noting that the luminous efficiencies of the present white light-emitting devices (60–93 lm W⁻¹) are comparable or even higher than those of the devices based on chemically synthesized CsPbBr₃ and CsPb(Br/I)₃ colloidal PQDs (14–61 lm W⁻¹),^45,46 glass-crystallized CsPbBr₃ and CsPb(Br/I)₃ PQDs (15–61 lm W⁻¹),³⁵ CsPbBr₃ NCs and commercial K₂SiF₆:Mn⁴⁺ phosphors (63–98 lm W⁻¹),^47,48 CsPbBr₃ NCs and CaAlSiN₃:Eu²⁺ phosphors (50–60 lm W⁻¹)³³ and CsPbBr₃/Eu³⁺/Tb³⁺ co-doped glass (63 lm W⁻¹).³² This is attributed to the high PLQYs and bright emissions of the present CsPbBr₃ and CsPb(Br/I)₃ PQDs@glass nanocomposites (Fig. S16†). Importantly, with the increase in the forward bias current, green, orange and red emissions from CsPbBr₃, CsPbBr₂I and CsPbBr_1.5I_1.5 PQDs, respectively, were proportionally enhanced and the color coordinates of the device remained unchanged (Fig. S17†). The results were quite different from those of a previously reported study,³⁵ where green emission intensity increased much faster than red emission intensity due to the low PLQY of red PQDs@glass.

In summary, fluoride additives were demonstrated to promote the precipitation of whole-family CsPbX₃ PQDs inside borosilicate glass. The 3% PLQY of CsPbCl₃ PQDs@glass was reported for the first time and the currently highest PLQYs of 50–60% for the glass-stabilized CsPb(Br/I)₃ and CsPbI₃ orange/red PQDs were obtained. All these colorful CsPbX₃ PQDs@glass products showed excellent long-term stability. Specifically, no obvious loss of PL intensity was observed after immersing them in water for up to 30 days and no detrimental anion exchange occurred among different PQDs due to the effective protection of robust inorganic oxide glass. By adopting the mixed green/red or green/orange/red glass-stabilized PQD powders as color converters, cold/warm white light diodes with tunable optoelectronic parameters could be easily achieved. This work exploits a new strategy for preparing high-performance CsPbX₃ PQDs and provides an important advancement in exploring their practical applications in lighting and displays.

Conflicts of interest

The authors declare no competing financial interests.

Acknowledgements

This research was supported by the National Natural Science Foundation of China (51572065, 51972060).

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Footnote

† Electronic supplementary information (ESI) available: Table S1, Fig. S1–S17. Experimental section, extra XRD, HAADF-STEM images, HRTEM image, PL spectra, decay curves, luminescence photographs, EL spectra and CIE diagrams. See DOI: 10.1039/c9nr07307h

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