In the past decade, colloidal lead halide perovskite nanocrystals (NCs) have aroused widespread interests due to their excellent photoelectric properties and potential application in light-emitting devices (LEDs) [1-3]. However, the instability in air and aqueous medium as well as lead toxicity issues hindered their employments [4-6]. Therefore, finding a lead-free alternative with optical properties comparable to the compounds has become the focus of current research. One of the most effective strategies was to replace the lead with divalent cations such as Sn²⁺ or Ge²⁺. Unfortunately, the related materials exhibited poor crystal structural stability which was attributed to the facile oxidation to Sn⁴⁺ or Ge⁴⁺ in air, bringing about defects to cause photoluminescence (PL) quenching [7]. Another popular alternative approach was to substitute the two Pb²⁺ ions with a pair of singly-charged and triply-charged metal ions to form two kinds of octahedron, which led to a A₂M'M''X₆ (A: monovalent cation, X: halogen anion, M′: + 1 cation, M'': + 3 cation) three dimensional (3D) double perovskite structure [8, 9]. The lead-free double perovskite NCs were expected to be excellent candidate in photoelectric field because of their 3D structure, eco-friendly and stability [10].

Currently, different colloidal double perovskite NCs such as Cs₂AgBiX₆, Cs₂AgInX₆, Cs₂AgSbX₆, Cs₂NaBiCl₆ and Cs₂NaInCl₆ have been synthesized via various methods. One of the mostly used approaches was hot-injection in which one precursor was injected into another solution under the condition of N₂ environment and high temperature. Another strategy was the ligand-assisted reprecipitation (LARP) route, where ions in the polar solvent was immediately destabilized by adding into an anti-solvent, thereby causing a nucleation explosion [11]. However, low photoluminescence quantum yield (PLQY) and ambient instability still remained to be the obstacles of wide application for double perovskite NCs, which were attributed to the indirect band gaps [12, 13]. For the purpose of improving the PLQYs of these NCs, different doping strategies have been explored such as incorporating Cu²⁺, Mn²⁺, and lanthanide ions [14, 15]. Alternative processes, mainly aimed at engineering the band gap of double perovskites, involved the synthesis of quinary compounds by employing a combination of B⁺ or B³⁺ ions. Previous reports showed that impurity ions (e.g., Mn²⁺, Yb³⁺/Er³⁺ and Bi³⁺ ions) could be successfully introduced into Cs₂AgInCl₆ NCs and remarkable PL improvements were obtained [16, 17]. PLQY of 36.6% for Bi-alloyed Cs₂AgInCl₆ NCs was observed which was comparable to traditional lead perovskite NCs [18, 19]. The PLQY of Na and Bi alloyed Cs₂AgInCl₆ NCs was as high as 64%, and the emission color varies with its composition and size [20, 21]. In addition to Ag-based perovskite NCs, Cs₂NaBiCl₆ NCs were mainly prepared by hot-injection with very low PLQY. Besides, Cs₂NaBiCl₆ NCs showed high flexibility for different ion doping to enhance its photoelectric properties. Recently, Mn-doped Cs₂NaBiCl₆ NCs with an improved PLQY of 15% were reported [22-24]. As investigated by Yao et al., Ag-doped Cs₂NaBiCl₆ NCs exhibited a high PLQY of 20%, which has attracted great attention on the doped lead-free Cs₂NaBiCl₆ NCs [25].

Sb³⁺, with an electronic structure of ns², is an important metal ion for double perovskites owing to its close energy levels to other B³⁺ ions and unique luminescence characteristics [26]. Currently, Sb-doping strategy was successfully applied in Cs₂SnCl₆, Cs₂AgInCl₆ and Cs₂NaInCl₆. For instance, Sb³⁺ alloyed Cs₂NaInCl₆ microcrystals showed PLQYs of 78.9% and 79% respectively. Also, by the incorporation of Sb³⁺, an excellent PLQY (~37%) of the vacancy-ordered double perovskite Cs₂SnCl₆ was obtained [27]. Besides, through a gradual substitution of Sb with In in a solid solution experiment of Cs₂AgSb_1−xIn_xCl₆, the bandgap transformed from direct to indirect [28]. In this paper, Cs₂NaBiCl₆ NCs prepared by a simple LARP approach showed bright blue photoluminescence emission with a record high PLQY of 39.05%. After incorporating Sb, the PLQY increased to 46.57%, that is, for the Cs₂NaBi_0.75Sb_0.25Cl₆ stoichiometry. Moreover, the powder and the solution of Cs₂NaBi_0.75Sb_0.25Cl₆ exhibited superior stability under air which would make it promising for commercial application.

The crystal structures of Cs₂NaBi_1-xSb_xCl₆ (x = 0, 0.25, 0.5, 0.75, 1) NCs were shown in Scheme 1. The Cs₂NaBiCl₆ crystal adopted a typical 3D framework of double perovskite with Cs atoms in the center and octahedrons consisting of corner-sharing [NaCl₆]⁵⁻ and [BiCl₆]³⁻ while Sb³⁺ ions were incorporated into Cs₂NaBiCl₆ NCs and occupied Bi³⁺ sites to form Cs₂NaBi_1-xSb_xCl₆ crystals [29]. The synthesis of Cs₂NaBi_1-xSb_xCl₆ (x = 0, 0.25, 0.5, 0.75, 1) NCs were carried out by a simple low-cost LARP route as well as by tuning the molar ratios of BiCl₃ and SbCl₃ in the precursor. Details were discussed in Supporting information. Also, the influence of different amounts of oleic acid (OA) and precursors were explored for better optical properties, where 100 µL precursor solution and 3% OA were found to be the optimal condition for the reaction (Figs. S1a-d in Supporting information). As shown in Fig. 1a, the X-ray diffraction (XRD) patterns of Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ NCs were consistent with Cs₂NaBiCl₆ (PDF#77–1831), indicating that the NCs were pure and highly crystalline in the Fm 3 m space group. The ordered arrangements of Na⁺ and Bi³⁺ were evidenced by the strong intensity of the (1 1 1) peak. The (4 0 0) peak shifted toward higher angles as the amount of Sb³⁺ increased, which was attributed to the lattice contraction (Fig. S2 in Supporting information) [26]. Notably, the intensity of (1 1 1) and (3 1 1) peaks gradually decreased with higher Sb³⁺ ratios (x > 0.25), while an impure phase Cs₃Sb₂Cl₉ appeared in Sb-rich condition (Fig. 1a) [30]. The low alloying level of Sb was probably due to the larger radius difference between Na⁺ (0.102 nm) and Sb³⁺ (0.076 nm) comparing to that of Na⁺ (0.102 nm) and Bi³⁺ (0.103 nm). In order to reveal the accurate stoichiometric ratios of the metal ions (Bi³⁺ and Sb³⁺⁾ in Cs₂NaBi_0.75Sb_0.25Cl₆, the inductively coupled plasma-optical emission spectrometry (ICP-OES) measurement was conducted. The feeding molar amount of (Bi³⁺ and Sb³⁺) was equal to that of Na⁺ so that the calculated molar ratios of Bi/Na and Sb/Na could represent the actual values of 1 − x and x (Table S1 in Supporting information). The results indicated that the actual amount of Sb³⁺ was much lower than the chemical formula composition, which was common in other doped double perovskites [26, 12]. As shown in transmission electron microscopy (TEM) images, the Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ NCs were spherical-shaped (Figs. 1c and d) with average sizes of 10.03 ± 0.34 nm and 7.01 ± 0.16 nm, respectively (Figs. 1e and f). From the high-resolution TEM (HRTEM), it could be observed that the lattice distances of the NCs were measured to be 0.33 nm, corresponding to the (3 1 1) plane (insets in Figs. 1c and d). The analysis above revealed that the alloy of Sb³⁺ would not produce additional vacancies and interstitial defects. Therefore, no additional surface trapping states would occur after doping and the intrinsic optical properties were significantly enhanced [31].

In order to further clarify the presence of Sb³⁺, the X-ray photoelectron spectroscopy (XPS) technique was conducted. The full spectra revealed that Cs, Na, Bi, Cl existed in both Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ NCs (Fig. S3 in Supporting information). As shown in Fig. S3d, two peaks of Sb 3d appearing at 539 eV and 531 eV in Cs₂NaBi_0.75Sb_0.25Cl₆ sample confirmed the successful incorporation of Sb³⁺ ions in Cs₂NaBiCl₆ NCs. Furthermore, compared the high-resolution XPS spectra of Bi 4f and Na 1s of Cs₂NaBi_0.75Sb_0.25Cl₆ and Cs₂NaBiCl₆ NCs (Fig. S4 in Supporting information), the binding energy of Bi 4f shifted slightly to lower energy while Na 1s has little variations. The slight shift could be attributed to the lattice contraction caused by the introduction of Sb³⁺ ions, probing the above speculation that Sb³⁺ ions occupied Bi³⁺ sites instead of Na⁺sites [32].

Next, the steady-state optical properties of Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ NCs were investigated. Bright blue-violet emission could be seen under UV lamp (Fig. S5 in Supporting information). From the absorption spectra, it was observed that a prominent absorption peak appeared at 295 nm for Cs₂NaBiCl₆ while a broad peak at 312 nm was shown for Cs₂NaBi_0.75Sb_0.25Cl₆ NCs (Fig. 1b). The 295 nm absorption peak could be assigned to ¹S₀→³P₂ forbidden transition whereas the 312 nm peak was associated with ¹S₀→³P₁ partially allowed transition of the localized [BiCl₆]³⁻ octahedron in colloidal solution [33]. The band gap, calculated from the corresponding Tauc plot, fell at 3.04 eV and 2.92 eV which illustrated a red-shift trend (Fig. S6 in Supporting information). According to the previous research, the valence band was shifted up by the incorporation of Sb³⁺ 5s orbitals while the 5p orbitals of Sb³⁺ caused a lower conduction band, resulting in a narrower band gap [31].

To look deep into the luminescence features of Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ NCs, PL spectra and PLQYs of these NCs were studied. The highest emission intensity was observed for 25% Sb-doped Cs₂NaBiCl₆ NCs. Meanwhile, the slight decrease in luminescence intensity was clearly visible in PL spectrum for the higher Sb³⁺ contents, which was due to the concentration quenching and the presence of by-product Cs₃Sb₂Cl₉. (Fig. S7 in Supporting information) [33]. In contrary to the weak emission with a low PLQY of Cs₂NaBiCl₆ NCs in published literature, the pristine Cs₂NaBiCl₆ prepared by LARP route exhibited a PLQY up to 39.05% (Fig. S8 in Supporting information). By using hot-injection method, inhomogeneous morphologies and impurity CsCl were found under different temperature which probably led to the low PLQY. Thus, the LARP approach used in this work effectively avoided such shortcomings and the ligand (OA) passivated the surface defects of the NCs, resulting in an improved PLQY [25]. Interestingly, the PL peak position of Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ showed no variation (Fig. 2a), which indicated that the PL transition was derived from the radiative recombination of the internal energy states instead of the band-edge electrons/holes recombination [31]. It was proposed that the self-trapped excitons (STEs) played an important role in the emission process. The observable PL in Cs₂NaBiCl₆ was attributed to the dark state STEs [31]. After trace Sb³⁺ doping, the emission intensity was greatly improved while maintaining the same peak position at 388 nm, which illustrated that the dark STEs in Cs₂NaBiCl₆ transformed to bright STEs in Cs₂NaBi_0.75Sb_0.25Cl₆ [12]. The transition from dark state to bright state was probably related to two aspects. Firstly, the parity-forbidden transition could be broken by the interaction of [SbCl₆]³⁻ and [BiCl₆]³⁻, resulting in effective PL emission. Secondly, reduced lattice vibrations could increase radiative recombination pathway [12]. To prove the smaller lattice vibrations caused by Sb³⁺ doping, Raman spectra was investigated (Fig. S9 in Supporting information). It was obvious that the peak position of Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ remained the same, but the vibration intensity decreased after alloying Sb [12]. More importantly, the PLQY value of Cs₂NaBi_0.75Sb_0.25Cl₆ NCs was measured to be 46.57% (Fig. S8), suggesting a significant improvement compared to other ion alloyed Cs₂NaBiCl₆ NCs (Table S2 in Supporting information). For the purpose of better understanding the mechanism, the wavelength-dependent PL emission spectrum and PL excitation (PLE) spectrum of Cs₂NaBi_0.75Sb_0.25Cl₆ were carried out. The PL curves were almost the same at different excited wavelength which proved that the emission was not aroused from permanent defects (Fig. S10a in Supporting information) [12]. The almost same PLE profile verified the same excited state origin of these emission bands (Fig. S10b in Supporting information) [31]. As shown in Scheme S1 (Supporting information), the emission energy of STE could be calculated by the formula: E_STE = E_g - E_b - E_st - E_d where E_g represents the band gap, E_b represents the binding energy of free exciton, E_st represents the self-trapping energy, and E_d represents the lattice deformation energy [34]. Based on the strong exciton-phonon coupling existed in pure Cs₂NaBiCl₆ and Cs₂NaSbCl₆ [26], Cs₂NaBi_0.75Sb_0.25Cl₆ promoted the formation of STEs and produced small E_st and E_d. Therefore, the values of E_STE and E_g was very close which could lead to the strong blue emission [31].

In order to shed more light on the recombination dynamics of Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ NCs and clarify the implication of Sb, time-resolved PL (TR-PL) measurements using time-correlated single photon counting (TCSPC) was carried out. The two decay curves could be fitted well by a biexponential function (Fig. 2b). For Cs₂NaBiCl₆, a fast decay component (short lifetime τ₁) assigned to charge carrier trapping processes was about 2.5 ns and a slow decay component (long lifetime τ₂) associated with the band-to-band radiative recombination was 9.8 ns. For Cs₂NaBi_0.75Sb_0.25Cl₆, a long-lived lifetime of 8.7 ns and a short-lived lifetime of 1.9 ns were obtained (Table S3 in Supporting information). It was well-known that the ground state of Sb³⁺ was represented by the ¹S₀ atomic term, and four energy levels of ¹P₁, ³P₀, ³P₁ and ³P₂ constituted the excited state (sp). The ¹S₀−¹P₁ transition was allowed, and the ¹S₀−³P₁ transition was partially-allowed because of the spin-orbit coupling, while the ¹S₀−³P₂ and ¹S₀−³P₀ transitions were completely prohibited at the electric dipole transition level [35]. Upon the theory and research of Reisfeld et al., it was reasonable to speculate that there could exist two recombination routes. One was the excited state of ³P₀→¹S₀ recombination in Sb³⁺ with a faster decay rate and the other was the excited state of ³P₁→¹S₀ with a slower decay rate [35]. It meant that the incorporation of Sb³⁺ increased the possibility of radiation recombination by breaking the local symmetry so that the initial parity-forbidden transition could occur, which led to the improvement of PLQY [21]. Besides, both the percentages of τ₁ and τ₂ of Cs₂NaBiCl₆ were 50% while the percentages of τ₁ and τ₂ of Cs₂NaBi_0.75Sb_0.25Cl₆ were 49% and 51%. The similar percentages of two NCs represented the same emission mechanism which is consistent with the inference that Sb alloying caused STEs to change from dark state to bright state.

The stability of perovskite NCs was also directly relevant to the industrial applications. To evaluate the air stability of Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ NCs solutions, the PL evolution was measured with different storage time in ambient atmosphere. As illustrated in Fig. 3a, the PL intensity in first few days of two solutions increased slightly which was assigned to the elimination of dangling bonds [36]. With the increasing time placed in air, the PL intensity gradually declined owing to the formation of surface defects after air exposure [36]. It was worth noting that the Cs₂NaBi_0.75Sb_0.25Cl₆ solution retained 91% of original photoluminescence after 40 days while Cs₂NaBiCl₆ solution retained only 83% of original photoluminescence after 26 days (Fig. 3a). Furthermore, the powder of Cs₂NaBi_0.75Sb_0.25Cl₆ NCs exhibited better stability in air, maintaining the XRD pattern even after 5 months (Fig. 3b), where impurity appeared in the XRD pattern of undoped Cs₂NaBiCl₆ sample (Fig. 3c). The higher air stability of Cs₂NaBi_0.75Sb_0.25Cl₆ NCs compared to undoped NCs was attributed to the reduction of volume defects which resulted in the improvement of short-range order of the Cs₂NaBi_0.75Sb_0.25Cl₆ NCs [12]. Moreover, the water-stability of the Cs₂NaBiCl₆ and Cs₂NaBi_0.75Sb_0.25Cl₆ colloidal solution were investigated by adding 3% deionized water. It was observed that Cs₂NaBi_0.75Sb_0.25Cl₆ remained 68% of initial intensity after 8 days while the intensity of Cs₂NaBiCl₆ was dropped to merely 20% after 5 days (Fig. 3d). The better water-stability was probably due to the formation of SbOCl on the surface. The XPS peaks located at 540.0 eV and 532.9 eV were associated with Sb³⁺ 3d_3/2 and 3d_5/2. An additional peak of Sb-O at 531.7 eV appeared (Fig. S11a in Supporting information), which was attributed to the presence of Sb-O on the crystal surface due to the hydrolysis of antimony chloride. Besides, the XRD patterns of Cs₂NaBi_0.75Sb_0.25Cl₆ before and after adding water were investigated (Fig. S11b in Supporting information), where SbOCl phase was identified in Cs₂NaBi_0.75Sb_0.25Cl₆ after adding 3% water [27]. All the results demonstrated the superior stability of Cs₂NaBi_0.75Sb_0.25Cl₆ comparing to various ion doping strategies used for Cs₂NaBiCl₆ NCs in the literatures (Table S4 in Supporting information). The excellent air-stability of Cs₂NaBi_0.75Sb_0.25Cl₆ NCs laid a solid foundation of its wide application and the research results provided a new example for the future study of Sb-based double perovskite.

In summary, Cs₂NaBiCl₆ NCs were fabricated via a room-temperature LARP approach with a high PLQY of 39.05%. The strategy was extended to the synthesis of Cs₂NaBi_0.75Sb_0.25Cl₆ NCs, which showed a strong blue photoluminescence with an improved PLQY of 46.57%. Additionally, the powder and the solution of Cs₂NaBi_0.75Sb_0.25Cl₆ NCs exhibited superior stability in ambient environment. The excellent optical performance and outstanding stability of Cs₂NaBi_0.75Sb_0.25Cl₆ perovskite NCs prepared by this simple route would expand its potential application for optoelectronic devices.

The authors declare no competing financial interest.

This work was sponsored by the National Natural Science Foundation of China (Nos. 21475021, 21427807 and 21777096), the Priority Academic Program Development of Jiangsu Higher Education Institutions.

Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.cclet.2021.05.071