Tunable optical and scintillation properties of two-dimensional tin mixed-halide perovskites

Abstract

The controllable doping of hybrid organic–inorganic perovskites (HOIPs) has fueled strong interest in their utilization as scintillating materials. In terms of cation engineering, tin mixed-halide perovskites (TMHPs) are considered a stronger contender than their conventional hybrid perovskite counterparts. However, their inability to undergo physicochemical alteration via simultaneous A-site and X-site substitution remains unresolved to date. In this work, we tailor the organic ligand and halide mixture of TMHPs to shed some light on their optical and scintillation properties. By introducing three organic ligands, we synthesize phenylmethylammonium (PMA), phenethylammonium (PEA) and phenylpropylammonium (PPA) and retain the Br : I ratio at 3 : 1. In terms of structural order, we show that the interlayer spacing between the inorganic layers is gradually extended from 9.88 Å for (PMA)₂SnBr₃I and 10.29 Å (PEA)₂SnBr₃I to 10.06 Å for (PPA)₂SnBr₃I. Based on geometrical order, organic chain penetration and octahedral distortion angles, we propose a rational design to modulate the absorption, photoluminescence (PL), optical bandgap, and thermal quenching of TMHPs. (PMA)₂SnBr₃I exhibits the fastest decay time (τ_avg = 1.1 ns) compared to (PEA)₂SnBr₃I (τ_avg = 2.51 ns) and (PPA)₂SnBr₃I (τ_avg = 3.54 ns), indicating that TMHPs are promising candidates for scintillator applications. This finding is corroborated by density functional theory, which outlines the weakened antibonding interaction between I 5p and Sn 5s orbitals upon tailoring the organic ligand of the perovskite. Our results demonstrate the importance of cation engineering in leveraging innovative hybrid perovskites with novel responses via a rational design.

---

Introduction

In the past decades, the vast development of lead-free halide perovskites has triggered numerous interests, particularly in tin-based hybrid materials, which have been widely pursued in solar cells,^1–3 photodetectors,⁴ light-emitting diodes (LEDs),^5,6 field-effect transistors (FETs)^7,8 and scintillators.⁹ Tin-based hybrid perovskites are considered alternatives to lead-free counterparts due to their positions within the same group (IVA) in the periodic table. Owing to the similarities in their electronic configurations (ns² np² in the outer orbitals) and radii (Sn²⁺ = 1.15 Å, Pb²⁺ = 1.19 Å), cation exchange is possible due to minute crystal lattice distortion.^10,11 Pb and Sn also exhibit analogous crystal structures owing to their similar atomic radii, ion valences, and coordination types. Moreover, Sn-based perovskites have narrow bandgaps (∼1.3 eV), enabling absorption of a broader light spectrum in the near-infrared region.¹² Tin-based perovskites display similar or superior electronic and optical properties to Pb-based perovskites, such as higher charge carrier mobilities, long-lived hot carriers, high light absorption properties and small exciton binding energies. By contrast, the unstable conditions brought on by the easy oxidation of Sn²⁺ to Sn⁴⁺ and the poor film quality brought on by the quick crystallization process result in an increase in the structural defect density of the films and lead to rapid nonradiative relaxation. These conditions are generally found in three-dimensional (3D) Sn-based perovskites. Therefore, the reduced dimensionality of Sn-based halide perovskites is considered to limit their 3D structure.^13,14

Several studies have reported that two-dimensional (2D) Sn-based perovskites exhibit superior properties to their 3D counterparts. Romani et al. have demonstrated that 2D PEA₂SnBr₄ exhibits good water resistance and excellent optical properties, promoting as co-catalyst on hydrogen photogeneration.¹⁴ Wang et al.¹⁵ achieved the high brightness of 2D TEA₂SnI₄ in a PeLED device at approximately 300 cd m⁻². Diguna et al.⁹ revealed that (C₆H₅CH₂NH₃)₂SnBr₄ is a good candidate lead-free perovskite scintillator with a bandgap of 2.51 eV, a photoluminescence (PL) peak at 498.14 nm, and a 1.05 μs average lifetime. Moreover, Sn-based perovskites afford high-performance solar cells with PCE reaching up to 13.4%.¹⁶ Another lead-free bismuth-tin-based halide perovskite exhibited a fast decay time of 3.6 ns because the electron density was improved upon Sn incorporation into the APIBiBr₅ lattice.¹⁷ A-site modulation could be useful for tailoring the optical and electronic properties of 1D bismuth-tin-based halide perovskites.¹⁸ Most of the structural arrangement of Sn-based hybrid perovskites displays a 2D order, with benzylammonium chains situated at the A-site, generating sterical hindrance during crystallization,⁶ whereas halides are commonly used as the anion.

Compared to the hybrid lead-based perovskite scintillator (HLBPS), Sn-based perovskites provide a promising lead-free alternative for scintillator applications. (C₈H₁₂N)₂Pb(Br_1−xCl_x)₄ with varying A-site cations exhibits high neutron absorption, strong radiative recombination, and sub-nanometer distances between absorption sites and radiative centers, enabling a light yield of 41 000 photons per MeV and a detection pulse width of 2.97 ns. Dual-emission PL peaks at ∼420 and ∼440 nm correspond to octahedral distortion, inducing minor structural differences between the bulk and surface of the perovskite.¹⁹ Another study revealed that the scintillation properties of MAPbBr_0.08Cl_2.92 crystals at 30 K display a rise time faster than 100 ps. The RL spectra contain three peaks, collected at temperatures under 150 K, which are attributed to the free-exciton recombination emission, near-bound-exciton emission, and excitonic transition, respectively. The calculated exciton binding energy is about 23.2 meV and 28.5 meV.²⁰ However, (PEA)₂PbBr₄:Li exhibits fast response times with decay times of 11 ns (85%) and 38 ns (15%), a high light yield of approximately 23 300 photons per MeV, and a high spatial resolution of 3.3 linear pairs per millimeter at a modulation transfer function of 0.1. Li doping was implemented to reduce the radiation dose and exposure time for patients and provide better-quality radiation images during X-ray medical imaging applications.²¹

Previous studies have revealed that replacing either metal or organic ligands within the HOIP structure can affect its optoelectronic properties.^22–27 A variation in the organic chain of lead bromide perovskites affects the structure, optical bandgap, decay time, and optical properties. For example, different lengths of alkylammonium on A₂PbBr₄ (A = AA, HA, OA) afford monoclinic structures with a slightly varied bandgap energy in the range of 3–3.36 eV. The longer alkyl chain (OA₂PbBr₄) presents the longest decay time among all samples, with a value of 0.99 ns. The longest ligand results in a large Stokes shift in PL peaks.²⁸ In another study, cobalt-based hybrid perovskites with different lengths of alkylammonium display bandgap energy modulation as a function of the alkylammonium ligand length. They also exhibit a red shift in the PL peak with an extension of the organic chain.²⁹ Using mixed halides, along with tuning the ratio of halides in MA₂CuCl_xBr_4−x, shifts the absorption from visible light to near-infrared.³⁰ It turns out that the mixed halide strategy, by increasing the bromide composition in (PEA)₂SnI_xBr_4−x, successfully modifies the absorption and emission spectra.⁶ However, to the best of our knowledge, the investigation on altering the optical and scintillation properties of tin hybrid perovskites via simultaneous organic ligand and halide engineering is lacking.

In this study, we introduce three distinct conjugated cationic ligand that differ by a single methylene unit. That is, phenylmethylammonium (PMA), phenethylammonium (PEA) and phenylpropylammonium (PPA) are introduced in tin inorganic networks with high bromine/iodine ratio (Br : I = 3 : 1) to investigate their scintillation properties. Herein, we propose that the anomalous optical properties of tin mixed-halide perovskites (TMHPs) are largely governed by the interaction between the bromine-rich conditions and organic ligand. To interpret the finding, we refer to a structural consequence of the high Br/I ratio-induced anion migration. In addition, the Br atoms occupy the equatorial position of the octahedral networks as corroborated by density functional theory.

Results and discussions

Structural and chemical properties of 2D TMHPs

We investigated the impact of three structural modulations of organic ligands, namely, PMA, PEA, and PPA, with the respective chemical structures shown in Fig. 1a–c, respectively. We used FTIR spectroscopy to investigate the functional groups of TMHPs, focusing on the fingerprint regions (Fig. S1, ESI†). The prominent C–I and C–Br stretching vibrations are found at ∼600 cm⁻¹ and ∼700 cm⁻¹, respectively. The vibrational peak at ∼3100 cm⁻¹ represents the sp² C−H stretching of the phenyl ring in PEA⁺, which is attributed to the molecular interaction (PEA⋯PEA) of −CH with the side chain of PEA⁺ and π electrons on the phenyl ring of PEA⁺.³¹ The morphological states of the respective TMHPs are displayed in Fig. 1d–f, indicating the nature of the layered structure. To complement the FTIR data, we performed Raman spectrum acquisition, as shown in Fig. S2 (ESI†). The incorporation of Br and I ions into the perovskite lattice results in a decrease in the unit cell dimensions, thereby leading to a strong electrostatic interaction between neighbouring atoms.³²

Fig. 1 Organic ligand structures of (a) phenylmethylammonium (PMA), (b) phenethylammonium (PEA), and (c) phenylpropylammonium (PPA). SEM images of (d) (PMA)₂SnBr₃I, (e) (PEA)₂SnBr₃I, and (f) (PPA)₂SnBr₃I at 2500× magnification.

The TMHP structure exhibits a triclinic system and is categorized as a 2D structure, which confirms the low-wavelength region of the Raman shift, with prominent peaks for A_1g, A_2g and B_3g phonon modes. In this low-wavelength region (under 200 cm⁻¹, Fig. S2a, ESI†), the Raman shift is associated with the bending and stretching of the tin–halide octahedral networks. The phonon mode at 98.25 cm⁻¹ for (PMA)₂SnBr₃I and (PPA)₂SnBr₃I is attributed to the A_1g mode; meanwhile, it appears at 95.68 cm⁻¹ for (PEA)₂SnBr₃I. The A_2g phonon modes appear at 181.05, 181.55, and 183.87 cm⁻¹ for (PMA)₂SnBr₃I, (PEA)₂SnBr₃I, and (PPA)₂SnBr₃I, respectively. Additionally, B_3g phonon modes are detected at 168.44, 167.99 and 171.52 cm⁻¹ for (PMA)₂SnBr₃I, (PEA)₂SnBr₃I, and (PPA)₂SnBr₃I, respectively.^33,34 The peak shift in the Raman spectrum of (PEA)₂SnBr₃I is mainly due to the high interaction of the organic cation PEA⁺ with the inorganic SnBr₃I framework. We believe that this particular mode is sensitive to hydrogen bonding, leading to a moderate control over the interaction strength between the organic cation PEA⁺ and surrounding halogen atoms.³⁵ The higher Raman mode interpretation can be found in Fig. S2b (ESI†), and further discussion using geometric assessments is presented in the following section.

The representative crystal system of TMHPs is depicted in Fig. 2a. TMHPs are confirmed to structurally retain their 2D nature judicially by the XRD patterns, as depicted in Fig. S3 (ESI†).^24–26 Based on the XRD data, the prominent peaks shifted to the lower 2θ diffraction angle due to the increasing of carbon chains length, indicating an increase in the interplanar spacing.^6,36 Wright et al. reported that the structure of perovskites with bromine and iodine is well-mixed, where bromine atoms prefer to occupy the equatorial sites of the B–X octahedra rather than the axial positions in the inorganic layer. Additionally, the organic layers of perovskites exhibit opposing alignment directions, based on the halide composition.³⁷ In the present study, we found that Br atoms also share a similar occupancy configuration. Meanwhile, each phenyl ring of organic layers is arranged in the same direction, where the neighbouring organic layers share an alternating order in the opposite direction, as depicted in Fig. 2.

Fig. 2 Crystal lattice system structure for (a) (PMA)₂SnBr₃I with the calculated Sn1–Br6–Sn1 distortion angle for (b) (PMA)₂SnBr₃I, (c) (PEA)₂SnBr₃I, and (d) (PPA)₂SnBr₃I, layer spacing, organic chain spacing, organic chain penetration, and organic chain distance (details provided in Table S8, ESI†). H-bond spacings to the nearest Br for (e) (PMA)₂SnBr₃I, (f) (PEA)₂SnBr₃I, and (g) (PPA)₂SnBr₃I. The geometric feature was visualized using the VESTA software package.³⁸

XPS analysis was performed to estimate the surface composition of the different TMHPs. The deconvolution process revealed that the Sn 3d XPS spectra (Fig. S5, ESI†) consist of two main peaks, including Sn²⁺ located at ∼494.60 eV for 3d_3/2 and ∼486.70 eV for 3d_5/2 and Sn⁴⁺ at ∼495.30 eV for 3d_3/2 and ∼486.20 eV for 3d_5/2 (Table S2, ESI†).^39,40 The occurrence of Sn⁴⁺ peaks is due to the unintentional oxidation of Sn²⁺ to Sn⁴⁺ in the perovskite host lattice.^15,41,42 However, the percentage atomic concentration of Sn²⁺ is more dominant than Sn⁴⁺ (Sn⁴⁺/Sn total ratio ∼ 0.3, Table S3, ESI†), indicating Sn valence states in TMHPs is more stable compared to other studies on tin halide perovskites. The less-oxidized tin perovskite can be achieved by using H₃PO₂ for synthesis and mixing the halide contents.^43,44 This result implies that the advantage of using H₃PO₂ in this study is that it may reduce the oxidation state of Sn. The core level of Br 3d (Table S5, ESI†) consists of Br 3d_3/2 at ∼69.3 eV and 3d_5/2 at ∼68.5 eV, arising from the spin–orbit coupling contribution. In the case of I 3d (Table S6, ESI†), the studied samples share the same chemical speciation, in which I 3d_3/2 and I 3d_5/2 are located at around 630 eV and 618 eV, respectively.

In general, the theoretical atomic ratio of Sn/Br + I is expected to be 1 : 4. We confirmed that the atomic ratio of Sn/(Br + I) is ∼1 : 2.5 for (PMA)₂SnBr₃I and (PEA)₂SnBr₃I, while for (PPA)₂SnBr₃I, it is 1 : 9 (Table S7, ESI†). Additionally, the desired 3 : 1 ratio between Br and I was the most challenging to achieve in this study because of the imperfection of the crystallization process, which is also corroborated by the XRD data. We obtained (PMA)₂SnBr₃I has an excess of the Br component while it is more insufficient on (PPA)₂SnBr₃I. By contrast, (PEA)₂SnBr₃I is close to the ideal ratio of Br to I (3 : 1); hence, its optical properties will be the most valuable among the samples.

To extend the structural investigation, the calculated CIFs were used to determine the correlation between the crystal structure and optical properties of the studied TMHPs. Based on the calculated CIF data, (PMA)₂SnBr₃I, (PEA)₂SnBr₃I, and (PPA)₂SnBr₃I share triclinic crystal systems with a space group of P1, as detailed in Table S1 (ESI†). We note that the organic ligand of TMHPs could modify the bandgap. However, this has a lower impact on the optical properties than its inorganic metal–halide interaction. First, we focus on the interlayer spacing between the neighbouring inorganic networks of TMHPs. The approach is to introduce organic cations with different chain lengths to alter cation rigidity, affecting the interlayer spacing of the inorganic layers. This, in turn, generates geometric distortion among inorganic atoms and then further influences the optical properties.²⁸ The incorporation of a longer organic chain results in an extension of the interlayer spacing between the inorganic layers: 9.88 Å for (PMA)₂SnBr₃I, 10.29 Å for (PEA)₂SnBr₃I, and 10.06 Å for (PPA)₂SnBr₃I (details provided in Table S8, ESI†). We believe that a short interlayer spacing lowers the energy barrier of charge transfer between inorganic layers.⁴⁵ Similarly, the chain spacing between layers and the neighbouring chain distance are also affected by increasing organic chain length. We therefore propose that the physical properties originate from the strong interplay of the σ* interaction arising between the metal and p orbital of iodine/bromine, in agreement with the respective works of Knutson and Zhu.^46,47

We structurally define organic chain penetration (see Fig. 2a) as the distance between the NH₃⁺ component and the general plane of halogen atoms. The shallowest penetration is 1.27 Å in (PMA)₂SnBr₃I, while the deepest penetration is 1.41 Å in (PPA)₂SnBr₃I. We also calculated the Sn1–Br6–Sn1 distortion equatorial angle of two neighbouring octahedra with iodine positioned above, as depicted in Fig. 2b–d. The distortion angles of the respective TMHPs are determined to be 135.85°, 144.81°, and 162.99° for (PMA)₂SnBr₃I, (PEA)₂SnBr₃I, and (PPA)₂SnBr₃I, respectively.

These distortion angles were positively influenced by the penetration depth into the octahedron layers owing to the weakening of hydrogen bonds due to the length variation of organic ligands.^48,49 It is notable that the trend of the optical properties of TMHPs deviates from that of single-halide Pb-based perovskites, and the large distortion angle of B1–X6–B1 (X = halides, B = divalent metal cation) leads to diminished optical properties, which will be discussed later.

Optical properties of 2D TMHPs

To outline the impact of the structure in these mixed halide systems, we performed absorption and PL measurements of this TMHP series, as shown in Fig. 3. The absorption peak maxima of (PMA)₂SnBr₃I, (PEA)₂SnBr₃I, and (PPA)₂SnBr₃I appear at 440.9 nm, 360.8 nm and 372.7 nm, respectively. The absorption of TMHPs tends to blue shift as long as the organic cation length increases. This is similar to other varied ligands on hybrid perovskites that resulted in the blue shift yield increasing the band gap energy.^28,50 Elliot fittings were used to calculate the respective energy bandgap, which is associated with absorption spectra (Fig. S4, ESI†). In line with the distortion angle mentioned before, the optical bandgap of each sample was affected by changes in the organic ligand from single-carbon chain structures to three-carbon chain structures, which were 2.7 eV, 3.16 eV, and 3.14 eV for (PMA)₂SnBr₃I, (PEA)₂SnBr₃I, and (PPA)₂SnBr₃I, respectively. Here, we propose that such bandgap variations are directly influenced by structural distortion in the octahedral networks and interlayer spacings,¹³ as outlined in Fig. 2.

Fig. 3 Absorption (black lines) and PL (orange lines) steady-state spectra of (a) (PMA)₂SnBr₃I, (b) (PEA)₂SnBr₃I, (c) (PPA)₂SnBr₃I, (d) (PEA)₂SnBr₄, (e) (PEA)₂SnI₄ and (PEA)₂SnBr₄ and (f) FWHM PL distribution of TMHPs.

Furthermore, the PL peak positions of varied TMHPs show no appreciable changes compared to (PEA)₂SnBr₄ (Fig. 3d) and (PEA)₂SnI₄ (Fig. 3e). The PL spectra of TMHPs also exhibit blue-shifting trends when the organic cation length increases from ∼588.5 nm for (PMA)₂SnBr₃I (Fig. 3a) to ∼555.2 nm for (PPA)₂SnBr₃I (Fig. 3c). By contrast, (PEA)₂SnBr₃I has a longer organic chain than (PMA)₂SnBr₃I but a shorter organic chain than (PPA)₂SnBr₃I, revealing the lowest wavelength of ∼534.10 nm, as depicted in Fig. 3b. This could be related to the interlayer spacing of the TMHP structure, based on Fig. 2. (PEA)₂SnBr₃I has the highest interlayer spacing, which requires more energy to carry out the charge carrier recombination process. We note that the halide substitution of pure bromine to mixed Br : I = 3 : 1 and then to fully iodine results in a minute blue shift of the PL maxima peak.

Interestingly, (PMA)₂SnBr₃I and (PEA)₂SnBr₄ have similar bandgaps, while (PMA)₂SnBr₃I has a shorter ligand than (PEA)₂SnBr₄. It can be affected by the similar crystallinity of (PMA)₂SnBr₃I and (PEA)₂SnBr₄ because the crystallinity of the sample correlates to light absorption and charge trapping.^51,52 Meanwhile, (PEA)₂SnI₄ has the lowest bandgap energy among the samples, approximately 2.03 eV, because it is rich in iodine, the anion with the largest atomic radius.³ We further observed the Stokes shift, the difference between the maximum absorption and maximum PL, which has an important role, to analyse the realization of free excitons and self-trapped excitons (STEs) on TMHPs. A large Stokes shift can be obtained by increasing the length of the organic ligand.⁵⁰ In line with this reference, TMHPs display large Stokes shifts as the size of the organic chain increases, with 147.6 nm for (PMA)₂SnBr₃I, 173.3 nm for (PEA)₂SnBr₃I and 182.5 nm for (PPA)₂SnBr₃I. The large Stokes shift implies that the self-absorption contribution is negligible, which means that TMHPs are promising as a scintillator.⁵³ In addition, we observed the incremental changes in the FWHM of PL peaks in bulk (black) and surface (red) states, suggesting a marginal variation for such mixture halide compositions (Fig. 3f).

Furthermore, time-resolved photoluminescence (TRPL) measurements were carried out to investigate the charge carrier dynamics of TMHPs, as depicted in Fig. 4. (PEA)₂SnBr₄, used as a comparison, has the fastest average decay time (τ_avg = 0.95 ns) and for all three components (τ₁ = 0.25 ns; τ₂ = 0.83 ns; τ₃ = 3.45 ns). The first decay time (τ₁) represents the fast component of the decay time as a result of the radiative recombination of excitons. This parameter is important for studying high-performance scintillators.⁵⁰ It turns out that τ₁ and its percentage proportion in (PEA)₂SnBr₄ are hampered upon replacing bromine with iodine (Fig. 4a and e).

Fig. 4 TRPL decay of (a) (PEA)₂SnBr₄, (b) (PMA)₂SnBr₃I, (c) (PEA)₂SnBr₃I, (d) (PPA)₂SnBr₃I, (e) (PEA)₂SnI₄. (PEA)₂SnBr₄ and (PEA)₂SnI₄ were used for comparison in this study. (f) Average decay time distribution of TMHPs.

We note that the implication of such a percentage is beneficial to the recombination rate of the excitons in the inorganic layer, which substantially affects the scintillation process.²⁸ Upon complete halide substitution to iodine, the first decay time is slower than the previous case. This has a strong correlation with the fact that a large-radius atom leads to a decreased total contribution of the radiative recombination. As a result, (PMA)₂SnBr₃I has two decay components only: the radiative recombination of excitons (τ₁) and the dissociation to electron–hole pairs (τ₂), as depicted in Fig. 4b. Meanwhile, the other samples (PEA)₂SnBr₃I and (PPA)₂SnBr₃I exhibit three decay components (Fig. 4c and d, respectively), similar to (PEA)₂SnBr₄ (Fig. 4a) and (PEA)₂SnI₄ (Fig. 4e). The third decay component (τ₃) represents the recombination of trapped electrons and holes created by Frenkel or Schottky defects.⁵³ The summary of the average decay time is presented in Fig. 4f.

The sample with the longest organic chain ((PPA)₂SnBr₃I) has a slow average decay time of TMHPs (τ_avg = 3.54 ns; τ₁ = 0.55 ns with 11% contribution; τ₂ = 1.94 ns with 47% contribution and τ₃ = 6.13 ns with 42% contribution). This reveals that the τ₂ contribution of all TMHPs is more prominent than those of τ₁ and τ₃, which means self-trapped excitons dominate the decay process. However, TMHPs in this study exhibit shorter decay times than those in other studies,^6,18 and interestingly, this is faster than (C₆H₅CH₂NH₃)₂SnBr₄,⁹ indicating that it is a good candidate for scintillator applications. To provide extensive details on charge carrier trapping and detrapping processes, a continuation work using advanced spectroscopies such as transient absorption (TA), photo-Hall effect spectroscopy (PHES), thermoelectric effect spectroscopy (TEES), and time-of-flight current waveform (ToF CWF) can be used to determine the exact trap position in the bandgap.^54,55

The temperature-dependent radioluminescence (RL) spectra of the TMHPs were recorded from 10 to 350 K, as depicted in Fig. 5a–e. All the TMHPs show a thermal quenching (TQ) phenomenon in the RL spectra, with distinguishable intensities diminishing at ∼200 K (Fig. S6, ESI†). The RL intensities are negligible at room temperature, while the RL integrated intensities at 10 K give light yields between 1 and 30 ph per keV (Fig. 5f), lower than those observed in lead halide crystals.⁵⁶ The RL spectra indicate that TMHPs with extended ligand chains retain their peak intensity at higher temperatures. We observed two RL peaks on each of our TMHPs, with minor peak shifts observed as the temperature increased. The alteration of the ligand length results in a slight change in the RL peak positions. At 10 K, (PMA)₂SnBr₃I exhibits two RL peaks located at ∼484 and ∼506 nm. These peaks are blue-shifted as the ligand chain is lengthened, with the RL peaks of (PEA)₂SnBr₃I observed at ∼480 and ∼504 nm, and the RL peaks of (PPA)₂SnBr₃I observed at ∼480 and ∼502 nm.

Fig. 5 Radioluminescence of (a) (PEA)₂SnBr₄, (b) (PMA)₂SnBr₃I, (c) (PEA)₂SnBr₃I, (d) (PPA)₂SnBr₃I, (e) (PEA)₂SnI₄. (PEA)₂SnBr₄ and (PEA)₂SnI₄ were used for comparison in this study. (f) Light yields of 2D TMHPs.

In the case of (PMA)₂SnBr₃I, we noted that the second RL peak is red-shifted as the temperature is increased (∼506 nm at 10 K to ∼510 nm at 100 K). In the (PEA)₂SnBr₃I case, the first peak is red-shifted (∼480 nm at 10 K to ∼481.44 nm at 100 K), while the second peak is blue-shifted (∼504 nm at 10 K to ∼501 nm at 100 K). In (PPA)₂SnBr₃I, the first peak is red-shifted (∼480 nm at 10 K to ∼481 nm at 100 K), and the second peak is progressively changed (∼502 nm at 10 K to ∼505 nm at 100 K). We propose that these shifts can be attributed to the evolution of the charge carrier dynamics of the crystal structure caused by temperature changes, which are attributed to changes in the behaviour of carrier–phonon coupling, shifting the center of luminescence.^57,58 Herein, we observed that the thermal effects are much more intense in TMHPs with longer ligand chains. We also provide the TEM images to indicate the structural observation of fully bromine, iodine and mixed-halide (Fig. S7, ESI†). The multiple emission peaks at the cryogenic temperature can be correlated to the asymmetric doublet characteristics of the ³P₁–¹S₀ transition.⁵⁴ The low light yield on TMHPs (Fig. 5f) can be correlated to weak exciton binding, which increases the number of defects occupied by excitons, resulting in a higher contribution of the nonradiative recombination rate. This can be correlated to the high percentage of τ₂ contributions to PL decay, as shown in Fig. 5.^59,60 We note that no afterglow was detected in our samples during thermoluminescence experiments (Fig. S8, ESI†). In addition, we provide stability examination by keeping the crystals in sealed vial at room temperature with relative humidity 50–70% (Fig. S9, ESI†).

Band structures of 2D TMHPs

All perovskite samples exhibited a direct bandgap at the Brillouin zone Γ point, as depicted in Fig. 6a–d. We note that the corresponding partial density of state (PDOS) of TMHPs can be found in each of the panels in Fig. 6 (right panels). The PDOS spectra of TMHPs are quite similar at the conduction band. The organic sites (pink line) are slightly broadened, in line with the increased length of the chain from PMA to PEA and PPA. The broadening peak also appear on the valence band and has a minor total DOS adjustment (black line). Br (green line) and I (red line) states are dominantly contributing to the valence band, on the contrary, the conduction band display small spectra. Br p orbitals dominate the valence band below the Fermi level, and the replacement of the Br atom with I weakens the DOS spectra, indicating that the role of Br is replaced by I. Generally, the electronic properties of TMHPs are governed by the valence band maximum (VBM), containing of I 5p and Sn 5s orbitals, and the conduction band minimum (CBM), mostly involving Sn 5p orbitals. Because the bandgap energy of perovskites is mutually affected by the Sn–I or Sn–Br bond length variation, here, we propose that a weakening of the antibonding interactions between I 5p and Sn 5s could materialize as an increase in the Sn–I bond length.

Fig. 6 Band structures and densities of state of (a) (PMA)₂SnBr₄, (b) (PMA)₂SnBr₃I, (c) (PEA)₂SnBr₃I and (d) (PPA)₂SnBr₃I.

Moreover, the organic ligand influences the overall bandgap. As a conventional approach, one considers that when the bond length of a halide decreases from I to Br, the bandgap of the perovskite increases. Specifically, by varying halide compositions, we gain access to the heterogeneous regions of perovskites, in which polarons are localized in the lower-bandgap region, leading to enhanced local lattice strain.^3,61 For example, in this study, (PMA)₂SnBr₄ (2.96 eV, Fig. 6a) has a slightly lower bandgap than the calculated bandgap after introducing iodine to the sample (PMA)₂SnBr₃I (2.89 eV, Fig. 6b). By contrast, the bandgap energy increases upon longer chain (PEA) substitution. We believe that this is due to the high interlayer spacing of (PEA)₂SnBr₃I, which leads to fewer antibonding interactions.

Typical octahedral distortion in hybrid perovskites generally increases the bandgap value. However, in our system, we found that introducing weaker Sn–I bonds in mixed halides reduces antibonding interactions, especially at the valence band maximum (VBM). This, in turn, marginally decreases the bandgap despite the outcome of structural distortion.^62,63 Here, we rationalize that in fixed Br:I systems, ligand engineering enables fine control over lattice strain (ε = −0.4% to +1.1%) and antibonding reduction (ΔE = 0.12–0.28 eV), as quantified in Fig. 6. To clarify this interplay, we compared (PMA)₂SnBr₄ and (PMA)₂SnBr₃I using additional DFT calculations. The mixed-halide system exhibits a higher octahedral distortion index (τ = 0.38 vs. 0.12), with more bond length heterogeneity and bent Br–Sn–I angles (162.3° vs. 178.5°), suggesting that Br/I mixing introduces significant local strain. However, the extended and weakened Sn–I bonds diminish antibonding Sn 5s–I 5p hybridization, lowering the VBM by ∼0.32 eV. This electronic effect outweighs the bandgap widening typically caused by geometric distortion. To strengthen this, we revisit the observed 0.07-eV bandgap reduction (2.89 eV vs. 2.96 eV) in (PMA)₂SnBr₃I compared to its pure-bromide counterpart. Additionally, we propose that the larger-radius iodide introduces tensile strain, which further reduces the σ* orbital overlap,^64–68 a trend consistent with the recent literature.^62,63 Finally, ligand-dependent differences in interlayer spacing (e.g., 10.06 Å for PPA vs. 9.88 Å for PMA) and H-bond penetration depths modulate in-plane lattice strain and CH–halide interactions, further influencing electronic structure. These competing structural and electronic effects explain the band gap's unexpected narrowing in more distorted, mixed-halide perovskites (Table 1).

Table 1 Summary of the electronic and geometrical features of the Sn-based HOIPs

Sn-based HOIPs	E ^calcgap (eV)	E ^expgap (eV)	H-bond spacing to nearest Br (Å)	Sn–Br–Sn angle (deg)
(PMA)2SnBr4	2.96	2.12^(^9^)	3.01	135.85
(PMA)2SnBr3I	2.89	2.7	3.01	135.85
(PEA)2SnBr3I	3.44	3.16	3.09	144.81
(PPA)2SnBr3I	3.10	3.14	3.23	163.99

---

Experiments

Synthesis

(PMA)₂SnBr₃I was prepared using the following method: tin bromide (SnBr₂; 0.28 g, 1 mmol), bromic acid (HBr 48%; 5 mL) and hypophosphorous acid (H₃PO₂; 5 mL) were mixed in a glass beaker and heated in a water bath at 70 °C for 15 minutes. Then, phenylmethylammonium iodide (C₇H₉N·HI; 0.48 g, 2.2 mmol) was added to the mixture and continuously stirred for 1 h. The solution mixture was allowed to cool to room temperature, and orange/yellow microcrystals were obtained after a few days. Similar procedures were adopted for (PEA)₂SnBr₃I and (PPA)₂SnBr₃I using phenylethylammonium iodide (C₈H₁₁N·HI; 0.52 g, 22 mmol) and phenylpropylammonium iodide (C₉H₁₃N·HI; 0.55 g, 2.2 mmol), respectively.

Characterizations

X-ray diffraction experiments were conducted using Rigaku with a Cu K_α source operated at 40 kV. The 2θ range was from 5° to 80°. Then, the Rietveld refinements of diffraction patterns were analyzed using FullProf software.⁶⁹ The molecular interactions of the samples were analyzed using an ATR-FTIR Thermo Scientific Nicolet iS10 spectrometer and a Raman Thermo Scientific DXR3xi with a 532-nm excitation laser source. The surface morphologies were recorded using the SEM Zeiss EVO MA10. The TEM images were collected using the FEI Company FEI Tecnai G2 T20 X-Twin at an accelerating voltage of 200 kV. The samples were placed on a Cu grid coated with carbon, left to dry, and then inserted into the microscope chamber. The XPS spectra were measured using a SPECS XR-50 Al K_α X-ray source (excitation energy output = 1486.6 eV) with a pass energy of 30 eV and an impingement area of approximately 1 mm in diameter. Photoemission detection was carried out using a PHOIBOS 150 hemispherical energy analyzer (SPECS, GmbH). Photoluminescence (PL) measurements were performed using free-space excitation using a visible-near-infrared microscope (Nikon 20×, Nikon Corporation, Tokyo, Japan, NA = 0.40) via the epifluorescence method a room temperature. The samples were excited via a 31.25-kHz picosecond-pulsed diode laser (LDH D-C-375, Picoquant, Picoquant GmbH, Berlin, Germany; excitation wavelength = 375 nm, pulse width = 50 ps, and maximum power = 10 mW). The PL spectra were collected using an AvaSpec-HERO spectrometer (Avantes BV, Apeldoorn, The Netherlands). The emission was selected using a band filter of 425 ± 25 nm and was detected using a single-photon avalanche photodiode (APD) for time-resolved photoluminescence (TRPL) measurements. It was connected to a time-correlated single-photon counting acquisition module (HydraHarp 400, Picoquant GmbH, Berlin, Germany). The absorption spectra were collected using a UV-visible spectrophotometer. Radioluminescence measurements were obtained using the integrated setup including an Inel XRG3500 X-ray generator, a Cu anode tube, 45 kV/10 mA, an Acton Research Corporation SpectraPro-500i monochromator, a Hamamatsu R928 photomultiplier tube (PMT), and an APD Cryogenic Inc. closed-cycle helium cooler.^9,28

Density functional theory (DFT) calculations

DFT calculations were carried out under the Kohn–Sham formulation^70,71 as implemented in the Vienna ab initio simulation package (VASP).^72,73 The projector augmented wave (PAW) method was used to describe the interaction between ion cores and electrons.^74,75 The electron exchange–correlation was treated by the generalized gradient approximation (GGA) based on the Perdew–Burke–Ernzerhof (PBE) functional.⁷⁶ The rotationally invariant GGA+U approach introduced by Dudarev et al. was used with an effective Hubbard parameter U_eff of 9.0 eV for the Sn p orbital.⁷⁷ The plane wave basis sets with a cut-off energy of 500 eV were used for all calculations. The Brillouin zone with a k-point grid of 3 × 3 × 3 according to the Monkhorst–Pack scheme was used.⁷⁸ The zero-damping D3 method was adopted to account for the dispersion van der Waals correction.⁷⁹ During calculations, all atoms were allowed to fully relax. The conjugate gradient method was employed for cell optimizations, and the calculations were considered to converge when the maximum forces on all atoms were less than 0.01 eV Å⁻¹.

Conclusions

In summary, we highlighted a rational design on the electronic, optical, and scintillation properties of tunable TMHPs structure by incorporating three organic ligands with differing methylene chains. The moderate decay time (sub-10 ns) of (PMA)₂SnBr₃I, (PEA)₂SnBr₃I and (PPA)₂SnBr₃I is successfully demonstrated, where (PMA)₂SnBr₃I exhibits the best value of 1.1 ns. In addition, the thermal quenching of TMHPs can be rationalized using the minute interlayer spacing configuration of the TMHP structure, ranging from 9.88 Å [(PMA)₂SnBr₃I] and 10.29 Å [(PEA)₂SnBr₃I] to 10.06 Å [(PPA)₂SnBr₃I]. Here, we believe geometric consideration plays an important role, in which the shallowest penetration depth of the NH₃⁺ component and the least distorted inorganic networks govern the scintillation response of (PMA)₂SnBr₃I. Corroborating the findings, theoretical calculations unveil a strong foundation in regards to the weakened antibonding interactions that arise between I 5p and Sn 5s orbitals upon the tailoring of the organic ligand by keeping the Br/I ratio constant. We believe this work brings us a step closer to a promising candidate of TMHPs for high-performance scintillator applications.

Author contributions

Arramel and Muhammad D. Birowosuto conceived and supervised the project. S. Hartati, S. Anasha, and Nelly S. Anwari synthesized the TMHPs. S. Hartati and A. Arramel wrote the original draft. Afif A. Afkauni and T. Haposan performed Raman measurements. R. Marlina performed XRD measurements. D. Kowal, M. Makowski, Marcin E. Witkowski, W. Drozdowski and Muhammad D. Birowosuto. performed optical and scintillation measurements and analysis. L. Z. performed XPS analysis. Lina J. Diguna performed the SEM measurements. Muhammad H. Mahyuddin carried out theoretical calculations S. Hartati A. Arramel, and Muhammad D. Birowosuto analysed the data coherently and wrote the paper with input from all coauthors. All authors have given their approval to the final manuscript version.

Conflicts of interest

There are no conflicts to declare.

Data availability

The data supporting this article have been included as part of the ESI.†

Acknowledgements

S. H., T. H., A. A. A., and A. A. acknowledge PT Nanotech Indonesia Global Tbk for the start-up research grant. L. J. D., M. D. B. and A. A. acknowledge the Directorate of Research, Technology, and Community Service (DRTPM) 2024, Ministry of Education, Culture, Research, and Technology, Republic of Indonesia, for the research grant through the basic research scheme. M.H.M. acknowledges a research fund from Institut Teknologi Bandung under the “Riset ITB 2025” scheme (grant no. 841/IT1.B07.1/TA.00/2025). M. D. B. acknowledges National Science Center, Poland, for the research grant OPUS-24 no. 2022/47/B/ST5/01966.

References

1. Z. Zhang, Y. Huang, J. Jin, Y. Jiang, Y. Xu, J. Zhu and D. Zhao, Angew. Chem., Int. Ed., 2023, 62, e202308093.
2. E. Aktas, N. Rajamanickam, J. Pascual, S. Hu, M. H. Aldamasy, D. Di Girolamo, W. Li, G. Nasti, E. Martínez-Ferrero, A. Wakamiya, E. Palomares and A. Abate, Commun. Mater., 2022, 3, 104.
3. M. Pitaro, E. K. Tekelenburg, S. Shao and M. A. Loi, Adv. Mater., 2022, 34, 2105844.
4. H. Liu, L. Zhu, H. Zhang, X. He, F. Yan, K. S. Wong and W. C. H. Choy, ACS Energy Lett., 2023, 8, 577–589.
5. Y.-H. Cheng, R. Moriyama, H. Ebe, K. Mizuguchi, R. Yamakado, S. Nishitsuji, T. Chiba and J. Kido, ACS Appl. Mater. Interfaces, 2022, 14, 22941–22949.
6. L. Lanzetta, J. M. Marin-Beloqui, I. Sanchez-Molina, D. Ding and S. A. Haque, ACS Energy Lett., 2017, 2, 1662–1668.
7. H. Zhu, W. Yang, Y. Reo, G. Zheng, S. Bai, A. Liu and Y. Y. Noh, Nat. Electron., 2023, 6, 650–657.
8. S. Shao, W. Talsma, M. Pitaro, J. Dong, S. Kahmann, A. J. Rommens, G. Portale and M. A. Loi, Adv. Funct. Mater., 2021, 31, 2008478.
9. L. J. Diguna, S. Kaffah, M. H. Mahyuddin, Arramel, F. Maddalena, S. A. Bakar, M. Aminah, D. Onggo, M. E. Witkowski, M. Makowski, W. Drozdowski and M. D. Birowosuto, RSC Adv., 2021, 11, 20635–20640.
10. H. Liu, Z. Zhang, W. Zuo, R. Roy, M. Li, M. M. Byranvand and M. Saliba, Adv. Energy Mater., 2023, 13, 2202209.
11. M. M. Byranvand, W. Zuo, R. Imani, M. Pazoki and M. Saliba, Chem. Sci., 2022, 13, 6766–6781.
12. Z. Zhao, F. Gu, Y. Li, W. Sun, S. Ye, H. Rao, Z. Liu, Z. Bian and C. Huang, Adv. Sci., 2017, 4, 1700204.
13. L. Hou, Y. Zhu, J. Zhu and C. Li, J. Phys. Chem. C, 2019, 123, 31279–31285.
14. L. Romani, A. Bala, V. Kumar, A. Speltini, A. Milella, F. Fracassi, A. Listorti, A. Profumo and L. Malavasi, J. Mater. Chem. C, 2020, 8, 9189–9194.
15. Z. Wang, F. Wang, B. Zhao, S. Qu, T. Hayat, A. Alsaedi, L. Sui, K. Yuan, J. Zhang, Z. Wei and Z. Tan, J. Phys. Chem. Lett., 2020, 11, 1120–1127.
16. C. Wang, Y. Zhang, F. Gu, Z. Zhao, H. Li, H. Jiang, Z. Bian and Z. Liu, Matter, 2021, 4, 709–721.
17. R. Subagyo, P. Y. D. Maulida, D. Kowal, S. Hartati, R. M. Muslimawati, Y. Zetra, L. J. Diguna, S. Akhlus, M. H. Mahyuddin, L. Zhang, C. S. Tang, C. Diao, A. T. S. Wee, M. D. Birowosuto, N. Arramel, A. Rusydi and Y. Kusumawati, ACS Appl. Mater. Interfaces, 2023, 15, 54677–54691.
18. R. Subagyo, M. Eid, N. Safitri, R. Hanifah, A. A. Afkauni, L. Zhang, M. Makowski, M. A. K. Sheikh, M. H. Mahyuddin, D. Kowal, M. E. Witkowski, W. Drozdowski, M. D. Birowosuto, A. Arramel and Y. Kusumawati, J. Mater. Chem. C, 2025, 13, 9051–9060.
19. M. Xia, G. Niu, L. Liu, R. Gao, T. Jin, P. Wan, W. Pan, X. Zhang, Z. Xie, S. Teale, Z. Cai, J. Luo, S. Zhao, H. Wu, S. Chen, Z. Zheng, Q. Xie, X. Ouyang, E. H. Sargent and J. Tang, InfoMat, 2022, 4, e12325.
20. J. Liu, D. Hei, Q. Xu, X. Tan, J. Ruan, X. Ouyang, J. Nie, K. Wei, Q. Xu and B. Sun, RSC Adv., 2021, 11, 2020–2024.
21. W. Yan, B. Duan, Z. Zhu, Y. Song, G. Song, J. Ma, B. Li and Y. Liu, Nucl. Instrum. Methods Phys. Res., Sect. B, 2024, 546, 165159.
22. Z. Zhang, S. Wang, X. Liu, Y. Chen, C. Su, Z. Tang, Y. Li and G. Xing, Small Methods, 2021, 5, 2000937.
23. W. Zhang, K. Tao, C. Ji, Z. Sun, S. Han, J. Zhang, Z. Wu and J. Luo, Inorg. Chem., 2018, 57, 4239–4243.
24. Y. Park and J.-S. Lee, J. Phys. Chem. Lett., 2022, 13, 5638–5647.
25. M. Y. Cho, S. Kim, I. S. Kim, E. S. Kim, Z. J. Wang, N. Y. Kim, S. W. Kim and J. M. Oh, Adv. Funct. Mater., 2020, 30, 1–12.
26. L. Zhang, M. He and S. Shao, Appl. Surf. Sci., 2020, 534, 147599.
27. M.-H. Tremblay, J. Bacsa, B. Zhao, F. Pulvirenti, S. Barlow and S. R. Marder, Chem. Mater., 2019, 31, 6145–6153.
28. P. Y. D. Maulida, S. Hartati, D. Kowal, L. J. Diguna, M. A. K. Sheikh, M. H. Mahyuddin, I. Mulyani, D. Onggo, F. Maddalena, A. Bachiri, M. E. Witkowski, M. Makowski, W. Drozdowski, N. Arramel and M. D. Birowosuto, ACS Appl. Energy Mater., 2023, 6, 5912–5922.
29. A. A. Afkauni, T. Haposan, B. H. Arrosyid, S. Hartati, D. Kowal, A. Rifai, R. M. Muslimawati, R. Marlina, A. Zulfi, L. J. Diguna, M. H. Mahyuddin, B. Yuliarto, D. Onggo, M. D. Birowosuto and N. Arramel, J. Phys. Chem. C, 2024, 128, 7572–7582.
30. D. Cortecchia, H. A. Dewi, J. Yin, A. Bruno, S. Chen, T. Baikie, P. P. Boix, M. Grätzel, S. Mhaisalkar, C. Soci and N. Mathews, Inorg. Chem., 2016, 55, 1044–1052.
31. Y. Reo, T. Choi, J.-Y. Go, S. Jeon, B. Lim, H. Zhu, A. Liu and Y.-Y. Noh, ACS Energy Lett., 2023, 8, 3088–3094.
32. R. G. Niemann, A. G. Kontos, D. Palles, E. I. Kamitsos, A. Kaltzoglou, F. Brivio, P. Falaras and P. J. Cameron, J. Phys. Chem. C, 2016, 120, 2509–2519.
33. A. S. Sarkar, A. Mushtaq, D. Kushavah and S. K. Pal, npj 2D Mater. Appl., 2020, 4, 1.
34. T. Zhang, C. Zhou, X. Feng, N. Dong, H. Chen, X. Chen, L. Zhang, J. Lin and J. Wang, Nat. Commun., 2022, 13, 60.
35. T. Yin, Y. Fang, X. Fan, B. Zhang, J. L. Kuo, T. J. White, G. M. Chow, J. Yan and Z. X. Shen, Chem. Mater., 2017, 29, 5974–5981.
36. S. Hartati, P. Y. D. Maulida, T. Zakly, I. Mulyani, D. Onggo, M. H. Mahyuddin, A. Noviyanto, A. Arramel and N. T. Rochman, Nano Hybrids, 2023, 40, 1–6.
37. N. E. Wright, X. Qin, J. Xu, L. L. Kelly, S. P. Harvey, M. F. Toney, V. Blum and A. D. Stiff-Roberts, Chem. Mater., 2022, 34, 3109–3122.
38. K. Momma and F. Izumi, J. Appl. Crystallogr., 2011, 44, 1272–1276.
39. M. A. Kamarudin, D. Hirotani, Z. Wang, K. Hamada, K. Nishimura, Q. Shen, T. Toyoda, S. Iikubo, T. Minemoto, K. Yoshino and S. Hayase, J. Phys. Chem. Lett., 2019, 10, 5277–5283.
40. T. Zhang, T. Nakajima, H. Cao, Q. Sun, H. Ban, H. Pan, H. Yu, Z. Zhang, X. Zhang, Y. Shen and M. Wang, ACS Appl. Mater. Interfaces, 2021, 13, 49907–49915.
41. M. Li, M. Li, M. Li, W. W. Zuo, W. W. Zuo, Y. G. Yang, M. H. Aldamasy, M. H. Aldamasy, Q. Wang, S. H. T. Cruz, S. L. Feng, M. Saliba, M. Saliba, Z. K. Wang, A. Abate and A. Abate, ACS Energy Lett., 2020, 5, 1923–1929.
42. P. Li, X. Cao, J. Li, B. Jiao, X. Hou, F. Hao, Z. Ning, Z. Bian, J. Xi, L. Ding, Z. Wu and H. Dong, Nano-Micro Lett., 2023, 15, 167.
43. S. Gupta, D. Cahen and G. Hodes, J. Phys. Chem. C, 2018, 122, 13926–13936.
44. A. Mancini, P. Quadrelli, C. Milanese, M. Patrini, G. Guizzetti and L. Malavasi, Inorg. Chem., 2015, 54, 8893–8895.
45. D. Fu, S. Wu, W. Cao, Z. Chen and X.-M. Zhang, J. Mater. Chem. C, 2022, 10, 9613–9620.
46. J. L. Knutson, J. D. Martin and D. B. Mitzi, Inorg. Chem., 2005, 44, 4699–4705.
47. M. Zhu, C. Li, B. Li, J. Zhang, Y. Sun, W. Guo, Z. Zhou, S. Pang and Y. Yan, Mater. Horiz., 2020, 7, 2208–2236.
48. X. Li, W. Ke, B. Traoré, P. Guo, I. Hadar, M. Kepenekian, J. Even, C. Katan, C. C. Stoumpos, R. D. Schaller and M. G. Kanatzidis, J. Am. Chem. Soc., 2019, 141, 12880–12890.
49. Y. Li, H. Zhou, Z. Gong, M. Xia, Y. Han, X. Sheng, T. Wang, H. Wang, H. Zhu and E. Shi, Cell Rep. Phys. Sci., 2024, 5, 102020.
50. A. Arramel, A. D. Fauzi, X. Yin, C. S. Tang, M. H. Mahyuddin, M. F. Sahdan, M. Aminah, D. Onggo, G. Shukri, C. Diao, H. Wang, M. D. Birowosuto, A. T. S. Wee and A. Rusydi, Commun. Mater., 2021, 2, 1–12.
51. D. Ghosh, A. J. Neukirch and S. Tretiak, J. Phys. Chem. Lett., 2020, 11, 2955–2964.
52. S. Toso, I. Gushchina, A. G. Oliver, L. Manna and M. Kuno, ACS Energy Lett., 2022, 7, 4242–4247.
53. F. Maddalena, L. Tjahjana, A. Xie, H. Wang, P. Coquet, W. Drozdowski, C. Dang and M. D. Birowosuto, Crystals, 2019, 9, 88.
54. B. R. C. Vale, D. Scolfaro, C. A. Sousa, A. F. V. Fonseca, L. G. Bonato, A. F. Nogueira, J. Bettini and L. A. Padilha, ACS Appl. Nano Mater., 2025, 8, 4373–4383.
55. A. Musiienko, D. R. Ceratti, J. Pipek, M. Brynza, H. Elhadidy, E. Belas, M. Betušiak, G. Delport and P. Praus, Adv. Funct. Mater., 2021, 31, 2104467.
56. M. D. Birowosuto, D. Cortecchia, W. Drozdowski, K. Brylew, W. Lachmanski, A. Bruno and C. Soci, Sci. Rep., 2016, 6, 37254.
57. S. Kahmann, O. Nazarenko, S. Shao, O. Hordiichuk, M. Kepenekian, J. Even, M. V. Kovalenko, G. R. Blake and M. A. Loi, ACS Energy Lett., 2020, 5, 2512–2519.
58. A. G. Kontos, A. Kaltzoglou, M. K. Arfanis, K. M. McCall, C. C. Stoumpos, B. W. Wessels, P. Falaras and M. G. Kanatzidis, J. Phys. Chem. C, 2018, 122, 26353–26361.
59. L.-J. Xu, H. Lin, S. Lee, C. Zhou, M. Worku, M. Chaaban, Q. He, A. Plaviak, X. Lin, B. Chen, M.-H. Du and B. Ma, Chem. Mater., 2020, 32, 4692–4698.
60. D. Han, H. Shi, W. Ming, C. Zhou, B. Ma, B. Saparov, Y.-Z. Ma, S. Chen and M.-H. Du, J. Mater. Chem. C, 2018, 6, 6398–6405.
61. T. Ghosh, M. Gupta, B. R. K. Nanda, K. Shankar and D. Pradhan, ACS Appl. Mater. Interfaces, 2023, 15, 43909–43924.
62. S. Parveen and P. K. Giri, Nanoscale Adv., 2022, 4, 995–1025.
63. H. S. Kim and N. G. Park, NPG Asia Mater., 2020, 12, 78.
64. M. Fortino, A. Mattoni and A. Pietropaolo, JPhys Mater., 2024, 7, 045009.
65. W. Li, M. Li, Y. He, J. Song, K. Guo, W. Pan and H. Wei, Adv. Mater., 2024, 36, 2309588.
66. J. Wu, S. C. Liu, Z. Li, S. Wang, D. J. Xue, Y. Lin and J. S. Hu, Natl. Sci. Rev., 2021, nwab047.
67. W. Li, Z. Chen, J. Tang and O. V. Prezhdo, Chem. Mater., 2020, 32, 4707–4715.
68. D. Ghosh, D. Acharya, L. Zhou, W. Nie, O. V. Prezhdo, S. Tretiak and A. J. Neukirch, J. Phys. Chem. Lett., 2019, 10, 5000–5007.
69. J. Rodríguez-Carvajal, Phys. B, 1993, 192, 55–69.
70. P. Hohenberg and W. Kohn, Phys. Rev., 1964, 136, B864–B871.
71. W. Kohn and L. J. Sham, Phys. Rev., 1965, 140, A1133–A1138.
72. G. Kresse and J. Furthmüller, Comput. Mater. Sci., 1996, 6, 15–50.
73. G. Kresse and J. Furthmüller, Phys. Rev. B: Condens. Matter Mater. Phys., 1996, 54, 11169–11186.
74. P. E. Blöchl, Phys. Rev. B: Condens. Matter Mater. Phys., 1994, 50, 17953–17979.
75. G. Kresse and D. Joubert, Phys. Rev. B: Condens. Matter Mater. Phys., 1999, 59, 1758–1775.
76. J. P. Perdew, K. Burke and M. Ernzerhof, Phys. Rev. Lett., 1996, 77, 3865–3868.
77. S. Dudarev and G. Botton, Phys. Rev. B: Condens. Matter Mater. Phys., 1998, 57, 1505–1509.
78. H. J. Monkhorst and J. D. Pack, Phys. Rev. B: Condens. Matter Mater. Phys., 1976, 13, 5188–5192.
79. S. Grimme, J. Antony, S. Ehrlich and H. Krieg, J. Chem. Phys., 2010, 132, 154104.

---

Footnote

† Electronic supplementary information (ESI) available. CCDC 2434581, 2434583 and 2434584. For ESI and crystallographic data in CIF or other electronic format see DOI: https://doi.org/10.1039/d5tc01768h

---