Highly Emissive and Stable Ge(II)‐ and Sn(II)‐Based Vacancy‐Ordered Iodide Double Perovskites

ABSTRACT

Vacancy‐ordered double perovskites (VODPs) of the A ₂ B□X ₆ type (□ = ordered B‐site vacancy) incorporating stereochemically active ns² metals are promising lead‐free self‐trapped exciton (STE) emitters. However, efficient and stable iodide VODPs with ns² metals as the lattice centers remain rare. Here, we develop N‐alkylated 1,4‐diazabicyclo[2.2.2]octane (dabco)‐derived diammonium cations Rdabco²⁺ (R = Me, Et, or Pr; mono‐alkylated or N,N′‐dialkylated), to access iodide VODPs (Rdabco)₂ B□I₆ with B = Ge(II) or Sn(II). Varying the R group and alkylation across six cations effectively tunes the steric hindrance and lattice strain, enhancing oxidation tolerance and yielding zero‐dimensional lattices of isolated [BI₆] ^{4 −} octahedra. The materials exhibit bright room‐temperature STE emission with photoluminescence quantum yields (PLQYs) up to 35.0% for Ge and 36.8% for Sn. Notably, the highest PLQY of the Ge VODP is approximately seven times that of the highest reported Ge(II) iodide compounds. Structure–photophysics correlations reveal distinct determinants: the Ge PLQYs correlate with [GeI₆] ^{4 −} octahedral distortion, whereas the Sn PLQYs reflect a more complex balance between radiative and nonradiative relaxation channels rather than a single distortion metric. These materials exhibit high ambient stability up to two months whereas the 3D perovskite analogs decompose within three days.

We report a series of new iodide vacancy‐ordered double perovskites (Rdabco)₂ B□I₆ (B = Ge^II or Sn^II), incorporating N‐alkylated dabco‐derived organic cations. By systematically varying the cation, we control the octahedral distortion and crystal symmetry, which directly tune and enhance the emission properties with record‐setting PLQY for the Ge(II) iodides. We further reveal that reduced octahedral distortion effectively suppresses nonradiative decay, while the zero‐dimensional isolated‐octahedra architecture ensures exceptional stability, establishing a versatile platform for efficient, stable, lead‐free luminescent materials.

1. Introduction

Lead‐free double perovskites, A ₂ B′B″X ₆, and their derivatives provide chemically robust and compositionally tunable halide frameworks for optoelectronic materials design [1, 2]. However, intrinsic photoluminescence (PL) is often weak because many lead‐free double perovskites possess indirect band gaps and/or weakly allowed band‐edge transitions (including symmetry‐limited transition strength), which together suppress radiative recombination [2, 3, 4, 5]. Although lanthanide doping can introduce narrow f–f emission [6] or enable multi‐channel luminescence (e.g., in high‐entropy double perovskites) [7], and self‐trapped‐exciton (STE)‐type luminescence has been established in selected double perovskites under strong electron–phonon coupling [4]. Yet, efficient dopant‐free visible emission from the host lattice remains uncommon. Vacancy‐ordered double perovskites (VODPs) with composition A ₂ B□X ₆ (□ denotes an ordered B‐site vacancy), by contrast, comprise a face‐centered lattice of isolated [BX₆] units [8, 9, 10], which promotes carrier localization and strong electron–phonon coupling—key ingredients for self‐trapping and dopant‐free broadband emission [11, 12]. Nevertheless, most reported A ₂ B ^IV X ₆ hosts feature closed‐shell ns⁰/d⁰ B‐site ions (e.g., Sn⁴⁺, Zr⁴⁺, Ti⁴⁺) and consequently lack strong intrinsic visible emission, as illustrated by Cs₂SnI₆ (Figure 1a) [8, 9]. As a result, intrinsic B‐site emission from undoped VODPs remains rare, motivating the exploration of lattice–composition strategies that combine intrinsically emissive centers with structural motifs that mitigate nonradiative losses [8, 9, 10].

FIGURE 1.

Structural comparison of vacancy‐ordered double perovskites. (a) Conventional ns⁰/d⁰ VODPs (A ₂ B ^IV X ₆) exemplified by Cs₂SnI₆, showing weak host emission that requires doping for efficient luminescence. (b) ns² lone‐pair‐active VODPs developed in this work: (Rdabco)₂ B ^III₆ with B = Ge, Sn, showing (Mdabco)₂GeI₆ as representative structure. Isolated [BI₆]^{4 −} octahedra enable bright dopant‐free emission.

ns² metal ions offer a promising route to intrinsic host‐lattice emission because stereochemically active lone pairs can induce local lattice distortion and strengthen electron–phonon coupling, thereby promoting exciton self‐trapping [13, 14, 15]. Nevertheless, widely used ns² centers—including Pb²⁺, Sb³⁺, and Bi³⁺—each face distinct practical constraints—toxicity, limited spectral tunability and emission efficiency across many host lattices [4, 16, 17, 18, 19]. Sn²⁺ and Ge²⁺ are attractive alternatives owing to their ns² electronic configurations and lower toxicity; however, their susceptibility to oxidation under O₂/H₂O can trigger local reorganization and increased nonradiative loss, ultimately compromising emission efficiency and long‐term stability [20, 21, 22, 23]. Accordingly, achieving efficient and durable ns²‐based emission requires not only selecting the appropriate B‐site ion but also structurally isolating and constraining [BX₆] units to help limit oxidant ingress and suppress deleterious local rearrangements [24, 25, 26, 27]. Prior work on Sn²⁺‐based VODPs in bromide frameworks provides useful precedent [28]. Iodide frameworks are particularly attractive for visible emission owing to their narrower band gaps compared to bromide or chloride analogues [29, 30]. For comparison, bulk Cs₄SnI₆, the closest fully inorganic analogue, shows no room‐temperature photoluminescence, and only Cs₄SnI₆ nanocrystals achieve a PLQY of 27% through surface‐induced symmetry breaking [11, 12]. Fully inorganic zero‐dimensional Ge(II) iodide phases remain unreported.

Here we introduce divalent N‐alkylated 1,4‐diazabicyclo[2.2.2]octane (dabco)‐derived diammonium cations, Rdabco²⁺ (R = methyl, ethyl, or propyl; spanning mono‑ and N,N′‑disubstituted variants), [28, 31, 32] to access iodide VODPs (Rdabco)₂ BI₆ with B = Ge(II) or Sn(II). These A ₂ B ^II X ₆‑type lattices adopt a zero‑dimensional (0D) architecture featuring discrete, electronically isolated [BI₆]^{4 −} octahedra separated by bulky organic cations and ordered B‑site vacancies, a structural motif that is widely associated with strong carrier localization and STE emission in low‑dimensional metal halides [33, 34, 35]. Figure 1b highlights (Mdabco)₂GeI₆ (Mdabco = N‐methylated dabco) as a representative member. Systematic cation tuning improves air/moisture tolerance and suppresses local rearrangement in this iodide VODP series. The materials show 575–689 nm STE emission with PLQYs up to 35.0% for the Ge series and 36.8% for the Sn series; the Ge(II) PLQY is about seven times prior Ge(II)–iodide emitters (∼5%) [15, 25, 36, 37, 38, 39, 40, 41]. Structure–photophysics trends are metal‐dependent: Ge PLQYs increase as the [GeI₆]^{4 −} octahedra become more symmetric, whereas Sn PLQYs are governed by radiative–nonradiative kinetic branching rather than a single distortion metric. Together, these results provide a molecular design route to efficient, durable ns²‑metal‑centered emitters and clarify structure–photophysics–exciton‐dynamics relationships in low‐dimensional metal halides [42, 43, 44, 45, 46].

2. Results and Discussion

Ge(H₂PO₂)I and Rdabco diiodide salts were prepared as described in the Supporting Information (Section S1), with the diiodide salts synthesized following established literature procedures [28, 31, 32]. Representative Ge(H₂PO₂)I was confirmed by single‐crystal and powder X‐ray diffraction (Figure S1, Table S1). Single crystals of vacancy‐ordered double perovskites, (Rdabco)₂ BI₆ with B = Ge(II) or Sn(II), were obtained from these reagents using solution‐based routes, including slow cooling from HI/H₃PO₂ at 100°C (∼1 day), solvent evaporation (several days), and vapor diffusion (several days). Figure S2 shows representative millimeter‐scale crystals (0.5–3 mm) grown by the slow‐cooling method. Despite providing crystals suitable for structural analysis, these solution approaches are time‐consuming, solvent‐intensive, and give modest isolated yields (30%–50%).

To enable rapid synthesis with reduced solvent usage, we developed a solvent‐assisted solid‐state grinding (SAG) protocol that delivers phase‐pure products within 15 min at room temperature via mechanochemical activation. Using (DMdabco)₂GeI₆ as an example (DMdabco = N,N′‐dimethylated dabco), Ge(H₂PO₂)I (pale yellow) and DMdabcoI₂ (white) were ground thoroughly in an agate mortar. With continued grinding, faint orange emission emerged under 365 nm irradiation, indicating formation of the emissive phase. A catalytic amount of an HI/H₃PO₂ mixture (∼0.3 mL) was then added; H₃PO₂ helps preserve the Ge²⁺/Sn²⁺ oxidation state under reducing conditions, while HI provides additional iodide to drive the reaction to completion. Further grinding yielded a homogeneous paste exhibiting bright orange emission under 365 nm light, consistent with near‐quantitative conversion (Figure S3). The resulting microcrystalline product was isolated by drying at 80°C and washing with hot ethanol to remove residual reagents, and was stored in sealed containers. This SAG protocol is general across all twelve compositions, affording phase‐pure materials in 75%–95% yield and enabling rapid screening and scalable preparation.

Having established an efficient synthetic route for all twelve compounds, we performed single‐crystal X‐ray diffraction (SCXRD) analysis to elucidate how systematic variations in the organic cation modulate the resulting crystal architectures (crystallographic data in Tables S2–S5). Powder X‐ray diffraction (PXRD) confirms that SAG‐prepared materials are phase‐pure and crystallographically identical to solution‐grown single crystals (Figures S4 and S5).

As shown in Figure 2a, six N‐alkylated dabco derivatives serve as A‐site cations: Mdabco (monomethyl), Edabco (monoethyl), DMdabco (dimethyl), Pdabco (monopropyl), MEdabco (methyl‐ethyl), and DEdabco (diethyl). This series enables systematic tuning of steric hindrance and molecular symmetry. All twelve compounds crystallize in the K₂PtCl₆‐type vacancy‐ordered double perovskite structure with an anti‐fluorite arrangement [47], in which [BI₆]^{4 −} octahedra alternate with ordered vacancies in a three‐dimensional checkerboard pattern (Figures 2b, S6 and S7). The B‐site cations (Ge²⁺ or Sn²⁺) are octahedrally coordinated by iodide to form discrete [BI₆]^{4 −} units that are spatially separated by the ordered vacancies, yielding a zero‐dimensional (0D) inorganic sublattice with electronically isolated octahedra. The organic diammonium cations occupy the cavities between neighboring inorganic units, providing charge balance while simultaneously templating the framework packing.

FIGURE 2.

Molecular design and crystallographic features of (Rdabco)₂BI₆ vacancy‐ordered double perovskites. (a) Molecular structures of six N‐alkylated dabco derivatives (abbreviated as MD = Mdabco, ED = Edabco, DMD = DMdabco, PD = Pdabco, MED = MEdabco, DED = DEdabco). (b) Representative crystal structure showing isolated [BI₆]^{4 −} octahedra alternating with ordered vacancies, with N‐alkylated dabco cations occupying A‐site positions. (c) Unit cell volume (normalized to Z = 4) and density as a function of organic cation for Ge‐ and Sn‐based compounds.

Crystallographic analysis shows that Ge‐based compounds exhibit shorter B–I bond lengths (3.04–3.08 Å) than the Sn‐based analogues (3.17–3.25 Å), consistent with the smaller ionic radius of Ge²⁺ relative to Sn²⁺ (Tables S6 and S7). [39] As shown in Figure 2c, the unit‐cell volume (normalized to Z = 4) increases monotonically from the Mdabco to DEdabco derivatives for both metal centers, spanning 2814–3398 ų (density: 2.57–2.30 g cm^{− 3}) for the Ge series and 2876–3440 ų (density: 2.63–2.36 g cm^{− 3}) for the Sn series, corresponding to volume expansions of 20.8% and 19.6%, respectively. This monotonic increase indicates that lattice expansion is governed primarily by the increasing steric demand of the organic cations.

Beyond lattice expansion, cation substitution also exerts a pronounced influence on crystallographic symmetry. The space group tracks the molecular symmetry of the organic cations: Mdabco‐ and DMdabco‐based compounds crystallize in the cubic space group $Pa\overline{3}$, whereas Edabco‐, Pdabco‐, and MEdabco‐based Ge/Sn derivatives, together with the Sn‐based DEdabco analogue, all crystallize in the orthorhombic space group Pbca. In contrast, (DEdabco)₂GeI₆ crystallizes in the monoclinic space group P2₁/n, consistent with reduced cation symmetry coupled with tighter packing constraints imposed by the smaller Ge²⁺ ion. Overall, the systematic symmetry reduction from cubic to orthorhombic to monoclinic mirrors decreasing molecular symmetry and increasing conformational complexity of the organic cations, highlighting the strong coupling between cation design and lattice symmetry.

Crystal packing density governs organic–inorganic interactions, and exciton confinement in hybrid perovskites, thereby influencing stability and optoelectronic properties [48, 49, 50]. We calculated void fractions using Mercury software [51] to quantify cation‐induced packing variations (Section S1, Supporting Information). Void‐fraction analysis reveals systematic trends across the series (Figure S8). Compounds with smaller cations (Mdabco, Edabco) exhibit accessible void fractions of 3.6%–8.6%, with Sn‐based compounds displaying systematically lower void fractions than Ge analogues, consistent with more efficient packing afforded by the larger Sn²⁺ ionic radius. In contrast, larger‐cation compounds (DMdabco, Pdabco, MEdabco, and DEdabco) show negligible accessible void space, with the organic cations occupying essentially all available volume. This evolution from open to densely packed structures demonstrates continuous tuning of packing density through cation size, establishing a structural design principle for property control in this materials family.

To elucidate how cation structure modulates [BI₆]^{4 −} octahedral geometry, we analyzed coordination environments and distortion metrics for representative structures (Figures 3, S9 and S10; Tables S6 and S7). Progressing from monoalkylated (Mdabco²⁺) to symmetric dialkylated (DMdabco²⁺) and bulkier dialkylated cations (DEdabco²⁺) produces clear changes in local contacts (hydrogen bonding versus steric packing) that are accompanied by systematic variations in octahedral distortion.

FIGURE 3.

Crystal structures and [GeI₆]^{4 −} octahedral coordination environments showing non‐monotonic distortion trends (two cations and hydrogen atoms omitted for clarity). (a) (Mdabco)₂GeI₆: asymmetric N‐H···I hydrogen bonding induces angular distortion (σ ² = 7.18°²). (b) (DMdabco)₂GeI₆: optimal size matching achieves minimal distortion (σ ² = 0.66°²). (c) (DEdabco)₂GeI₆: steric overload causes symmetry‐breaking and maximum distortion (σ ² = 5.40°², Δ_d = 0.00439).

As shown in Figure 3a, (Mdabco)₂GeI₆ crystallizes in the cubic $Pa\overline{3}$ space group. Weak N‐H···I hydrogen bonding (N···I = 3.825 Å) links the protonated nitrogen to the [GeI₆]^{4 −} octahedra, whereas the methylated site engages primarily through van der Waals contacts. This asymmetric interaction mode is associated with pronounced angular distortion: while all Ge–I bond lengths remain equivalent (3.04 Å), the twelve I–Ge–I angles split into two sets of six (87.44° and 92.56°), deviating substantially from the ideal 90° and yielding a large bond‐angle variance (σ ² = 7.18°²). In contrast, (DMdabco)₂GeI₆ (Figure 3b) also adopts $Pa\overline{3}$ but exhibits markedly reduced distortion. The high‐symmetry DMdabco²⁺ cation lacks N–H groups, replacing directional N–H···I bonding with weaker, more isotropic C–H···I contacts and van der Waals interactions. Despite the absence of strong directional bonding, the cation size matches the [GeI₆]^{4 −} cavity, efficiently filling the available volume and enforcing a well‐defined orientation. Accordingly, the octahedra are significantly more regular: the I–Ge–I angles cluster near 90° (90.77° and 89.23°), giving the lowest bond‐angle variance in the Ge series (σ ² = 0.66°²). For the larger (DEdabco)₂GeI₆ (Figure 3c), the longer ethyl substituents drive a symmetry reduction to monoclinic P2₁/n. Relative to Figure 3a,b (dashed arrows), the cation–framework contacts reorient substantially, and the [GeI₆]^{4 −} octahedra undergo strong steric‐induced distortion. Ge–I distances split into three pairs (3.100, 3.080, and 3.060 Å), and the I–Ge–I angles span 86.91–93.09°, further departing from 90°. This member exhibits the highest distortion index (Δ_d = 0.00439) and a large bond‐angle variance (σ ² = 5.40°²) among the Ge‐based compounds, providing a structural basis for investigating how ground‐state octahedral irregularity influences excited‐state dynamics and emission efficiency.

The remaining Ge‐based compounds (Edabco²⁺, Pdabco²⁺, and MEdabco²⁺; Figure S9) and all Sn‐based analogues (Figure S10) follow the same nonmonotonic trend (Tables S6 and S7): octahedral regularity improves from Mdabco²⁺ to DMdabco²⁺ as cation size and symmetry better match the cavity, but decreases for bulkier cations as steric constraints dominate. Beyond packing effects, ns² B‐site ions such as Ge²⁺ and Sn²⁺ are intrinsically susceptible to symmetry‐lowering distortions via the second‐order Jahn–Teller effect (SOJT), which enables s–p mixing and electronic stabilization [52, 53]. These structural variations are expected to modulate band‐edge character and electron–phonon coupling, influencing the photophysical properties examined below [54, 55].

To investigate how structural variation influences luminescence, we characterized the photophysical properties of all 12 compounds. For the Ge‐based series (Figure 4), optical bandgaps estimated from UV–vis diffuse‐reflectance spectra using the Kubelka–Munk transformation span 2.87–3.45 eV (Figure 4a). DFT‐calculated band structures are provided in Figure S11. Under 365 nm excitation, all Ge‐based compounds exhibit bright photoluminescence spanning orange to deep red (Figure 4b). The emission maximum progressively red‐shifts from 625 nm for (Mdabco)₂GeI₆ to 689 nm for (Edabco)₂GeI₆ (Figure 4c), and the longest‐wavelength emitters approach the visible/NIR‐I boundary (∼700 nm), a spectral region relevant to in vivo fluorescence imaging [56, 57]. The close correspondence between PLE spectra and absorption onsets supports host‐lattice excitation of the isolated [GeI₆]^{4 −} units rather than extrinsic impurity centers. Consistent with STE emission, these compounds display characteristic features summarized in Table S8: large Stokes shifts (1.41–1.51 eV), broad emission bands (FWHM = 121–153 nm), and sub‐microsecond lifetimes (0.14–0.82 µs; Figure 4d), indicative of strong exciton–phonon coupling and substantial excited‐state lattice relaxation [58].

FIGURE 4.

Photophysical properties of Ge‐based (Rdabco)₂GeI₆ vacancy‐ordered iodide perovskites. (a) Optical band gaps (2.87–3.45 eV) obtained from UV–vis diffuse reflectance spectra using the Kubelka–Munk transformation. (b) Photographs of representative single crystals under 365 nm excitation and the corresponding room‐temperature PLQYs. (c) Normalized absorption (Abs., dashed), excitation (Ex.), and emission (Em.) spectra. (d) Time‐resolved photoluminescence decay traces (symbols) with biexponential fits (solid lines); average lifetimes (τ _ave) are indicated.

Room‐temperature PLQY measurements show substantial variation across the Ge‐based series (Figure 4b, Table S8). (DMdabco)₂GeI₆ exhibits the highest PLQY (35.0%), followed by (Pdabco)₂GeI₆ (21.4%) and (Mdabco)₂GeI₆ (14.2%). In contrast, (Edabco)₂GeI₆, (MEdabco)₂GeI₆, and (DEdabco)₂GeI₆ display much lower PLQYs of 4.6%, 3.2%, and 1.7%, respectively, corresponding to a >20‐fold spread across the series. Although all members share the same 0D structural motif of isolated [GeI₆]^{4 −} octahedra, the PLQY trend correlates with the cation‐dependent evolution of octahedral distortion and crystallographic symmetry. For example, the most efficient emitter, (DMdabco)₂GeI₆ (cubic $Pa\overline{3}$, σ ² = 0.66°², Δ_d = 0), features the smallest bond‐angle variance and negligible bond‐length distortion, whereas (DEdabco)₂GeI₆ (monoclinic P2₁/n, σ ² = 5.40 °², Δ_d = 0.00439) shows markedly reduced PLQY (1.7%) alongside pronounced distortion (Tables S6 and S8). These results suggest that ground‐state octahedral regularity influences emission efficiency, with the mechanistic link between structural distortion and nonradiative losses further elucidated through excited‐state geometry analysis (vide infra). By contrast, metal–metal separations do not account for the PLQY spread: Ge–Ge distances (9.68–10.98 Å) and Sn–Sn distances (10.06–10.86 Å) vary within a relatively narrow window and do not correlate with the >20‐fold PLQY variation (Table S6).

The emission lifetimes vary systematically across the Ge‐based series (0.14–0.82 µs; Figure 4d) and are generally longer for higher‐PLQY compositions. For (Edabco)₂GeI₆ and (MEdabco)₂GeI₆, SCXRD indicates pronounced cation disorder/dynamics (Figure S12), evidenced by unusually large anisotropic displacement parameters (ADPs) and positional disorder of the alkyl substituents (coexisting flipped conformations). Such dynamic disorder can amplify local lattice fluctuations and strengthen electron–phonon coupling, thereby facilitating additional nonradiative decay pathways, consistent with the shortened lifetimes and reduced quantum efficiencies [59, 60].

To elucidate the microscopic origin of the PLQY variations, we extracted radiative (k_r) and nonradiative (k_nr) decay rates from the measured PLQYs and lifetimes (Table S9). Across the Ge‐based series, PLQY differences are governed primarily by changes in k_nr rather than k_r. The high‐efficiency emitter, (DMdabco)₂GeI₆ (PLQY = 35.0%), combines a relatively fast radiative rate (k_r ≈ 0.43 × 10⁶ s^{− 1}) with the lowest nonradiative rate in the series (k_nr ≈ 0.79 × 10⁶ s^{− 1}). In contrast, low‐efficiency members such as (MEdabco)₂GeI₆ (3.2%) and (Edabco)₂GeI₆ (4.6%) exhibit smaller k_r values (∼0.23 × 10⁶ s^{− 1}) together with substantially elevated k_nr values (>4.7 × 10⁶ s^{− 1}). Overall, k_nr spans 0.79–6.91 × 10⁶ s^{− 1}, whereas k_r varies over a narrower range of 0.23–0.43 × 10⁶ s^{− 1}, indicating that suppression of nonradiative decay is the dominant determinant of quantum efficiency in this family.

To establish the mechanistic link between ground‐state structural regularity and nonradiative decay rates, we performed DFT optimization of excited‐state geometries based on the experimental ground‐state structures from SCXRD for the highest‐PLQY (DMdabco)₂GeI₆ and lowest‐PLQY (DEdabco)₂GeI₆ representatives (Figure S13). Upon photoexcitation, both compounds undergo Jahn–Teller distortion of the [GeI₆]^{4 −}octahedra, but the extent of structural reorganization differs markedly. (DEdabco)₂GeI₆ exhibits pronounced distortion with axial Ge–I bonds elongating by ∼1.0 Å (from ∼3.08 Å to ∼4.09 Å) and asymmetric equatorial compression, whereas (DMdabco)₂GeI₆ undergoes substantially reduced reorganization with axial elongation of only ∼0.6 Å (from ∼3.06 Å to ∼3.66 Å) and uniform equatorial compression. The substantially larger excited‐state reorganization in (DEdabco)₂GeI₆ compared to (DMdabco)₂GeI₆ is consistent with its markedly reduced PLQY and inefficient emission, establishing that compounds with symmetric, well‐matched ground‐state octahedra experience smaller excited‐state lattice reorganization and consequently exhibit lower nonradiative recombination losses. The correlation between ground‐state octahedral regularity (σ ²) and excited‐state reorganization energy thus provides a unified structure–dynamics–efficiency relationship across the Ge‐based series. Collectively, organic‐cation selection enables concurrent tuning of radiative and nonradiative pathways, with the highest‐PLQY compositions achieving a favorable combination of relatively fast k_r and strongly suppressed k_nr through minimized excited‐state structural reorganization. Notably, even polycrystalline powders prepared by the rapid SAG method retain high PLQYs of 28.9% for (DMdabco)₂GeI₆ and 20.6% for (Pdabco)₂GeI₆ (Table S8), indicating that the high quantum efficiency is intrinsic to these compositions and underscoring the generality of the SAG route.

Compared with the Ge analogues, the Sn‐based (Rdabco)₂SnI₆ series shows a narrower bandgap tunability (2.83–2.99 eV; Figure 5a), consistent with stronger Sn–I orbital hybridization and the stereochemically active 5s² lone pair of Sn²⁺, which can favor SOJT‐type symmetry‐lowering distortions [14, 53, 61]. DFT calculations reproduce the experimentally observed bandgap trend (Figure S14). Under 365 nm excitation, all Sn compounds display intense broadband emission from 575 to 650 nm (Figure 5c), and the PLE spectra closely follow the absorption edges, supporting intrinsic emission from isolated [SnI₆]^{4 −} octahedra. The Sn series retains the hallmarks of STE emission (Table S10), including large Stokes shifts (1.02–1.40 eV), broad bandwidths (FWHM = 121–136 nm), and sub‐microsecond lifetimes (0.11–0.37 µs; Figure 5d). Relative to the Ge series, the shorter lifetimes and smaller Stokes shifts of the Sn compounds indicate faster excited‐state decay dynamics, consistent with stronger spin–orbit coupling (SOC) of the heavier ns² center and enhanced singlet–triplet mixing in Sn‐based lattices [15].

FIGURE 5.

Photophysical properties of Sn‐based (Rdabco)₂SnI₆ vacancy‐ordered iodide perovskites. (a) Optical band gaps (2.83–2.99 eV) obtained from UV–vis diffuse reflectance spectra using the Kubelka–Munk transformation. (b) Photographs of representative single crystals under 365 nm excitation and the corresponding room‐temperature PLQYs. (c) Normalized absorption (Abs., dashed), excitation (Ex.), and emission (Em.) spectra. (d) Time‐resolved photoluminescence decay traces (symbols) with biexponential fits (solid lines); τ _ave values are indicated.

The PLQY distribution of the Sn series differs from that of the Ge analogues (Figure 5b, Table S10). (Pdabco)₂SnI₆ reaches the highest PLQY of 36.8%, while (DEdabco)₂SnI₆, (DMdabco)₂SnI₆, (Mdabco)₂SnI₆, (Edabco)₂SnI₆, and (MEdabco)₂SnI₆ exhibit PLQYs of 18.1%, 16.1%, 16.0%, 14.5%, and 9.4%, respectively. Overall, the PLQY spread in the Sn series (∼4‐fold) is markedly smaller than that observed for the Ge series (∼20‐fold), with multiple compositions clustering in a narrow range (14%–18%), indicating a reduced sensitivity of Sn emitters to cation‐driven structural perturbations.

Consistent with the compressed PLQY distribution, the Sn lifetimes span a narrower window (0.11–0.37 µs; Figure 5d). To probe the determinants of efficiency, we extracted k_r and k_nr for the Sn series (Table S11). Unlike the Ge series, where the PLQY spread is dominated by k_nr, the Sn series exhibits comparable variations in k_r and k_nr, and similar PLQYs can arise from distinct rate combinations. For example, (Pdabco)₂SnI₆ achieves the highest PLQY through a favorable combination of the lowest nonradiative loss in the series, whereas other members reach comparable efficiencies through different balances between radiative and nonradiative channels. These results suggest that Sn emission efficiencies are governed by competition among multiple relaxation pathways rather than by a single structural descriptor. (Edabco)₂SnI₆ and (MEdabco)₂SnI₆ show positional disorder of the ethyl substituents (Figure S15), which may contribute to their comparatively lower PLQYs. Notably, SAG‐prepared polycrystalline powders retain high efficiencies (e.g., PLQY = 30.7% for (Pdabco)₂SnI₆; Table S10), and SAG‐ and solution‐grown samples exhibit nearly identical PL and PLE peak positions for both the Ge and Sn series (Figures S16 and S17), indicating that the synthetic route does not alter the intrinsic photophysical characteristics.

Remarkably, the Ge(II) iodide VODP (DMdabco)₂GeI₆ achieves a room‐temperature PLQY of 35.0%, setting a new benchmark for Ge(II)‐based metal‐halide emitters [15, 25, 36, 37, 38, 39, 40, 41, 62]. This value exceeds the prior Ge(II)–iodide record, [(PEA)₆GeI₈(Bz)₂] (PLQY = 5.0%) [25], and also surpasses the best reported Ge(II)–bromide emitter, (C₅N₂H₁₄)GeBr₄ (PLQY = 7.15%) [62]. By comparison, (Pdabco)₂SnI₆ reaches a PLQY of 36.8%, placing it among the highest‐efficiency low‐dimensional Sn(II) iodide emitters [14, 15, 23, 24, 25, 26, 27, 28]. These high efficiencies are consistent with the present structural design: N‐alkylated dabco cations template a vacancy‐ordered 0D lattice of electronically isolated [BI₆]^{4 −} octahedra, providing an effective host for ns²‐centered luminescence.

To probe the electronic origin of the optical transitions, we analyzed the projected density of states (PDOS) (Figures S18 and S19). The valence‐band maximum is dominated by I 5p states with antibonding contributions from metal ns²‐derived lone‐pair states, whereas the conduction‐band minimum mainly originates from metal np states. In contrast, the organic‐cation states lie far from the band edges and contribute negligibly to the frontier electronic structure. These results indicate that optical absorption and emission are governed by transitions within the isolated [BI₆]^{4 −} units, while the organic cations modulate luminescence indirectly through structural effects such as packing and local distortion.

To elucidate the STE emission mechanism and compare exciton dynamics in Ge‐ and Sn‐centered lattices, we performed variable‐temperature PL and TRPL measurements on the highest‐PLQY representatives, (DMdabco)₂GeI₆ and (Pdabco)₂SnI₆, over 83–325 K (Figures 6 and 7). The packing views (Figures 6a and 7a) highlight the vacancy‐ordered framework that spatially isolates the [BI₆]^{4 −} octahedra, establishing a 0D structural motif conducive to exciton localization and self‐trapping. Below, we first describe the key temperature‐dependent spectral and kinetic signatures (Figures 6b,c and 7b,c) and then quantify thermal quenching and exciton–phonon coupling using fits to the integrated intensity and linewidth data (Figures 6d–f and 7d–f).

FIGURE 6.

Temperature‐dependent STE emission analysis of (DMdabco)₂GeI₆. (a) Crystal packing highlighting the vacancy‐ordered lattice and isolated [GeI₆]^{4 −} units. (b) PL spectra (83–325 K). (c) TRPL decay traces with biexponential fits; τ _ave values are indicated. (d) Arrhenius analysis of thermal quenching, yielding E_a and E_b (shown). (e) FWHM versus temperature fitted to extract S and ħω_phonon (shown). (f) FWHM versus 1/T analysis to evaluate Γ _op (shown).

FIGURE 7.

Temperature‐dependent STE emission analysis of (Pdabco)₂SnI₆. (a) Crystal packing highlighting the vacancy‐ordered lattice and isolated [SnI₆]^{4 −} units. (b) PL spectra (83–325 K). (c) TRPL decay traces with biexponential fits; τ _ave values are indicated. (d) Arrhenius analysis of thermal quenching, yielding E_a and E_b (shown). (e) FWHM versus temperature fitted to extract S and ħω_phonon (shown). (f) FWHM versus 1/T analysis to evaluate Γ _op (shown).

Consistent with STE emission, both compounds exhibit pronounced thermal quenching in their variable‐temperature PL spectra (Figures 6b and 7b). Upon heating from 83 to 325 K, the integrated PL intensity decreases substantially, with I(83 K)/I(325 K) ≈ 44 for (DMdabco)₂GeI₆ and ≈ 12 for (Pdabco)₂SnI₆, indicative of thermally activated nonradiative relaxation pathways [26]. The Ge compound shows an essentially invariant emission maximum (622 nm at 83 K vs 625 nm at 325 K; Δλ = 3 nm), whereas the Sn compound displays a modest shift to shorter wavelength at low temperature (587 nm at 325 K to 578 nm at 83 K), suggesting a more temperature‐sensitive emissive profile. In both cases, the emission linewidth increases markedly upon heating (Ge: 73 → 141 nm; Sn: 76 → 128 nm), consistent with increased phonon population and enhanced exciton–phonon interactions at elevated temperature [23].

Time‐resolved PL (TRPL) measurements further clarify the temperature‐dependent dynamics (Figures 6c and 7c). Upon heating from 83 to 325 K, τ _ave decreases markedly—from 3.34 to 0.27 µs for (DMdabco)₂GeI₆ and from 1.13 to 0.16 µs for (Pdabco)₂SnI₆—paralleling the steady‐state intensity quenching and confirming that nonradiative losses dominate at elevated temperature. Temperature‐dependent average lifetimes plotted over the full 83–325 K range further confirm the onset of thermally activated de‐trapping above ∼225 K, consistent with enhanced electron–phonon coupling at elevated temperatures (Figure S20) [26].

To quantify the energetics underlying thermal quenching, we fitted the temperature‐dependent integrated PL intensity using an Arrhenius model (Figures 6d and 7d) [63, 64]. For (DMdabco)₂GeI₆, the fit yields E_a = 470.90 meV and E_b = 75.53 meV (R² = 0.9989), whereas for (Pdabco)₂SnI₆, E_a = 341.99 meV and E_b = 31.56 meV (R² = 0.9977). The large E_a values (>300 meV) for both compounds are consistent with deep self‐trapping, and the larger E_a for the Ge compound indicates a deeper trapping potential and a higher energetic barrier against thermal deactivation.

Temperature‐dependent full width at half maximum (FWHM) analysis provides further evidence for strong exciton–phonon coupling underlying the broadband STE emission (Figures 6e and 7e). Fitting the FWHM(T) dependence yields large Huang–Rhys factors [65] of S = 52.45 for (DMdabco)₂GeI₆ and S = 43.45 for (Pdabco)₂SnI₆, indicative of very strong coupling and deep self‐trapping. The extracted effective phonon energies are ℏω_phonon = 11.31 meV (Ge) and 15.46 meV (Sn), consistent with low‐energy lattice vibrations associated with the [BI₆]^{4 −} octahedra that participate in the self‐trapping and emission processes [66].

Notably, despite similarly large Huang–Rhys factors, the two systems show markedly different linewidth–temperature behavior in the 1/T representation (Figures 6f and 7f). The extracted optical‐phonon broadening parameter, Γ_{o p} , is 149 meV for (DMdabco)₂GeI₆ but increases to 420 meV for (Pdabco)₂SnI₆ [67, 68]. This difference is consistent with more pronounced phonon‐assisted dephasing/broadening in the Sn compound and points to different excited‐state relaxation behavior in Ge‐ versus Sn‐centered STEs that is not fully captured by S alone.

Taken together, the variable‐temperature analysis and the structure–photophysics dataset across all twelve compounds point to a clear difference between the Ge and Sn series: Ge emitters display a much stronger dependence of PLQY on octahedral distortion, whereas Sn emitters maintain comparatively high efficiencies over a wider distortion range. For example, (DMdabco)₂GeI₆ (cubic $Pa\overline{3}$) features a minimally distorted [GeI₆]^{4 −} unit (Δ_d = 0, σ ² = 0.66°²) with six equivalent Ge–I bonds (3.0554 Å) and delivers the highest PLQY (35.0%) together with the lowest k_{n r} (0.79 × 10⁶ s^{− 1}), in line with its relatively small linewidth‐broadening term (Γ_op = 149 meV). In contrast, (DEdabco)₂GeI₆ exhibits larger distortion metrics (Δ_d = 0.00439; σ ² = 5.41 °²) accompanied by pronounced excited‐state reorganization (∼1.0 Å axial elongation), resulting in a much lower PLQY (1.7%; >20 times lower). In the Sn series, efficient emission is retained even at appreciable distortion, as exemplified by (Pdabco)₂SnI₆ (PLQY = 36.8%; σ ² = 2.31 °²) and (DEdabco)₂SnI₆ (PLQY = 18.1%; σ ² = 11.69 °²), indicating that Sn²⁺ emitters are less sensitive to ground‐state structural irregularity owing to the reduced stereochemical expression of the 5s² lone pair.

The STE formation and emission pathways are illustrated in the configuration‐coordinate diagrams in Figure 8. Both compounds exhibit broadband emission with large Stokes shifts (1.48 eV for (DMdabco)₂GeI₆ and 1.10 eV for (Pdabco)₂SnI₆), consistent with the characteristic of pronounced lattice relaxation during STE formation. Upon photoexcitation, the excited electrons relax from the excited state to the minimum energy state of STEs, followed by radiative recombination to the ground state. The larger Stokes shift and higher activation energy of (DMdabco)₂GeI₆ (E_a = 470.90 meV) relative to (Pdabco)₂SnI₆ (E_a = 341.99 meV) reflect a deeper self‐trapping potential well and stronger lattice reorganization in the Ge‐centered system.

FIGURE 8.

Configuration‐coordinate diagrams illustrating STE formation and emission in (a) (DMdabco)₂GeI₆ and (b) (Pdabco)₂SnI₆. GS and ES denote the ground state and excited state, respectively.

The difference in distortion dependence may reflect the distinct stereochemical activity of the ns² lone pair and its coupling to excited‐state lattice relaxation. Ge²⁺ (4s²) is often considered to exhibit stronger stereochemical activity, such that local symmetry breaking can more readily open additional nonradiative relaxation pathways and increase phonon‐assisted energy loss.[13, 55] In Sn²⁺ (5s²) systems, the stereochemical expression of the lone pair is weaker, and the heavier cation can introduce stronger SOC that promotes singlet–triplet mixing within the STE manifold, thereby partially offsetting distortion‐induced nonradiative loss [11, 12, 15, 24, 25, 26, 27]. Together with the suppression of exciton migration by the electronically isolated [BI₆]^{4 −} motif in 0D structure, the above factors provide a plausible rationale for why Ge‐based emitters require higher structural regularity, whereas Sn analogues maintain relatively high luminescence efficiency across a wider range of distortion [24, 25, 26, 27].

Overall, the variable‐temperature analysis across the full compositional series underscores a clear contrast between the two families: Ge emitters are more sensitive to octahedral irregularity, whereas Sn emitters retain comparatively high efficiencies over a broader distortion window. Variable‐temperature data collected for four additional compositions support this conclusion (Figure S21). Specifically, (DMdabco)₂SnI₆ and (Pdabco)₂GeI₆ display the same qualitative signatures as the benchmark compounds—pronounced thermal quenching upon heating, concomitant lifetime shortening, and progressive linewidth broadening—whereas (Edabco)₂GeI₆ and (Edabco)₂SnI₆ show an unusual blue shift of the emission maximum (∼40 nm) upon cooling. This anomalous shift may reflect temperature‐dependent changes in cation dynamics (e.g., partial freezing of ethyl‐substituent disorder) that perturb the local octahedral environment and, in turn, reshape the STE energy landscape [59, 60].

The structural analysis above establishes that high PLQY correlates with low octahedral distortion and high crystal symmetry. To verify whether these structural features can enhance environmental stability, we monitored phase stability under ambient laboratory conditions using PXRD (Figure 9).

FIGURE 9.

Environmental stability of 3D hybrid perovskites and 0D vacancy‐ordered double perovskites. (a, b) PXRD of (MA)GeI₃ and (MA)SnI₃ with degradation peaks (*). (c, d) PXRD of (DMdabco)₂GeI₆ and (Pdabco)₂SnI₆ stable for 60 days. (e, f) Structures showing 3D corner‐sharing distorted [GeI₆]^{4 −} octahedra (faded octahedra indicate neighboring units) with small MA⁺ and 0D isolated perfect [GeI₆]^{4 −} octahedra protected by large DMdabco²⁺ spatial barrier.

For Ge/Sn‐based 3D hybrid perovskites, rapid degradation under ambient exposure is well documented and is commonly associated with moisture‐assisted decomposition and oxidation of the metal–halide framework [69, 70, 71]. (MA)GeI₃ shows impurity peaks (marked with asterisks) within one day (Figure 9a), while (MA)SnI₃ degrades within 3 days (Figure 9b). In stark contrast, the 0D compounds exhibit exceptional stability. (DMdabco)₂GeI₆ (Figure 9c) and (Pdabco)₂SnI₆ (Figure 9d) both retain phase purity with no detectable impurity reflections after 60 days of continuous ambient exposure, with no obvious loss of diffraction intensity. Additional stability tests were conducted to further assess the structural and optical stability of both compounds. After heating at 105°C for 24 h, the PXRD patterns show no noticeable changes, indicating that the crystal structures remain intact under thermal treatment (Figure S22). Both powders also remain light‐colored under daylight and highly emissive under 365 nm UV irradiation after the stability tests (Figure S23). After exposure to 100% relative humidity for 7 days, no obvious change is observed in the PL emission profile or intensity for either compound. In contrast, after thermal treatment, both compounds show reduced PL intensity and PLQY, whereas the emission profiles remain largely unchanged. These results indicate that thermal treatment affects the emission efficiency without altering the overall crystal structure.

The structural origins of this stability difference are illustrated in Figure 9e,f. (MA)GeI₃ adopts a 3D corner‐sharing metal–iodide network with small MA⁺ cations (Figure 9e), which provides continuous pathways and accessible surfaces for interaction with environmental degradants [70]. In contrast, (DMdabco)₂GeI₆ features 0D isolated [GeI₆]^{4 −} octahedra completely surrounded by large DMdabco²⁺ cations (Figure 9f). The large organic cations are expected to hinder H₂O/O₂ ingress by reducing accessible free volume and interrupting diffusion pathways, while the highly regular octahedral geometry (σ ² = 0.66°²) may further mitigate distortion‐driven bond rearrangements that promote nonradiative and chemical degradation.

Thermal stability measurements further confirm the robustness of these materials, with decomposition temperatures of 245°C for (DMdabco)₂GeI₆ and 200°C for (Pdabco)₂SnI₆ (Figure S24). The combination of 0D structural isolation, protective cation encapsulation, and optimized octahedral geometry establishes exceptional stability while maintaining high PLQYs (35.0% for Ge, 36.8% for Sn), positioning these materials as promising candidates for solid‐state lighting and scintillator applications [72, 73].

3. Conclusion

In conclusion, we report a series of highly emissive and stable iodide vacancy‐ordered double perovskites, (Rdabco)₂BI₆, spanning twelve compositions assembled from six N‐alkylated dabco‐derived cations and stereochemically active ns² Ge(II) and Sn(II) centers. Rapid solvent‐assisted grinding provides a convenient route to phase‐pure powders for comparative measurements across the compositional series. Cation engineering tunes crystal symmetry, octahedral distortion, and emission wavelength, delivering 575–689 nm self‐trapped exciton emission with photoluminescence quantum yields up to 35.0% for the Ge series and 36.8% for the Sn series. Across the series, higher PLQYs coincide with reduced nonradiative decay and lower octahedral distortion, highlighting symmetry control as an effective handle to suppress nonradiative losses. The 0D lattice of isolated octahedra, together with bulky organic cations, imparts robust environmental and thermal stability. This work establishes iodide vacancy‐ordered double perovskites as a tunable platform for efficient, stable, lead‐free emitters enabled by cation control of ns² metal–halide units.

Author Contributions

Le Ye: methodology, conceptualization, investigation, visualization, validation, formal analysis, data curation, software, writing – original draft. Yarou Duan: writing – original draft, investigation, methodology, validation, software, formal analysis, data curation. Jun Luo: methodology, software, validation, visualization, writing – original draft. Yutong Lin: methodology, investigation, validation. Mengqiu Cai: software, validation, supervision. Lingling Mao: conceptualization, investigation, funding acquisition, writing – review and editing, writing – original draft, methodology, validation, visualization, supervision, project administration, resources.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File 1: anie72845‐sup‐0001‐SuppMat.docx.

Supporting File 2: anie72845‐sup‐0002‐Data.zip.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.