Multisite Molecular Coordination for Defect and Structural Regulation in Carbon‐Cathode Hole‐Transport‐Layer‐Free CsPbI₂Br Solar Cells

ABSTRACT

Carbon‐cathode hole‐transport‐layer‐free (HTL‐free) CsPbI₂Br solar cells offer a promising route toward high‐performance‐to‐cost photovoltaics, yet their performance is largely constrained by defect‐related losses. Here, a molecular strategy based on terpyridine derivatives is developed to regulate defect states, lattice distortion, and interfacial charge‐transfer energetics in CsPbI₂Br films. Owing to cooperative N/O coordination, [2,2′:6′,2″‐terpyridine]‐6,6″‐dicarboxylic interacts with undercoordinated metal ions (Pb²⁺/Sn⁴⁺) and vacancy defect sites (I/O) at defect‐rich surfaces, grain boundaries, and buried CsPbI₂Br/SnO₂ interfaces. This coordination‐assisted regulation is associated with improved film crystallinity, reduced electronic non‐uniformity, and more efficient interfacial charge extraction. As a result, carbon‐cathode HTL‐free CsPbI₂Br devices achieve a leading comprehensive performance with a power conversion efficiency of 16.02%, minimal hysteresis, improved operational stability, and exceptional reproducibility. Furthermore, the applicability of this molecular approach is demonstrated in MAPbI₃‐based devices, highlighting its potential as a general strategy for defect regulation in the related optoelectronic devices.

A terpyridine‐based molecular (TPY‐DCA) strategy achieves cooperative N/O coordination, concurrently passivating undercoordinated Pb²⁺/Sn⁴⁺ ions and suppressing vacancy‐related defects. This multisite defect healing, spanning surfaces, grain boundaries, and buried interfaces, enables carbon‐cathode HTL‐free CsPbI₂Br solar cells (ITO/SnO₂/CsPbI₂Br/carbon) to achieve a PCE of 16.02% with significantly enhanced operational stability, providing a generalizable molecular approach for perovskite photovoltaics.

1. Introduction

Perovskite solar cells (PSCs) have rapidly emerged as one of the most promising photovoltaic technologies, with certified power conversion efficiencies (PCEs) now exceeding 27% [1]. Their exceptional properties, including strong optical absorption, long carrier diffusion lengths, defect tolerance, and solution‐processability, make them desirable candidates for next‐generation cost‐effective photovoltaics [2, 3]. However, widespread commercialization is still hampered by inadequate long‐term stability [4]. Organic A‐site cations used in conventional perovskites are vulnerable to thermal, humidity, and electrical stresses, while the commonly used hole‐transport layers (HTLs) are prone to undesirable chemical reactions, phase transitions, and morphological changes during operation [5, 6]. These degradation pathways aggravate non‐radiative recombination and promote ion migration, ultimately leading to substantial efficiency losses [7].

All‐inorganic CsPbI₂Br, characterized by high thermal robustness and an appropriate bandgap, offers a credible route to overcoming these stability limitations [8, 9, 10]. In parallel, carbon‐cathode HTL‐free CsPbI₂Br PSCs have attracted growing attention owing to their inherent material stability, cost‐efficiency, and compatibility with scalable manufacturing processes [11]. Removing the HTL eliminates a major failure‐prone component in conventional PSCs and simplifies device fabrication. Nonetheless, HTL‐free CsPbI₂Br PSCs face native limitations because the direct CsPbI₂Br/carbon interface contact lacks electronic selectivity, resulting in substantial charge recombination [12, 13, 14, 15]. Moreover, the CsPbI₂Br absorber inherently contains abundant defects at interfaces more than in the bulk, like uncoordinated Pb²⁺ sites and mobile I⁻ vacancies (V_I), which persist as active recombination centers that severely limit open‐circuit voltage (V _oc) and fill factor (FF) without the passivation of HTL [16, 17]. Consequently, the PCE of carbon‐cathode HTL‐free CsPbI₂Br devices has so far mainly been confined to ∼15%, well below their intrinsic potential of 20% [18].

Numerous defect‐passivation approaches have been reported, including surface modification, buried interface optimization, and crystallization regulation [19, 20, 21, 22, 23, 24, 25]. However, most of these strategies suffer from incomplete or strongly site‐specific passivation, as they typically target only a single defect family and fail to establish sufficiently robust interactions with multiple defect types simultaneously. Very few molecular systems are capable of concurrently healing uncoordinated Pb²⁺ defects, suppressing V_I, and establishing a strong charge‐extraction bridge at the CsPbI₂Br/SnO₂ electron‐transport layer (ETL) interface [26]. Thereby, achieving simultaneous control of defects on the surface, at grain boundaries, and across buried interfaces remains a critical challenge, particularly in carbon‐cathode HTL‐free CsPbI₂Br PSCs.

In this work, a multisite molecular coordination strategy employing a series of terpyridine (TPY) derivatives [2,2′:6′,2″‐terpyridine]‐4′‐carboxylic acid (TPY‐COOH), [2,2′:6′,2″‐terpyridine]‐6,6″‐dicarboxylic (TPY‐DCA), and [2,2′:6′,2″‐terpyridine]‐4,4′,4″‐tricarboxylic (TPY‐TCA) was introduced, each characterized by one, two, and three carboxyl (COOH) groups, respectively. These molecules combine three key functions within a single scaffold: a tridentate TPY core that chelates uncoordinated metal ions Pb²⁺ and Sn⁴⁺ via N‐donor coordination, COOH groups that bind Pb²⁺ and stabilize V_I through hydrogen bonding, and an extended π‐conjugated backbone that facilitates charge transfer. Among these ligands, TPY‐DCA offers the most favorable balance of multidentate coordination and minimal steric encumbrance, enabling strong multisite interactions at CsPbI₂Br surfaces, grain boundaries, and buried interfaces. This cooperative binding substantially lowers defect density, improves crystallinity, modulates lattice strain, and enhances charge transport. The optimized devices deliver a V _oc of 1.342 V and a PCE of 16.02%, ranking among the highest PCEs reported for carbon‐cathode HTL‐free CsPbI₂Br PSCs, while simultaneously exhibiting remarkable operational stability. This multisite coordination strategy offers a general route to defect regulation in carbon‐cathode HTL‐free PSCs.

2. Results and Discussion

To elucidate the atomic‐level interactions between TPY derivatives and CsPbI₂Br, Density Functional Theory (DFT) calculations were conducted to analyze their essential electron characteristics. The electrostatic potential (ESP) distributions in Figure 1a reveal pronounced negative‐potential regions localized on the pyridyl N atoms and the carbonyl O atoms of the COOH groups in all TPY derivatives, identifying these sites as potential Lewis‐basic coordination centers capable of interacting with undercoordinated Pb²⁺ species at defect‐rich regions. The charge‐density difference (CDD) plots in Figure 1b reveal a redistribution of electron density upon adsorption, characterized by electron accumulation around the N/O donor sites and electron depletion at surface Pb sites. This ligand‐to‐metal charge‐transfer behavior provides strong mechanistic support for the defect‐passivation capability of TPY‐based ligands [27, 28]. Notably, TPY‐DCA exhibits a more extensive charge redistribution, suggesting a relatively stronger and more cooperative multisite coordination interaction with the CsPbI₂Br surface.

FIGURE 1.

(a) DFT‐calculated ESP maps of TPY‐COOH, TPY‐DCA, and TPY‐TCA molecules; (b) Side and top views of optimized CDD distributions of TPY derivatives adsorbed on the CsPbI₂Br surface; (c) Schematic illustration of TPY derivative coordination at the CsPbI₂Br surface, buried CsPbI₂Br/SnO₂ interface, and grain boundaries.

Geometric analysis (Figure S1) indicates that the molecular dimensions of the TPY derivatives favor their localization at defect‐rich regions of CsPbI₂Br films, including surfaces, buried interfaces, and grain boundaries, rather than incorporation into the CsPbI₂Br lattice. Notably, TPY‐DCA exhibits a well‐balanced molecular geometry that enables simultaneous interactions with multiple undercoordinated metal ions, thereby circumventing the inherent coordination‐site and steric constraints of its counterparts. In comparison, TPY‐COOH, with only a single anchoring group, displays limited coordination interactions and thus a relatively weak defect‐passivation effect. Although TPY‐TCA features multiple coordination sites, it suffers from excessive steric hindrance, which limits effective binding to defect sites, particularly at grain boundaries and buried interfaces. Accordingly, the unique structural merits of TPY‐DCA enable the establishment of a more continuous and energetically favorable coordination environment across multiple defect‐rich regions of the CsPbI₂Br films. The schematic illustration in Figure 1c summarizes this multisite coordination, highlighting the concurrent interactions of TPY‐DCA with Pb‐ and I‐related defects in CsPbI₂Br as well as trap‐active sites on SnO₂.

The influence of TPY derivatives on the electronic structure of CsPbI₂Br and the underlying SnO₂ ETL was investigated using surface‐sensitive x‐ray photoelectron spectroscopy (XPS) to probe ligand‐metal interactions. As shown in Figure 2a, the Pb 4f core levels of pristine CsPbI₂Br are located at 142.89 and 137.98 eV, respectively, and exhibit a downward shift to lower binding energies (142.62 and 137.71 eV) upon TPY‐DCA modification. This shift indicates an increased local electron density around Pb²⁺, consistent with electron donation from N/O donor groups through coordination interactions [29, 30]. A concurrent binding energy downshift is observed for the I 3d peaks, as depicted in Figure 2b, indicating a stabilized I⁻ chemical environment and a reduced electrostatic driving force for I⁻ migration [31]. In comparison, TPY‐COOH and TPY‐TCA induce only marginal shifts in both Pb 4f and I 3d spectra (Figure S2), consistent with their comparatively weaker coordination capability. These spectroscopic observations agree with DFT calculations (Figure 2c), which reveal that TPY‐DCA possesses the most favorable adsorption energy (E _ads) of −2.57 eV, indicating a stronger thermodynamic propensity for surface binding and coordination.

FIGURE 2.

XPS spectra of (a) Pb 4f and (b) I 3d on the surface of pristine and TPY‐DCA‐treated CsPbI₂Br films; (c) DFT‐calculated E _ads of TPY derivatives on the CsPbI₂Br surface; The XPS spectra of (d) Sn 3d and (e) O 1s on the SnO₂ side of the buried interface of CsPbI₂Br/SnO₂ before and after TPY‐DCA modification; (f) DFT‐calculated E _ads of TPY derivatives at the CsPbI₂Br/SnO₂ interface; (g) Comparison of E _ads at the surface, grain boundary, and buried interface; (h) FTIR spectra of TPY‐DCA, TPY‐DCA coordinated with PbI₂ and SnO₂, respectively; (i) Raman spectra of pristine PbI₂ and TPY‐DCA‐treated PbI₂.

To further resolve coordination effects at the buried CsPbI₂Br/SnO₂ interface, a UV‐curable photoresist lift‐off strategy was employed to expose interfacial regions (Figure S3) selectively. XPS spectra collected from the CsPbI₂Br side (Figure S4) show that TPY‐DCA induces the largest downshift in both Pb and I core levels, demonstrating enhanced electronic stabilization of defect‐rich regions at the buried interface [32]. On the SnO₂ side, TPY‐DCA drives substantial downward shifts in the Sn 3d peaks from 495.29 and 486.69 eV to 494.83 and 486.23 eV, as shown in Figure 2d and Figure S5a, consistent with electron donation into Sn⁴⁺ defect states. Meanwhile, the oxygen vacancy (V_O) component in the O 1s spectrum is significantly diminished in intensity (Figure 2e and Figure S5b), suggesting effective passivation of V_O defects through interfacial charge compensation [33, 34]. The energetic origin of this interfacial stabilization is further supported by DFT calculations, which indicate that TPY‐DCA preferentially anchors to the SnO₂ surface through N‐coordination with an E _ads of −3.64 eV, which is thermodynamically more favorable than O‐anchoring (−3.31 eV), as summarized in Figure 2f and Figure S6. In contrast, TPY‐COOH and TPY‐TCA display considerably less negative E _ads, confirming their limited ability to regulate buried interfacial defects. Beyond the interface, TPY‐DCA also exhibits the most favorable adsorption at grain boundaries (Figure S7), underscoring its capability to stabilize high‐energy defect sites.

The multiscale energetic trends are summarized in Figure 2g, which compares the relative binding strengths of TPY‐DCA at the CsPbI₂Br surface, buried interface, and grain boundaries, consistently showing more favorable interactions than those of the other TPY derivatives. Vibrational spectroscopic analyses further support these theoretical insights. As illustrated in Figure 2h, Fourier‐transform infrared (FTIR) spectra display systematic redshifts of the C═N stretching vibration, along with both asymmetric and symmetric C═O modes, upon complexation with PbI₂ or SnO₂. These shifts are consistent with coordination‐induced perturbation of the corresponding bonds arising from Lewis acid‐base interactions between Pb²⁺/Sn⁴⁺ centers and N/O donor groups [35]. In addition, the full‐range FTIR spectra in Figure S8 reveal newly emerged low‐frequency features attributable to metal‐ligand vibrational modes, further supporting the formation of coordination interactions. Raman spectroscopy offers complementary structural information. As illustrated in Figure 2i, the Pb─I vibrational modes shift from 103 and 198 cm⁻¹ in pristine CsPbI₂Br to 106 and 214 cm⁻¹ after TPY‐DCA treatment, reflecting ligand‐induced electronic polarization of the local bonding environment. Moreover, the emergence of an additional vibration near 141 cm⁻¹, which can be assigned to Pb─O stretching, constitutes direct proof of metal‐O coordination [36].

The structural evolution of CsPbI₂Br films upon incorporation of TPY‐DCA is reflected in noticeable changes in crystallization behavior and film morphology. Top‐view scanning electron microscopy (SEM) images of the pristine film (Figure 3a) show compact but relatively small grains with an average size of 205 nm, indicative of a high nucleation density during rapid crystallization. In contrast, TPY‐DCA incorporation (Figure 3b) enlarges the average grain size to 412 nm, suggesting suppressed excessive nucleation and enhanced lateral grain growth. By comparison, CsPbI₂Br films modified with TPY‐COOH and TPY‐TCA exhibit only moderate grain‐size enlargement (Figure S9), consistent with their weaker coordination capabilities. Cross‐sectional SEM images (Figure S10) further reveal that TPY‐DCA promotes the formation of vertically oriented columnar grains extending across the full absorber thickness, which is beneficial for continuous out‐of‐plane charge transport. The improved film uniformity is further supported by atomic force microscopy (AFM). As shown in Figure 3c,d, and Figure S11, TPY‐DCA substantially reduces the surface roughness of CsPbI₂Br films, lowering the root‐mean‐square (RMS) roughness from 44.9 nm in the pristine film to 23.3 nm. Such a smoother and more compact surface morphology is essential for forming intimate and homogeneous contact with the porous carbon electrode, thereby supporting efficient charge extraction [37].

FIGURE 3.

Top‐view SEM images and corresponding grain‐size statistics of (a) pristine and (b) TPY‐DCA‐modified CsPbI₂Br films; AFM topographies of (c) pristine and (d) TPY‐DCA‐modified CsPbI₂Br films; GIWAXS patterns of (e) pristine and (f) TPY‐DCA‐modified CsPbI₂Br films; GIXRD patterns of the (100) reflection collected at different incident angles (ω) for (g) pristine and (h) TPY‐DCA‐modified films; (i) d‐spacing extracted from the (100) peak as a function of ω for pristine and TPY‐DCA modified films; KPFM surface‐potential maps of (j) pristine and (k) TPY‐DCA films; (l) Corresponding surface potential line profiles.

Two‐dimensional grazing‐incidence wide‐angle x‐ray scattering (GIWAXS) was employed to examine the crystallographic characteristics of CsPbI₂Br films with and without TPY‐DCA modification, as shown in Figure 3e,f. The identical diffraction‐ring positions observed for both pristine and TPY‐DCA‐treated CsPbI₂Br films indicate preservation of the cubic CsPbI₂Br phase. Notably, the TPY‐DCA‐modified film exhibits sharper and more continuous diffraction features, suggesting enhanced long‐range lattice ordering and reduced structural heterogeneity [38]. To further probe lattice distortion along the film depth, grazing‐incidence x‐ray diffraction (GIXRD) measurements were performed at varying incident angles (ω). For the pristine film, the (100) diffraction peak systematically shifts toward lower 2θ values with increasing ω (Figure 3g), revealing a distinct depth‐dependent tensile‐strain gradient [39, 40]. Such behavior indicates nonuniform lattice distortion across the film thickness. In contrast, the TPY‐DCA‐treated CsPbI₂Br shows nearly invariant peak positions over the full ω range (Figure 3h), implying a substantially more uniform strain distribution throughout the absorber layer. Only partial strain relaxation is observed for films treated with TPY‐COOH and TPY‐TCA (Figure S12a,b). Consistently, the corresponding d‐spacing profiles in Figure 3i and Figure S12c present a nearly constant evolution for TPY‐DCA, confirming its effectiveness in mitigating lattice strain variations. These crystallographic observations are consistent with chemically regulated interactions that promote more homogeneous lattice environments in CsPbI₂Br films [41, 42].

Kelvin probe force microscopy (KPFM) was employed to probe the surface electronic landscape of CsPbI₂Br films and the modulation by TPY‐DCA treatment. As shown in Figure 3j,k, the TPY‐DCA‐modified CsPbI₂Br film exhibits a higher average surface potential together with a more uniform potential distribution compared to the pristine counterpart, reflecting a more homogeneous surface electronic environment, which is consistent with suppressed surface electronic disorder upon TPY‐DCA modification [43, 44]. The corresponding surface‐potential line‐scan analysis in Figure 3l further displays reduced potential fluctuations and an elevated baseline level for the TPY‐DCA‐treated film. This suggests mitigated nanoscale potential variation across the CsPbI₂Br surface, supporting a more uniform local electronic landscape near the surface [45].

The optical and optoelectronic properties of CsPbI₂Br films with different TPY derivatives were further examined to assess the improvement in electronic uniformity induced by the modifiers. The UV–vis absorption spectra shown in Figure 4a, together with the corresponding Tauc plots in Figure S13, consistently reveal a similar absorption onset at 650 nm and an optical bandgap of 1.91 eV for all samples, indicating that the incorporation of TPY derivatives does not alter the fundamental band structure of CsPbI₂Br [46, 47]. Compared with the pristine film, the TPY‐DCA‐treated sample exhibits moderately enhanced absorption across the visible region, which is well correlated with its enlarged grain size and improved crystallinity. Sub‐bandgap disorder was quantitatively evaluated using Urbach energy (E _u) analysis (Figure S14). The pristine film displays a relatively large E _u of 40.85 ± 0.19 meV, reflecting a high density of band‐tail states [48]. In contrast, TPY‐COOH and TPY‐TCA lead to only partial reduction of E _u, whereas TPY‐DCA reduces E _u to 31.82 ± 0.08 meV, demonstrating more effective suppression of electronic disorder. Consistent with these optical characteristics, XRD patterns depicted in Figure 4b confirm that all CsPbI₂Br films maintain the cubic phase. Notably, the TPY‐DCA results in enhanced diffraction intensity and reduced full width at half maximum (FWHM) (Figure S15), reflecting improved crystalline ordering and reduced lattice disorder [49].

FIGURE 4.

(a) UV–vis absorption spectra of pristine and modified CsPbI₂Br films (Note: A. U. stands for absorbance unit); (b) XRD patterns of pristine and modified CsPbI₂Br films; Dark I–V characteristics of (c) electron‐only and (d) hole‐only CsPbI₂Br devices measured for pristine and TPY‐DCA‐treated samples; (e) DFT calculated formation energy of V_I; TRPL decay curves of CsPbI₂Br films deposited on (f) glass and (g) SnO₂ substrates; (h) Small‐bias J–V characteristics of electron‐only devices measured under AM 1.5G illumination; (i) UPS spectra of the pristine and TPY‐DCA‐modified CsPbI₂Br films.

The space‐charge‐limited current (SCLC) measurements were performed to quantitatively assess the effect of TPY‐DCA on the bulk defect density of CsPbI₂Br films. In the electron‐only architecture (Figure 4c and Figure S16a), TPY‐DCA treatment reduces the trap‐filled‐limit voltage (V _TFL) from 1.77 to 0.85 V, corresponding to a decrease in electron trap density (n _trap) from 4.66 × 10¹⁵ to 2.24 × 10¹⁵ cm⁻³, calculated using n _trap = 2εV _TFL/eL ², where ε denotes the permittivity of CsPbI₂Br, e is the elementary charge, and L represents the thickness of the CsPbI₂Br layers. A similar trend is observed in the hole‐only configuration (Figure 4d and Figure S16b), where V _TFL decreases from 1.69 to 0.94 V after TPY‐DCA modification, indicating a substantial reduction in n _trap of the hole. The concurrent suppression of both n _trap of electron and hole suggests an overall improvement in the bulk electronic quality of CsPbI₂Br [50, 51, 52]. These experimental results are consistent with DFT calculations of V_I formation energies (Figure 4e), which reveal that TPY‐DCA increases the energetic cost of V_I formation relative to the other TPY derivatives, providing thermodynamic evidence for the reduced n _trap observed experimentally.

Steady‐state and time‐resolved photoluminescence (SSPL/TRPL) measurements were performed to probe carrier recombination dynamics in CsPbI₂Br films. As shown in Figure S17a, the TPY‐DCA‐modified film delivers a higher PL intensity centered at 650 nm when deposited on glass, suggesting reduced non‐radiative recombination losses [53, 54]. Correspondingly, TRPL analysis reveals an increase in the average carrier lifetime (τ) from 38.71 to 94.73 ns after TPY‐DCA treatment (Figure 4f and Table S1), consistent with the reduced n _trap identified by SCLC measurements. To evaluate interfacial charge‐transfer behavior, SSPL and TRPL were further measured for CsPbI₂Br films deposited on SnO₂. Compared with the pristine film, the TPY‐DCA‐modified sample exhibits more substantial PL quenching (Figure S17b) accompanied by a shortened interfacial decay time from 16.69 to 2.41 ns (Figure 4g and Table S2), indicating more efficient electron extraction into the SnO₂ ETL [55, 56].

The small‐bias J–V characteristics in Figure 4h exhibit linear and symmetric responses around the origin, indicating near‐ohmic charge transport across the CsPbI₂Br/SnO₂ and CsPbI₂Br/PC₆₁BM interfaces. Upon TPY‐DCA incorporation, the steeper slope indicates enhanced charge transport, attributed to reduced trap‐limited scattering and more efficient carrier conduction [57, 58], which is consistent with the reduced n _trap extracted from SCLC and the prolonged τ revealed by TRPL. Ultraviolet photoelectron spectroscopy (UPS) provides complementary insight into the interfacial energetics upon additive incorporation. As shown in Figure 4i, all TPY‐based additives induce an increase in the secondary electron cutoff (E _cutoff), accompanied by a corresponding decrease in work function (WF) compared with the pristine film. Specifically, the WF decreases from 4.58 (pristine) to 4.07 eV for TPY‐DCA, while intermediate values are observed for TPY‐COOH (4.27 eV) and TPY‐TCA (4.46 eV). This systematic reduction in WF indicates an upward shift of the Fermi level (E _F) toward the conduction band, suggesting a more n‐type surface character. Such a shift is consistent with improved energetic alignment at the CsPbI₂Br/SnO₂ interface, which is expected to facilitate electron extraction and reduce interfacial carrier accumulation [59].

Building on the above structural, electronic, and interfacial optimizations enabled by multisite coordination, the photovoltaic performance of carbon‐cathode HTL‐free CsPbI₂Br PSCs was systematically evaluated. Figure 5a compares the J–V characteristics of pristine and additive‐modified devices. The pristine device delivers a champion PCE of 11.80%, with a V _oc of 1.228 V, a short‐circuit current density (J _sc) of 14.30 mA cm⁻², and an FF of 0.672. Upon incorporation of TPY‐DCA, the device performance is substantially improved, achieving a PCE of 16.02%, together with simultaneous increases in V _oc to 1.342 V, J _sc to 15.09 mA cm⁻², and FF to 0.791. Concentration‐dependent statistics in Figure S18 further identify 0.05% molar ratio of CsPbI₂Br precursor as the optimal additive loading for all derivatives, underscoring the importance of precise concentration control in maximizing coordination‐mediated passivation and device performance.

FIGURE 5.

(a) J–V curves and extracted photovoltaic parameters of pristine and modified devices; (b) EQE spectra and the corresponding integrated J _sc; (c) EIS Nyquist plots; (d) V _oc as a function of light intensity; (e) Mott‐Schottky plots of pristine and TPY‐DCA‐modified devices; (f) SPO measured at MMP; (g) Normalized power conversion efficiency (PCE) of unencapsulated devices stored under Ar atmosphere; (h) Thermal stability of unencapsulated devices aged under dry air.

The J _sc enhancement is further supported by the external quantum efficiency (EQE) spectra shown in Figure 5b. The TPY‐DCA‐modified device maintains EQE values above 85% across the 430–650 nm spectral range, delivering an integrated J _sc of 14.97 mA cm⁻² that closely agrees with the value extracted from J–V measurements. In contrast, the pristine device exhibits lower EQE and a reduced integrated J _sc over the same wavelength, confirming increased recombination losses and less efficient photocarrier collection [60]. Electrochemical impedance spectroscopy (EIS) measured under dark conditions provides additional insight into recombination‐related interfacial processes. As evidenced by the Nyquist plots in Figure 5c, TPY‐DCA increases the recombination resistance (R _rec) to 31.6 kΩ, compared with 6.8 kΩ for the pristine device, indicating suppressed interfacial recombination pathways [61]. Devices incorporating TPY‐COOH and TPY‐TCA exhibit intermediate R _rec values (Figure S19a), consistent with their comparatively weaker defect passivation effects. In parallel, TPY‐DCA yields the lowest series resistance (R _s), as reflected by the reduced high‐frequency intercept on the Z′ axis in Figure S19b, suggesting improved interfacial contact and transport quality [62].

The recombination was further examined through light‐intensity‐dependent V _oc analysis. As shown in Figure 5d, the pristine device exhibits an ideality factor (n _id) of 1.91, indicating that trap‐assisted Shockley‐Read‐Hall (SRH) recombination dominates device operation. Upon TPY‐DCA incorporation, n _id decreases to 1.25, suggesting a transition toward recombination close to bimolecular recombination, which is consistent with the effective suppression of deep trap states and non‐radiative losses [63]. Devices modified with TPY‐COOH and TPY‐TCA follow the same qualitative trend but retain higher n _id values (Figure S20a), reflecting less complete defect passivation. Complementary photoluminescence quantum yield (PLQY) measurements (Figure S20b) corroborate this interpretation, with the TPY‐DCA‐treated film exhibiting the highest PLQY, while the pristine film shows the lowest value, consistent with enhanced radiative recombination efficiency and reduced non‐radiative losses. Further insight into charge‐collection behavior is provided by the light‐intensity dependence of J _sc in Figure S21. TPY‐DCA elevates the power‐law exponent α from 0.892 to 0.982, shifting the dominant charge collection mechanism from trap‐assisted to diffusion‐assisted extraction [64].

The Mott–Schottky analysis (Figure 5e and Figure S22) reveals that TPY‐DCA increases the built‐in potential of the device. The extracted built‐in voltage (V _bi) rises from 1.19 V for the pristine device to 1.30 V after TPY‐DCA modification, indicating enhanced band bending at the CsPbI₂Br/SnO₂ interface. The increased V _bi promotes more efficient charge separation and drift, aligning with the higher V _oc observed in TPY‐DCA‐modified devices [65, 66]. Dark J–V measurements in Figure S23 demonstrate suppressed reverse‐bias leakage currents after TPY‐DCA incorporation, indicating reduced shunt pathways and improved interfacial selectivity associated with coordinated defect passivation [67, 68]. The stabilized power output (SPO) measurements (Figure 5f) show that the TPY‐DCA‐modified device maintains a steady PCE of 15.87% at a maximum power point (MPPT) voltage of 1.33 V, outperforming the pristine device (10.51% at 1.20 V). This enhancement is attributed to suppressed interfacial recombination and mitigated ion migration [69, 70]. Device statistics collected from 20 independent cells (Figure S24) show simultaneous improvements in the average J _sc, V _oc, and FF, together with narrower parameter distributions for TPY‐DCA‐treated devices, indicating enhanced fabrication reproducibility. Consistently, J–V hysteresis measurements (Figure S25) reveal a reduced hysteresis index (HI) after TPY‐DCA modification, in line with a lower density of mobile charged defects [71]. Beyond all‐inorganic CsPbI₂Br devices, the applicability of TPY‐DCA was further examined in MAPbI₃‐based PSCs fabricated under ambient air. As shown in Figure S26, incorporating TPY‐DCA increases PCE from 11.18% to 15.11%. Although MAPbI₃ and CsPbI₂Br differ in composition and defect chemistry, the consistent performance enhancement observed in both devices suggests that TPY‐DCA can effectively modulate defect states and facilitate charge transport in chemically distinct perovskite absorbers.

Device durability was evaluated under a series of environmental and thermal stress conditions to assess the operational stability imparted by TPY‐DCA coordination. Under inert storage (Ar, 20 ± 5°C), the pristine device exhibits continuous efficiency decay, retaining 64.6% of its initial PCE after 4500 h, whereas the TPY‐DCA‐modified device preserves 93.4% (Figure 5g). Consistently, no apparent phase evolution is observed in the TPY‐DCA‐treated films (Figure S27a,b), indicating suppressed intrinsic degradation under inert conditions. Under accelerated thermal aging in dry air (<10% RH, 60°C), the pristine device degrades to 50.7% of its initial PCE within 1500 h, while the TPY‐DCA‐treated counterpart maintains 79.3% even after 2500 h (Figure 5h). This improved thermal stability correlates with the suppressed formation of Pb⁰ and the retention of the original absorption profile (Figure S27c,d), suggesting a stabilized local Pb─X coordination environment under thermal stress. The TPY‐DCA modification markedly retards ambient degradation of unencapsulated devices (35 ± 10% RH, 20 ± 5°C), with the modified device retaining >50% of its initial PCE after 1000 h, whereas the pristine device retains >50% of its initial PCE only after 500 h (Figure S28a). This improvement stems from suppressed phase decomposition, as evidenced by the much weaker δ‐phase formation in XRD and the correspondingly sharper absorption edge, without distortion, compared to the pristine film (Figure S28b,c) [72]. In addition, water contact‐angle measurements show an increase from 45.4° to 74.5° upon TPY‐DCA incorporation (Figure S29), indicating an enhanced surface hydrophobicity that limits phase decomposition induced by moisture ingress. Collectively, these observations confirm that TPY‐DCA improves the material stability of carbon‐cathode HTL‐free CsPbI₂Br PSCs, which underpins their substantially prolonged operational lifetimes under diverse stress conditions.

3. Conclusions

In summary, a multisite coordination‐driven molecular strategy employing TPY derivatives is presented to address key limitations of defect states, interfacial energetics, lattice distortion, and carrier transport in carbon‐cathode HTL‐free CsPbI₂Br PSCs. By balancing coordination capability with steric accessibility, TPY‐DCA enables cooperative N/O coordination at defect‐rich surfaces, grain boundaries, and buried interfaces. This interaction is associated with reduced undercoordinated Pb²⁺ and Sn⁴⁺ defects, improved stability of V_I and V_O, a more uniform lattice strain distribution, and a more favorable interfacial energy environment. As a result, CsPbI₂Br films exhibit improved structural uniformity, reduced electronic heterogeneity, and more efficient charge extraction across critical interfaces. These improvements translate into enhanced device performance, delivering a PCE of 16.02% together with reduced J–V hysteresis, improved device reproducibility, and extended operational stability under inert, thermal, and humid conditions. The successful application of this approach to MAPbI₃ devices further indicates the broader applicability of this molecular strategy. Overall, this work demonstrates that multisite coordination molecules provide a practical route to simultaneously regulate structural and electronic properties in perovskite photovoltaics, offering a valuable framework for the development of efficient, scalable, and stable carbon‐cathode HTL‐free PSCs.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: smll73857‐sup‐0001‐SuppMat.doc.

Data Availability Statement

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