Unveiling the doping effect of mixed-halide CsPb(Br_1−nX_n)₃ (X = I, Cl) single crystals toward high-sensitivity radiation detection†

Rong Wu^a, Yansong Yue^bc, Aosheng Zhang^bc, Jing Wei^a, Fangze Liu*^bc and Hongbo Li*^a
aBeijing Key Laboratory of Construction-Tailorable Advanced Functional Materials and Green Applications Experimental Center of Advanced Materials, School of Materials Science and Engineering, Beijing Institute of Technology, Beijing 100081, China. E-mail: hongbo.li@bit.edu.cn
^bSchool of Interdisciplinary Science, Beijing Institute of Technology, Beijing 100081, China. E-mail: fliu@bit.edu.cn
^cSchool of Interdisciplinary Science, Beijing Institute of Technology, Zhuhai 519088, China

First published on 15th July 2025

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1. Introduction

Metal halide perovskites have emerged as promising semiconductor materials for optoelectronic applications, spanning photovoltaics,^1–3 light-emitting diodes^4–6 and radiation detectors.^7–12 Notably, perovskite X-ray detectors have shown higher sensitivity and lower detection limits compared with conventional selenium (Se) and cadmium zinc telluride (CdZnTe) detectors. These outstanding characteristics primarily benefit from their unique optoelectronic properties, including tunable bandgaps (E_g = 1.5–3 eV), heavy atom compositions, high resistivities (R = 10⁷–10¹¹ Ω cm) and large mobility–lifetime products (μ–τ = 10⁻⁵–10⁻² cm² V⁻¹).^9,13,14 Despite the rapid progress in perovskite X-ray detectors, practical implementations are restricted by inherent material limitations, including environmental instability, moderate resistivity and pronounced ion migration. Compared to polycrystalline films, perovskite single crystals (PSCs) exhibit several advantages through their grain-boundary-free structure, demonstrating enhanced environmental stability, reduced defect density (10⁹–10¹⁰ cm⁻³ vs. 10¹⁴–10¹⁶ cm⁻³ in polycrystalline films) and extended carrier diffusion lengths.^15,16 These superior properties position PSCs as particularly interesting materials for radiation detectors. Consequently, the development of large-area, high-quality PSCs has become a critical research focus for next-generation detector technologies.

Recent advances in the synthesis of PSCs have yielded several material systems for photodetection applications, including organic–inorganic hybrid perovskites, all-inorganic perovskites and lead-free perovskites. However, the intrinsic instability of organic constituents in hybrid perovskites poses fundamental limitations for their practical applications. In contrast, inorganic perovskites (CsPbX₃, X = Cl, Br, and I) address these issues through enhanced environmental stability and higher atomic numbers. Among the CsPbX₃ family, CsPbI₃ suffers from phase transition under ambient conditions, whereas CsPbCl₃ has relatively large exciton binding energy and less optimal charge transport properties.¹⁷ These limitations position CsPbBr₃, which benefits from its balanced bandgap (2.3 eV), high X-ray attenuation coefficient and good environmental stability, as the most variable candidate for X-ray detection applications. The crystal growth methods have a significant impact on the detector performance. The CsPbBr₃ SCs grown using the Bridgman method exhibit exceptionally high purity and centimeter-scale dimensions, enabling gamma-ray spectroscopy with high energy resolution.¹⁸ In addition to the Bridgman method, solution growth methods, including inverse-temperature crystallization (ITC), temperature-lowering crystallization (TLC), the slow solvent evaporation method (SSE) and antisolvent vapor assisted crystallization (AVC), also offer several advantages due to their lower cost and simpler process.¹⁹ The ITC method, which has been shown to grow PSCs with performance comparable to those grown by the Bridgeman method, utilizes the negative temperature coefficient of solubility, facilitating crystal growth under metastable conditions with precisely controlled supersaturation levels.^18,20–24

The performance of CsPbBr₃ X-ray detectors is significantly limited by their relatively low resistivity and severe ion migration, leading to a high detection limit and fast dark current drift.²⁵ Various strategies have been developed to address these issues, including preparing 2D–3D heterostructured perovskites, surface passivation, voltage cycling technology and constructing asymmetric electrodes.^26–31 However, these strategies mainly focus on improving the detector performance, without fundamentally increasing resistivity or inhibiting ion migration. For perovskite semiconductors, their bandgaps can be effectively adjusted through engineering the halogen composition. For example, CsPbBr_3–3nX_3n (X = Cl, I) SCs have been successfully prepared using the modified Bridgman method with tunable bandgaps between 2.88 eV and 1.90 eV by varying the halide composition.³² Moreover, doping A-site cations or halogens can improve the stability and resistivity of PSCs. Previous studies have shown that the undesirable phase transition and phase segregation in FAPbI₃, which increase the dark current and reduce material stability, can be mitigated by incorporating smaller ions to release lattice stress, including methylammonium (MA⁺), cesium (Cs⁺) and bromine (Br⁻). The inch-sized high-quality triple-cation mixed-halide FA_0.85MA_0.1Cs_0.05PbI_2.55Br_0.45 (FAMACs) SC exhibited suppressed ion migration, low defect density and improved environmental stability. These FAMACs SC detectors showed high sensitivity and a low detection limit.³³ Similarly, doping CsPbBr₃ SCs with I⁻ greatly increased the resistivity, and the CsPbBr_2.9I_0.1 SC achieved a high sensitivity of 6.3 × 10⁴ μC Gy_air⁻¹ cm⁻² with a low detection limit.²⁵ These results demonstrate that high-quality and large-size PSCs can be grown by composition engineering with improved resistivity, inhibited ion migration and increased material stability. While previous studies have separately shown that both iodine doping and chlorine doping at levels of 3.3% can significantly improve the crystal quality and detector performance,^25,34 the effects of light halide doping on the electron and hole transport properties of solution-grown CsPb(Br_1−nCl_n)₃ and CsPb(Br_1−nI_n)₃ SCs with n < 0.033 remain unresolved.

In this study, we report the synthesis of lightly doped mixed-halide perovskite CsPb(Br_1−nX_n)₃ (X = Cl, I) SCs with high crystallinity and compositional uniformity using the ITC method. While Cl⁻ can be effectively doped into the crystals, I⁻ doping was limited due to its large ionic radius. Through a comprehensive study of both hole and electron transport properties by constructing symmetric detectors with high and low work function contacts, we found that all SCs showed relatively lower hole resistivities (∼10⁸ Ω cm) than electron resistivities (∼ 10¹⁰ Ω cm), indicating that the majority of defects in CsPb(Br_1−nX_n)₃ SCs are halide vacancies. Cl⁻ doped SCs showed partly inhibited ion migration, while the higher hole resistivity of I⁻ doped SCs demonstrates the effectiveness of suppression of ion migration by I⁻. The high hole conductivity can induce current gain, leading to higher detection sensitivity. Therefore, the Au/CsPb(Br_0.98Cl_0.02)₃/Au hole detector achieved the highest sensitivity of 6388 μC Gy_air⁻¹ cm⁻². With low work function contacts, the electron dark current was substantially reduced by two orders of magnitude compared with the hole current due to the elimination of hole contributions, and the dark current drift was also effectively suppressed. The carrier transport dynamics was further revealed by α-particle spectroscopy, proving that Cl⁻ doping reduces the defect density, increases the carrier lifetime and achieves improved charge collection efficiency. In contrast, although trace-amount I⁻ doping can better inhibit ion migration, I⁻ doped SCs showed lower μ–τ products as well as slower carrier transport.

2. Results and discussion

CsPb(Br_1−nX_n)₃ (where X = Cl, I) SCs were synthesized using the ITC method with different percentages of Cl⁻ and I⁻ precursors (2%, 5% and 10%), which are referred to as Cl_0.02 (CsPb(Br_0.98Cl_0.02)₃), Cl_0.05 (CsPb(Br_0.95Cl_0.05)₃), Cl_0.1 (CsPb(Br_0.9Cl_0.1)₃), I_0.02 (CsPb(Br_0.98I_0.02)₃), I_0.05 (CsPb(Br_0.95I_0.05)₃) and I_0.1 (CsPb(Br_0.9I_0.1)₃). The photographs of typical SCs are shown in Fig. 1a. The geometrically faceted morphology and high optical transparency of these crystals confirm their phase purity and structural integrity. In Fig. 1b and Fig. S1, ESI,† X-ray diffraction (XRD) patterns verify the orthorhombic CsPbBr₃ structure (Pnma space group, PDF#97851) composed of corner-sharing PbBr₆⁴⁻ octahedra with Cs⁺ cations occupying interstitial sites.³⁵ The negligible peak shifts across all compositions indicate minimal lattice distortion, suggesting limited halogen incorporation into the CsPbBr₃ crystal structure.³⁶ Note that the broad peak around 17° is from the substrate.

In order to study the actual doping amount and compositional uniformity of CsPb(Br_1−nX_n)₃ SCs, energy-dispersive spectroscopy (EDS) was performed. The EDS mappings of Cl_0.1 SCs confirm the uniform distribution of Cs, Pb, Br and Cl elements (Fig. S2, ESI†). However, the measured Cl/Br ratio deviates substantially from that of the precursors, where the n values of Cl_0.02, Cl_0.05, Cl_0.1, Cl_0.15 and Cl_0.2 were actually 0.008, 0.016, 0.032, 0.04 and 0.055, respectively. The corresponding doping ratios calculated by dividing the measured percentage with the precursor value were 40.0%, 32.0%, 32.0%, 26.7% and 27.5%, respectively (Fig. S3, ESI†). For simplicity, we keep using the precursor ratios to represent the SCs. In addition, the content of I⁻ is hardly detected, suggesting that I⁻ is more difficult to be incorporated into CsPbBr₃ crystals.³⁶ Nevertheless, the effective doping of Cl⁻ in CsPbBr₃ SCs fills the Br⁻ vacancies and reduces the trap density, which inhibits ion migration and crystal phase transition.²⁵

The chemical states of the SCs were analyzed by X-ray photoelectron spectroscopy (XPS) for Cs 3d, Pb 4f and Br 3d peaks. As shown in Fig. 1c and Fig. S4, ESI†, an increase in Cl⁻ content induced systematic reductions in the binding energies of Cs 3d (0.3 eV), Pb 4f (0.3 eV) and Br 3d (0.3 eV) XPS peaks. While Cl has higher electronegativity than Br, the decreasing binding energies can be attributed to the effective passivation of Br⁻ vacancies by Cl⁻.^37–39 These observations confirm the effective doping of Cl⁻ in CsPbBr₃ SCs. I⁻ doped SCs also exhibit lowered binding energies of Cs 3d, Pb 4f and Br 3d XPS peaks due to the lower electronegativity of I compared to Cl.

Steady-state photoluminescence (PL) spectra displayed that the peaks were blue-shifted from about 532.4 nm to 528.7 nm as the Cl⁻ doping amount increases to 0.1 (Fig. 1d). In contrast, the PL peak only red-shifted by 1 nm for I⁻ doped SCs, demonstrating the very low doping of I⁻. PL mapping was measured to investigate the composition uniformity. The deviations of the PL position for the cross sections of Cl_0.05, Br₃ and I_0.05 were all around 0.4 nm (Fig. 1e–g). These narrow PL position distributions reveal uniform doping in CsPbBr₃ SCs. Furthermore, the full width at half maximum (FWHM) values of these SCs were 16–17 nm (Fig. S5, ESI†), which are comparable to those of SCs grown by the Bridgman method in previous reports.¹⁸ Time-resolved photoluminescence (TRPL) curves of all SCs can be well fitted by a biexponential decay function (Fig. S6, ESI†), which yields a fast decay lifetime (τ₁) and a slow decay lifetime (τ₂), corresponding to surface trap-assisted recombination and bulk exciton recombination, respectively.^30,40 The calculated average decay lifetime (τ_av, see the ESI†) peaked at 62.35 ns for Cl_0.05, while I⁻doped SCs generally exhibit significantly shorter lifetimes (τ_av < 35 ns), suggesting faster electron–hole recombinations in I⁻ doped SCs. The UV-visible absorption spectra showed sharp band edge absorption at around 560 nm, and the Tauc plot matched the bandgaps of about 2.21–2.22 eV (Fig. S7, ESI†).

To minimize the dark current and improve the signal-to-noise ratio in X-ray detectors, semiconductors generally need to exhibit high resistivity greater than 10⁹ Ω cm. The hole resistivities of CsPb(Br_1−nX_n)₃ SCs were evaluated by depositing high work function Au contacts to form a vertical device structure of Au/CsPb(Br_1−nX_n)₃/Au. At a bias ranging from 0 to 100 V, these symmetrical devices showed ohmic contact behaviors (Fig. S8, ESI†). The measured hole resistivities of CsPb(Br_1−nX_n)₃ were about 10⁸ Ω cm. The X-ray induced photocurrent was measured using an X-ray tube with a tungsten target operating at an acceleration voltage of 40 kV and varying tube current between 30 μA and 90 μA (Fig. 2a and b, Fig. S9a and b, ESI†). All SCs were biased under the same electric field of 50 V cm⁻¹. The dark current of Br₃ increased by 394 nA cm⁻² over 200 s, whereas for Cl_0.02, it increased by only 125 nA cm⁻². The reduction in dark current drift for Cl⁻ doped SCs compared with CsPbBr₃ SCs is attributed to the effective suppression of Br⁻ vacancy defects by Cl⁻, which inhibits ion migration.²⁵ In contrast, the I_0.05 and I_0.1 SCs exhibited the lowest dark current of less than 15 nA cm⁻² with a minimal drift of a few nA cm⁻². The much lower current level of I⁻ doped SCs can be attributed to the suppressed formation of I⁻ and Br⁻ vacancies, effectively reducing the hole defect density. The X-ray photocurrent measured at varying dose rates revealed a strong linear relationship, indicating excellent X-ray response (Fig. 2c, Fig. S9c, ESI†). The X-ray sensitivities evaluated using the hole currents of CsPb(Br_1−nX_n)₃ SCs were 179 μC Gy_air⁻¹ cm⁻² (Cl_0.1), 1066 μC Gy_air⁻¹ cm⁻² (Cl_0.05), 6388 μC Gy_air⁻¹ cm⁻² (Cl_0.02), 875 μC Gy_air⁻¹ cm⁻² (Br₃), 560 μC Gy_air⁻¹ cm⁻² (I_0.02), 479 μC Gy_air⁻¹ cm⁻² (I_0.05) and 251 μC Gy_air⁻¹ cm⁻² (I_0.1) (Fig. 2d). The enhanced electron affinity of Cl⁻ effectively attracts the photogenerated electrons produced in the SCs, reducing the radiation recombination rate, which is beneficial for improving the hole sensitivity of the detector. For I⁻ doped SCs, the reduced hole density also shortened the carrier lifetime, resulting in lower sensitivity, which agrees with the TRPL results.

To investigate electron transport characteristics, we fabricated Ga/CsPb(Br_1−nX_n)₃/Ga vertical devices to measure electron current, and the electron resistivities were determined to be 1.43 × 10¹⁰ Ω cm (Cl_0.1), 7.60 × 10⁹ Ω cm (Cl_0.05), 1.27 × 10¹⁰ Ω cm (Cl_0.02), 1.87 × 10¹⁰ Ω cm (Br₃), 8.79 × 10⁹ Ω cm (I_0.02), 3.87 × 10¹⁰ Ω cm (I_0.05) and 5.24 × 10¹⁰ Ω cm (I_0.1) (Fig. 3a and Fig. S10, ESI†). High electron resistivities are advantageous for reducing the dark current and improving the signal-to-noise ratio of the devices. The much higher electron resistivities compared to hole resistivities indicate that the major defects in CsPb(Br_1−nX_n)₃ SCs are halide vacancies. The charge transport properties of electrons can be evaluated by measuring the X-ray photocurrent as a function of applied voltage at an X-ray tube acceleration voltage of 90 kV and a current of 20 μA (Fig. S11a, ESI†). The μ–τ products of PSCs were derived by fitting the photocurrent with the Hecht equation:

where I is the measured photocurrent, I₀ is the saturated photocurrent, V is the applied bias, L is the thickness of the sample, and s is the surface recombination rate. The obtained μ–τ products were 5.99 × 10⁻⁴ cm² V⁻¹ (Cl_0.1), 8.35 × 10⁻⁴ cm² V⁻¹ (Cl_0.05), 5.70 × 10⁻⁴ cm² V⁻¹ (Cl_0.02), 5.06 × 10⁻⁴ cm² V⁻¹ (Br₃), 7.73 × 10⁻⁴ cm² V⁻¹ (I_0.02), 2.22 × 10⁻⁴ cm² V⁻¹ (I_0.05) and 1.15 × 10⁻⁴ cm² V⁻¹ (I_0.1) (Fig. 3b). These results demonstrate that Cl⁻ doping improves the charge transport properties of CsPbBr₃ SCs by increasing the electron lifetime, which is also shown by the prolonged TRPL lifetime. The I_0.05 and I_0.1 SCs, in contrast, exhibited much lower μ–τ products due to the shorter electron lifetime, which agrees with the hole transport case. The high μ–τ product of the I_0.02 SC can be explained by the very low doping concentration, and the I_0.02 SC is close to intrinsic CsPbBr₃.

The X-ray detection performances of Ga/CsPb(Br_1−nX_n)₃/Ga devices were evaluated at varying X-ray tube acceleration voltages of 40 kV, 60 kV and 90 kV. The high electron resistivity reduced the dark current of the device to below 100 nA cm⁻² under an electric field of 500 V cm⁻¹ (Fig. 3c and Fig. S11b and c, ESI†), lower than 1/100 of the dark current density observed with high work function Au contacts (Fig. S12, ESI†), where the hole dark current densities of Cl_0.05 and Br₃ increased to 60 and 30 μA cm⁻², respectively, after 150 s when measured under the same electric field. This significant hole current drift overwhelmed the X-ray induced photocurrents.

The electron sensitivities were calculated by plotting the X-ray photocurrent as a function of dose rate (Fig. 3c–f, Fig. S11b and c, ESI†). The I_0.05 and I_0.1 SCs showed much lower sensitivities compared to intrinsic and Cl⁻ doped SCs owing to their relatively lower μ–τ products. The higher sensitivity of I_0.02 agrees with its higher μ–τ product (Fig. 3g and h). All SCs exhibited the highest sensitivities at an X-ray tube voltage of 90 kV, which were around 400 μC Gy_air⁻¹ cm⁻² for the Cl⁻ doped and intrinsic CsPbBr₃ SCs (Fig. 3i). To further reveal the charge transport dynamics inside CsPb(Br_1−nX_n)₃ SCs, we use α-particle spectroscopy to probe the carrier transit time. The Ga/CsPb(Br_1−nX_n)₃/Ga devices were biased under an electric field of 500 V cm⁻¹ with the cathode grounded. The anode was connected to a preamplifier and a multi-channel analyzer. An ²⁴¹Am source emitting 5.48 MeV α-particles was placed within 5 mm of the cathode to minimize α-particle attenuation in air. Due to the short penetration depth of α-particles in CsPbBr₃ (∼10 μm), all electrons generated by the absorption of an α-particle at the cathode can be viewed as travelling toward the anode simultaneously, which generates an electric pulse to the preamplifier. Note that α-particle spectroscopy cannot be performed for hole devices due to the fast drifting dark current. The width and height of the pulse rising edge therefore correspond to the moving speed and amount of the radiation-induced electrons, respectively. A higher pulse height corresponds to higher charge collection efficiency, and a shorter rising time represents a faster moving speed. Consequently, by analyzing the pulse rise time, we can quantitatively extract the electron travelling time across the crystal. In Fig. 4a, the energy spectra of 5.48 MeV α-particles gradually shift to higher channel numbers after Cl⁻ and I⁻ doping, indicating higher charge collection efficiency. To further analyze the carrier behavior in SCs, typical pulses of the full-energy peaks are shown in Fig. 4b. The increasing trend of the rise time from Cl⁻ doped SCs to I⁻ doped SCs can be clearly observed. More quantitative analysis was performed by plotting the statistics of the pulse rise time in Fig. 4c, and the average rise times were determined to be 2.4 ± 1.4 μs (Cl_0.1), 2.0 ± 1.2 μs (Cl_0.05), 4.7 ± 2.7 μs (Cl_0.02), 5.4 ± 2.9 μs (Br₃), 6.0 ± 3.2 μs (I_0.02), 6.3 ± 2.6 μs (I_0.05) and 5.0 ± 1.9 μs (I_0.1) (Fig. 4d). Notably, Cl⁻ doped SCs exhibited shorter average rise times and narrower distributions, indicating optimized carrier transport pathways through the bulk crystal with suppressed defect-related carrier trapping. The reduced defect density achieved through Cl⁻ doping accounts for the improved α-particle spectral performance, consistent with the observed enhancements in X-ray detection sensitivity and μ–τ products for Cl⁻ doped SCs. On the other hand, the slower electron transport in I⁻ doped SCs agrees with their lower μ–τ products, which may originate from the combined effect of decreased electron mobility and shortened lifetime.

3. Conclusion

In summary, we successfully synthesized lightly doped mixed-halide CsPb(Br_1−nX_n)₃ SCs with high compositional uniformity and high crystallinity using the ITC method. The Cl⁻ can be more effectively doped into CsPbBr₃ SCs than I⁻. Cl⁻ doped CsPb(Br_1−nCl_n)₃ SCs exhibited longer exciton lifetimes than I⁻ doped SCs. Through a comprehensive study of both hole and electron transport properties, we found that all SCs showed relatively lower hole resistivities (∼10⁸ Ω cm) than electron resistivities (∼10¹⁰ Ω cm), indicating that the majority of defects in CsPb(Br_1−nX_n)₃ SCs are halide vacancies. Cl⁻ doped SCs showed partly inhibited ion migration, while the higher hole resistivities of I⁻ doped SCs demonstrate the effectiveness of suppression of ion migration by I⁻. The high hole conductivity can induce current gain, leading to higher detection sensitivity. Therefore, the Au/CsPb(Br_0.98Cl_0.02)₃/Au hole detector achieved the highest sensitivity of 6388 μC Gy_air⁻¹ cm⁻². Compared to hole transport, the much higher electron resistivity is beneficial for reducing the dark current of the detectors, and dark currents of electrons measured from Ga/CsPb(Br_1−nX_n)₃/Ga detectors were more than 100 times lower than those of holes, indicating that the low work function contact of Ga can effectively suppress the dark current drift caused by ion migration. The carrier transport dynamics was further revealed by α-particle spectroscopy, proving that Cl⁻ doping reduces the defect density, increases the carrier lifetime and achieves improved charge collection efficiency. In contrast, although trace-amount I⁻ doping can better inhibit ion migration, I⁻ doped SCs showed lower μ–τ products as well as slower carrier transport.

4. Experimental section

4.1 Materials and chemicals

Chemicals were used as obtained: cesium chloride (CsCl, 99.0%), cesium bromide (CsBr, 99.5%), cesium iodide (CsI, 99.9%) and lead bromide (PbBr₂, 99.0%) were purchased from Aladdin and Energy-Chemical. N,N-Dimethylformamide (DMF, super dry, 99.9%) and dimethyl sulfoxide (DMSO, super dry, 99.8%) were purchased from J&K. All chemicals were used as received without further purification.

4.2 Crystal growth

A total of 10 mmol of CsBr and CsCl or CsI, together with 20 mmol of PbBr₂, were dissolved in DMSO (10 mL) and stirred at room temperature for 2–3 hours. After all chemicals were dissolved, the solution was then filtered with a 0.45 μm filter and transferred into a 20 mL vial. The precursor solution was placed in a silicon oil bath at 95 °C and then heated at a rate of 1 °C per 6 hours for 24 hours. After oversaturation, many small crystals appeared, which were used as seeds for growing large crystals. The filtrate was transferred into a new vial, where a seed was placed at the bottom. The temperature was increased from 95 °C to 120 °C at a rate of 1 °C per 6 hours. In the process of slow heating, the seeds gradually grew into large-sized single crystals.

4.3 Characterization measurements

Author contributions

F. Liu and H. Li conceived the project. R. Wu synthesized single crystals. R. Wu, Y. Yue and A. Zhang fabricated and characterized detectors. J. Wei assisted with data analysis. R. Wu and F. Liu wrote the manuscript with input from all authors.

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

This work was supported by the National Natural Science Foundation of China (22105018, 52372135, 22379017, 22179009, and U22A2072), the Fundamental Research Funds for the Central Universities and the Beijing Institute of Technology.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d5nr01246e

This journal is © The Royal Society of Chemistry 2025