Construction of nano-lamellar expressways and multidimensional defects to realize the decoupling of carrier–phonon transport in BiSbSe_1.25Te_1.75

Abstract

BiSbSe_1.25Te_1.75, a typical multi-layered compound, has great potential for use in the manufacture of high-efficiency thermoelectric conversion devices due to its ability to be fabricated with p–n junctions of identical chemical composition by defect engineering. However, the thermoelectric properties of n-type BiSbSe_1.25Te_1.75 remain limited due to its poor electrical transport properties. Herein, we report an effective strategy to decouple its electrical and thermal transport properties, which can be realized by simple hot deformation of BiSbSe_1.25Te_1.75. Nanoscale lamellar structures with large surface areas and strongly preferred orientation formed by preferred growth along the ab planes provide expressways for electron transport. These structures are beneficial for promoting S while maintaining high σ because the expressways will effectively reduce the sacrifice through μ_H. Meanwhile, multidimensional defects are also introduced into samples by hot deformation, evoking strong scattering locations for phonons of different frequencies. Benefiting from the decoupling of carrier–phonon transport via hot deformation, a high average ZT value of 0.53 from 323 to 550 K (∼112% increase) and a high ZT value of 0.60 at 470 K (∼107% increase) are achieved in BiSbSe_1.25Te_1.75. This work undoubtedly paves the way for the utilization of TE materials with identical chemical composition in the fabrication of well-matched p–n junctions.

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Introduction

Thermoelectric materials are promising candidates for addressing the increasing demand for energy consumption and achieve carbon-neutral targets, which could be attributed to their advantages such as direct heat-to-electricity conversion capability, silent and reliable operation, and sustainable green technology without the production of undesirable and harmful substances. The energy conversion efficiency of TE materials is mainly evaluated by its dimensionless figure of merit, given by ZT = S²σT/κ, where S is the Seebeck coefficient, σ is the electrical conductivity, T is the absolute temperature and κ is the total thermal conductivity, which generally includes the carrier contribution κ_e and the lattice contribution κ_L. The quantities S, σ, and κ_e are considered to be coupled via the carrier concentration n.^1,2 The value of κ_L is only closely related to the heat-carrying phonon transport process. Therefore, two main efficacious strategies have been proposed to optimize and decouple electrical and thermal transport. One is to enhance the power factor, PF = S²σ, by band structure modification via tuning the resonance level and band convergence; the other is to suppress the κ_L value with no obvious degradation of electrical properties by introducing hierarchical architecture, disordered ions, and phonon engineering.

TE devices, which can directly realize waste-heat recovery and solid-state refrigeration, have been constructed from an array of n- and p-type TE materials with varying chemical compositions. Although many TE materials have attained the high ZT value of >1.0,^3–9 TE devices with high-performance and homogeneous structures are still lacking and less frequently reported, due to the discrepancy in structural and chemical compositions between n- and p-type TE materials. Accordingly, it is imperative to create n- and p-type thermoelectric materials with similar and well-matched systems.

As promising TE materials, typical multiple-layered compounds have produced a surge of interest in the thermoelectric field because of their intrinsically ultralow κ and excellent TE performance.^10–16 More recently, based on a particular layered structure, we found that p- and n-type TE materials with the same nominal chemical composition can be obtained by point-defect engineering in multiple-layered BiSbSe_1.5Te_1.5.¹⁷ Furthermore, an enhanced TE performance for n-type BiSbSe_1.25Te_1.75 by employing entropy engineering has also been reported in our previous work.¹⁸ Although a peak ZT of 0.54 at 475 K and a remarkable average ZT value of 0.45 (over 300–550 K) are achievable for the n-type BiSbSe_1.25Te_1.75 sample by entropy engineering, it is highly significant that a productive approach is designed to further optimize the ZT value for the n-type polycrystalline counterpart.

For TE materials composed of multiple-layered compounds, it is possible to enhance the PF to a level close to that of a single-crystal sample by grain alignment.¹⁹ Manipulating the grain alignment of TE materials by introducing hot deformation can tune the grain microstructure and promote the performance of layered-compound TE materials. The “top-down” hot deformation approach has been developed to improve the TE performance of p-type Bi_0.5Sb_1.5Te₃ by inducing in situ nanostructures and high-density lattice defects.²⁰ Meanwhile, the TE performance of n-type Bi₂Te₃ can also be distinctively improved by the hot deformation approach owing to its multi-scale microstructural effects and recrystallization-induced local nanostructures.²¹ The enhanced preferential orientation induced by the hot deformation approach can improve maximum ZT from 0.85 to 1.04 at 398 K in n-type Bi₂Te_2.7 Se_0.3¹⁹ and maximize PF at 923 K from 6.3 to 8.1 μW cm⁻¹ K⁻² in BiCuSeO.²² Based on these studies, it is confirmed that hot deformation can improve dramatically the performance of layered-compound TE materials.

In this work, repeated hot deformation has been employed to optimize the TE properties of n-type BiSbSe_1.25Te_1.75, and materials are denoted HDx (x = 0, 1, 2, and 3) bulk samples. It was found that nanoscale lamellar structures with large surface areas and strongly preferred orientation facilitate the electrical transport by providing expressways for carriers, giving rise to an enhanced PF of 1062.8 μW m⁻¹ K⁻² at 400 K. Meanwhile, various and abundant multiscale distortions are also introduced into samples subjected to hot deformation, evoking strong scattering for phonons of different frequencies. Overall, the mechanism of simultaneous optimization for electrical and thermal transport properties is shown in Fig. 1. Therefore, the electrical and thermal transport properties have been realized and effectively decoupled by hot deformation, leading to remarkable improvements of TE properties in n-type BiSbSe_1.25Te_1.75. As a consequence, a high ZT value of 0.60 at 470 K with an average ZT of 0.53 (over 323 to 550 K) for the n-type HD3 sample was achieved. This work demonstrates an innovative route to decouple electron and phonon transport by controlling the preferred orientation of grains via hot deformation for developing efficient thermoelectric materials.

Fig. 1 Schematic illustration of electron and phonon transport processes.

Results and discussion

Phase and microstructure

Fig. 2 shows the XRD patterns of the powder and HD_x (x = 0, 1, 2, and 3) bulk samples of BiSbSe_1.25Te_1.75. Compared with the powder sample, the peak intensities of the HD_x bulk samples vary significantly along (00l) faces. The great discrepancy in the diffraction intensities of these XRD patterns indicates a significant change in the grain orientations before and after hot deformation. The orientation degree F values of (00l) planes for the bulk HD_x samples were calculated from XRD patterns and formulas (S1) and (S2).†^23–26 The values of F = 0.19, 0.20, 0.25, and 0.28 are obtained for the HD_x (x = 0, 1, 2, and 3) bulk samples, respectively. It is apparent that hot deformation can cause the alignment of grains and lead to a highly oriented texture in the samples.^22,27

Fig. 2 XRD patterns of the BiSbSe_1.25Te_1.75 samples.

To evaluate the distribution of crystallographic orientations and demonstrate the formation of texture in the samples after hot deformation, the (006), (015), (1010), and (0018) pole figures of HD0 and HD3 are depicted in Fig. 3a. The noticeable enhancement along the (00l) preferred orientation is detected for the HD3 sample compared to that of HD0. Additionally, inverse pole figures (IPFs) also display a concentrated distribution that is higher along the (00l) direction for HD3, confirming again that strong (00l) textures are generated after hot deformation, as shown in Fig. 3b. Based on the pole figures, the orientation distribution function (ODF) can be calculated as shown in Fig. 3c. The HD3 sample shows a much stronger texture than that of HD0 and the texture strength initially displays a decrease followed by an increase in the range of 0°–90°. It is reasonable to conclude that the higher hot pressure temperature and larger plastic deformation can effectively promote the preferred growth of grains. All these results correspond to the variation in the XRD patterns, as given in Fig. 2.

Fig. 3 (a) Pole figures along the (006), (015), (1010), and (0018) directions, (b) inverse pole figures (IPFs), and (c) orientation distribution function (ODF) patterns of the HD0 and HD3 samples, respectively.

The morphologies of the freshly fractured surfaces and the corresponding reconstructed 3D images of the BiSbSe_1.25Te_1.75 bulk samples after different hot deformation times are presented in Fig. 4. As shown in the reconstructed 3D morphologies for the bulk samples (Fig. 4b, d, f, and h), the darker contrasts indicate a higher position of the corresponding fracture surface. With increasing hot deformation time, the stacked sheets of grains with a layered structure and distinct stripes are clearly observed. The corresponding reconstructed 3D morphologies reveal that undulating “peaks” are also gradually formed and tend to arrange along a specific orientation, demonstrating the preferred growth of grains during hot deformation. Moreover, the shape of the “peaks” gradually change from scattered columns into continuous flat layers. These grains prefer growth along the ab planes, which can be attributed to lateral flow during hot deformation¹⁹ and this lateral flow furthermore makes the nanoscale lamellae become thinner due to shear forces. These nano-lamellar structures can act as effective centers for strong phonon scattering, leading to a low thermal conductivity for the samples. Moreover, compared to the grain size of HD0, apparent grain growth can be observed as being driven due to the high temperature after hot deformation.^1,22,28

Fig. 4 (a), (c), (e), and (g) SEM images of the fractured surface and (b), (d), (f), and (h) the corresponding reconstructed 3D morphologies of all the HD_x (x = 0, 1, 2, and 3) bulk samples, respectively.

Considering the potential effects of defects on TE performance, TEM investigations were carried out for HD3. As shown in Fig. 5a, the local nanoprecipitates can be easily found in HD3, which is consistent with our previous works.¹⁷ The corresponding energy dispersive spectroscopy (EDS) maps of the enlarged region indicated in Fig. 5a is displayed in Fig. 5b, revealing that elemental Sb is remarkably enriched compared with other elements. The Sb element is more enriched than the Bi, Te, and Se elements in the matrix, which can be attributed to the considerable discrepancy in the formation energy of anti-site defects, E_AS, and anion vacancies, E_V.^21,29 The largest E_V(Sb–Te) and lowest E_AS(Sb–Te) values between the Sb cation and the Te anion make it more difficult for the Sb atoms occupying the Te sites to diffuse back into their original sublattices, and some Sb atoms are thereupon preserved. As a consequence, the Sb nanoparticle is precipitated in the n-type BiSbSe_1.25Te_1.75 samples, which is also reported in our previous work.^17,18 The high-angle annular dark-field HAADF-STEM image shown in Fig. 5c reveals that the hot deformation introduces high-density dislocations and numerous nanoscale-distorted regions (marked by yellow dashed-line borders). These distorted domains for the hot deformation sample are most likely related to (i) the significant discrepancy in atomic size that exists and (ii) slips caused by residual strain during synthesis and further strengthened by deformation.¹ Meanwhile, generous numbers of dislocations are also generated in the matrix, which also originate from lattice distortion via hot deformation. Fig. 5d shows a typical HAADF-STEM image, highlighting the atomic Z-contrast for HD3. The edge dislocations (marked with red ⊥ symbols) and distorted regions at the interfaces (marked with green rectangles) can be clearly identified from the corresponding inverse fast Fourier transform (IFFT) image (Fig. 5e). Geometric phase analysis (GPA) was carried out to illustrate strain field distribution along the xx and xy directions, as shown in Fig. 5f and g, respectively. Various and abundant multiscale defects, which can range from nanoprecipitates to a high density of nanoscale-distorted domains, to lattice distortions and dislocations, and to atomic-scale extrinsic and intrinsic point defects, effectively evoke strong phonon scattering with different frequencies and extremely reduced κ_L values over a wider temperature range. The BiSbSe_1.25Te_1.75 is composed of five layers of bright atoms with a dark contrast layer of van der Waals gaps that demarcates adjacent five-layered atomic lamellar structures, as shown in Fig. 5h (the enlarged image of area 2 in Fig. 5d).

Fig. 5 Structural characterization of the HD3 sample using STEM. (a) Medium-resolution HAADF-STEM image. (b) The corresponding elemental mappings of the nanoprecipitates in (a). (c) High-magnification HAADF-STEM image. (d) Atomic resolution high-magnification HAADF-STEM image. (e) IFFT image for (d). The inset shows an enlarged region image corresponding to area 1 in (d). (f) Stress and strain mapping along the xx (ε_xx) and (g) xy directions (ε_xy) confirmed by geometric phase analysis (GPA). (h) Enlarged image of area 2 in (d). (i) Selected area electron diffraction (SAED) pattern from (d).

Electrical properties

The electrical transport properties of all HD_x (x = 0, 1, 2, 3) samples are illustrated in Fig. 6. The dependence of S for all HD samples on temperature exhibits a typical n-type conduction behavior and initially an increase followed by a decrease, achieving the maximum value at 470 K, as shown in Fig. 6a. During the high-temperature and high-pressure sintering, additional energy was provided for atom diffusion. Cation atoms that occupied anion sites in the anti-site defects diffused back to the cation vacancies, significantly promoting the donor-like effect.¹⁷ As a result, the concentration of electrons was increased, resulting in stable n-type conduction for the BiSbSe_1.25Te_1.75 bulk sample. The degeneracy in S above 470 K is due to the emergence of bipolar conduction at high temperatures.^26,30–35 A noticeable enhancement in S (−153.6 μV K⁻¹ for HD3) with increasing hot deformation time is achieved, which presents an improvement of ∼20.3% in S compared with that of HD0 (−127.7 μV K⁻¹) at 323 K. Hot deformation causes nanoscale lamellar structures with large surface areas and the emergence of strongly preferred orientation in the HD samples while also introducing substantial grain boundaries. Both the electron and phonon transport behaviours are different in nanoscale lamellae and at grain boundaries.^36–38 Due to a charge barrier at the grain boundary having the effect of inducing interfacial resistance at the grain boundary areas, the potential barrier at the grain boundaries is generated, which leads to these regions exhibiting a larger magnitude of the Seebeck coefficient compared to the remainder of the sample.³⁹ Moreover, interfacial thermal resistance (Kapitza resistance) is also introduced at the grain boundaries, resulting in phonon scattering and thus a reduction in κ_L.^40–42 For the heterogeneous sample of nanoscale lamellar structures and grain boundary areas with a total temperature drop ΔT_t, the Seebeck coefficient can be expressed by the following equations:⁴³where S_g and S_gb represent the Seebeck coefficient of nanoscale lamellar structures and grain boundary areas, ΔT_gb is the temperature drop at the grain boundary areas, and d, κ_g and ρ_Kapitza are the grain size, the thermal conductivity of nanoscale lamellar structures and the Kapitza resistance, respectively. While increasing via increasing the Kapitza resistance will achieve a larger magnitude of the Seebeck coefficient, in TE materials, an increase in the Seebeck coefficient is coupled with a decrease in the carrier concentration. However, a decrease in the carrier concentration is not observed in the HD sample, indicating that the increased Seebeck coefficient comes from grain boundary formation rather than a reduction in the carrier concentration.

Fig. 6 Electrical transport properties of the HD_x bulk samples. (a) Temperature-dependent Seebeck coefficient S and (b) electrical conductivity σ. (c) Carrier concentration n_H and carrier mobility μ_H at 323 K. (d) Temperature-dependent power factor PF. (e) n_H dependence of S of the HD_x bulk samples. (f) Weighted carrier mobility μ_W for the HD_x bulk samples.

Fig. 6b shows the temperature-dependent σ values of all HD_x samples parallel to the pressure direction. All the samples exhibit similar negative temperature dependence, implying the typical metal-like conduction behavior. However, the negative trend gradually slows above 470 K for all samples and this can be mainly attributed to the existence of intrinsic excitation at high temperatures, which is consistent with the temperature transition point of S. Moreover, over the whole range of temperature measurement, the remarkable improvement in the σ value with increasing hot deformation time originates from the augmentation of both n_H and μ_H, as shown in Fig. 6c. Two types of slips, namely non-basal and basal slips, can be induced by the hot deformation process, respectively.⁴⁴ Abundant anion vacancies can be formed by non-basal slip, which promotes the donor-like effect and introduces into the lattice excess electrons.^21,45 Besides, basal slip also gives anion vacancies and releases electrons. These phenomena are responsible for the observed increase in n_H for the hot deformation samples. The increase of μ_H in the hot deformation samples is mainly attributed to the nanoscale lamellar structures with large surface areas and strongly preferred orientation (Fig. 4), which can produce transport expressways for electrons.⁴⁶ This is beneficial for promoting S while maintaining high σ because expressways can effectively reduce the sacrifice in μ_H. However, the gradually increased n_H with increasing hot deformation time leads to scattering between electrons, which slows the increasing tendency in μ_H gradually after repeated deformation. The corresponding PF values of all HD_x samples are shown in Fig. 6d. The temperature-dependent PF value primarily increases and then decreases, which is consistent with that of S. The significant enhancement in PF with increasing hot deformation time can be attributed to the increase in both S and σ, an ∼112.32% enhancement in PF from 478.2 (HD0) to 1015.4 (HD3) μW m⁻¹ K⁻² at 323 K. Meanwhile, a peak value of 1062.8 μW m⁻¹ K⁻² is achieved at 400 K for HD3, which is approximately an increase of 84.6% compared to that of HD0 (575.8 μW m⁻¹ K⁻²).

According to the S and n_H values at 323 K, the carrier effective mass m* and reduced Fermi level ξ_F are determined based on the single parabolic band (SPB) model, which can provide information on the electronic band structure after introducing hot deformation,^27,47 as shown in Fig. 6e. The m* value gradually increases with increasing hot deformation time and the larger m* value is indicative of the enhanced S, resulting in an improvement in PF.⁴⁸ The S values for the repeatedly deformed samples lie above the Pisarenko line, signifying that hot deformation markedly modifies the electronic band structure and enhances the density of states of the effective mass.⁴⁷ Additionally, the increase in n_H with increasing deformation time is accompanied by a gradual decrease in ξ_F, suggesting that the Fermi level gradually moves deeper into the conduction band. To further investigate the intrinsic electrical transport properties of the HD_x samples, the weighted carrier mobility μ_w was calculated based on the measured S and σ values,^27,49–51 as shown in Fig. 6f. Apparently, hot deformation can observably enhance μ_w, implying that the electrical properties can be effectively improved by optimization of the band structure.^5,47,48,51

Thermal properties

Fig. 7a presents the variation of κ_total for all HD_x samples as a function of temperature. An obvious reduction in κ_total is achieved for all samples after hot deformation, which is decreased from 0.97 W m⁻¹ K⁻¹ for HD0 to 0.82 W m⁻¹ K⁻¹ for HD1 at 323 K. The κ_total for all HD_x samples initially decreases and then begins to significantly increase at high temperatures, indicating the emergence of the bipolar effect, which is in accordance with the electrical transport results.

Fig. 7 Temperature-dependent total thermal conductivity κ_totalS (a), and electrical thermal conductivity κ_e (b) for the HD_x bulk samples. (c) Summary of the lattice thermal conductivity and bipolar thermal conductivity κ_L + κ_b, and (d) the μ_w/(κ_L + κ_b) ratio as a function of temperature for the HD_x bulk samples.

The κ_e value was calculated according to the Wiedemann–Franz law, κ_e = LσT, where L is the Lorenz number and can be determined using the single parabolic band (SPB) model.⁵² As shown in Fig. 7b, all samples after hot deformation show a high κ_e over the entire temperature interval compared to that of HD0, especially in HD2 and HD3. The dramatic increase in κ_e after hot deformation is mainly ascribed to the enhancement in σ. The κ_L + κ_b values of all samples in Fig. 7c were estimated using the formula κ_L + κ_b = κ_total − κ_e. The reduction in κ_total of the HD_x samples is mainly due to the remarkable decrease in κ_L + κ_b after hot deformation. As mentioned above, the broad nano-laminate grains with mesoscale size, local nanoprecipitates and high density of nanoscale-distorted domains (including lattice distortions and dislocations), and atomic-scale extrinsic and intrinsic point defects can be generated after hot deformation, which effectively evokes strong scattering for the low-, medium- and high-frequency phonons, respectively.^4,46,53,54 Therefore, the κ_L + κ_b value still remains at ∼0.60 W m⁻¹ K⁻¹ for the hot deformation samples at 323 K and a minimum of 0.52 W m⁻¹ K⁻¹ is achieved at 420 K for the HD3, which is a decrease of approximately 28.8% compared to that of HD0 (0.73 W m⁻¹ K⁻¹).

To further assess the optimization of TE performance, the temperature dependence of the ratio of μ_w to κ_L + κ_b was calculated. A high value of μ_w/(κ_L + κ_b) derives from not only the improved μ_w, but also from the reduced κ_L. As shown in Fig. 7d, the values of μ_w/(κ_L + κ_b) for the hot deformation samples are noticeably higher than that of HD0 over the whole temperature measurement range, suggesting that the electrical and thermal transport properties can be synergistically optimized via hot deformation.

Figure of merit ZT

The dimensionless quality factor B can be objectively used to estimate the electron transport properties of a given material, and is a good descriptor of whether the TE performance for a specified material is optimized.^55,56 The temperature-dependent B values of the HD_x bulk samples are shown in Fig. 8a. Notably, increasing B values are achieved for all samples after hot deformation over the entire temperature range. Specifically, the B value for HD3 reaches ∼1.84 at 323 K, which is almost triple that of HD0 (0.68). Moreover, the maximum B value of 2.93 is achieved at 450 K for HD3, an increase of ∼138% compared with the B value of 1.23 for the HD0 sample. The boosted B value in the temperature range of 323–550 K suggests that hot deformation is an available strategy for obtaining a high ZT value. Fig. 8b presents the temperature dependence of ZT for the HD_x samples. The ZT value increases first and then decreases for all samples, which corresponds to the trend in S. An improvement of ∼150% in ZT from 0.16 (HD0) to 0.40 (HD3) at 323 K is witnessed, benefitting from the synergistic optimization of electrical and thermal transport properties via hot deformation. The maximum ZT value reaches 0.6 at 470 K for HD3, an increase of ∼107% compared with the ZT value of 0.29 for HD0. Simultaneously, the HD3 sample also displays the highest average ZT value of 0.53 in the temperature of 323–550 K, as shown in Fig. 8c, which reveals a remarkable boosting of ∼112% in the average ZT value compared with that of HD0 (0.25). Additionally, the ZT value at low temperatures and the peak ZT value of the HD3 sample are higher than those reported in the previous literature,^57–59 as shown in Fig. 8d. Although the ZT value at low temperature in this work is still far from that of other state-of-the-art systems (Table S1†),^60–65 it undeniably establishes a foundation for further advancement of this strategy in preparing well-matched p–n junctions using TE materials with identical chemical composition.

Fig. 8 Temperature-dependent quality factor B (a) and ZT (b) of the HD_x bulk samples. (c) Average ZT of the HD_x bulk samples in the temperature range of 323–550 K. (d) Comparison of ZT values with those of other works.^57–59

Conclusion

In summary, the electrical and thermal transport properties have been synergistically optimized by hot deformation, which realizes effective improvement of TE properties for n-type BiSbSe_1.25Te_1.75. Not only does the donor-like effect elevate the n_H, but also nanoscale lamellar structures with large surface areas and strongly preferred orientation facilitate higher μ_H by providing transport expressways for electrons, giving rise to larger σ values. An enhanced PF value of 1062.8 μW m⁻¹ K⁻² was achieved at 400 K. Additionally, various and abundant multiscale distortions were introduced into the hot deformation samples, evoking strong scattering for phonons with various frequencies. Therefore, a minimum κ_L + κ_b value of 0.52 W m⁻¹ K⁻¹ was achieved at 420 K for HD3. As a result, the combination of improved PF and low κ_total values contributes to a high average ZT value of 0.53 (323 to 550 K) and a high ZT value of 0.60 (470 K) for the n-type HD3 sample. This work provides a design procedure to further boost the TE properties for n-type BiSbSe_1.25Te_1.75 by introducing nanoscale lamellae and multiscale distortions to decouple electron and phonon transport properties.

Experimental section

The experimental details consisting of the sample preparation, characterization studies, TE property measurements, and theoretical calculations are provided in the ESI.†

Author contributions

Zhen Tian: conceptualization, data curation, formal analysis, methodology and writing – original draft. Quanwei Jiang: software and writing – review & editing. Keqiang Su: software and writing – review & editing. Xiaowei Shi: software and writing – review & editing. Jianbo Li: software and writing – review & editing. Huijun Kang: conceptualization, funding acquisition, methodology and writing – review & editing. Zongning Chen: methodology and writing – review & editing. Enyu Guo: methodology and writing – review & editing. Tongmin Wang: conceptualization, funding acquisition, methodology and writing – review & editing.

Data availability

The data supporting this article have been included as part of the ESI.† The details consist of the experimental details, structural characterization, thermoelectric transport properties, the preferred orientations, the single parabolic band (SPB) model, weighted mobility, and the dimensionless quality factor B.

Conflicts of interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgements

This work was financially supported by the National Natural Science Foundation of China (Grant No. 52271025, 51927801 and U22A20174), the Science and Technology Planning Project of Liaoning Province (2023JH2/101700295), and the Innovation Foundation of Science and the Technology of Dalian (No. 2023JJ12GX021).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4qi02874k

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