Facile template-free preparation of silver-coated Cu₃SbS₄ hollow spheres with enhanced photoelectric properties

Abstract

Hollow polycrystalline structures and noble-metal-coating endow ternary sulfides with promising light absorption and electrical properties for solar energy conversion applications. We developed a facile one-pot solvothermal chemical route to synthesize hollow Cu₃SbS₄ microspheres with an opening. Additionally, four phase-pure Cu–Sb–S high-quality nanocrystals can be produced in a selectable way by the same route. The detailed evolution process and self-driven breakage mechanism have been proposed in detail. Subsequently, in situ reduction of the Ag-ion precursor was conducted and Ag clusters were uniformly deposited in the hollow spherical architecture with a nanoparticle-textured surface and interior to mitigate the oxidization of Cu₃SbS₄ nanoparticles. The created hollow metal-semiconductor nanocomposite opens up a favorable new situation. The nanocomposites exhibited a remarkably enhanced absorption of over 2 × 10⁵ cm⁻¹, covering the UV and visible regions of the electromagnetic spectrum. The electrical resistivity decreased drastically by 90% after the coating of noble metals. A reduction in the bandgap energy after the addition of Ag to Cu₃SbS₄ was confirmed from extrapolated optical results and theoretical calculation validation. The synthetic effect of hollow structures and nanocomposites results in the superior performance of Cu–Sb–S compounds, thus highlighting their potential applications in light absorbers and energy conversion technologies.

---

1. Introduction

Solar energy conversion devices have been enhanced incessantly to solve the urgent problems of non-renewable resource shortage and environmental pollution. Among these devices, thin-film solar cells are a constant pursuit because of their inherent low cost, effectiveness, lightweight, feasible flexibility and ongoing rising conversion efficiency.^1,2 Considerable effort has been devoted to discovering nontoxic and more earth-abundant alternative materials with a higher light absorption than that of commercially available CuInGaSe₂ (CIGS) and CdTe.^3,4 Many chalcogenides with a high absorption coefficient of over 10⁴ cm⁻¹ and an optimum bandgap range of 0.9–1.5 eV meet the prerequisite above, such as Cu–Zn–Sn–S–Se (CZTS or CZTSSe),^5,6 SnS,⁷ Ba–Cu–S⁸ and Cu–Sb–S,^2,9–14 and have been proposed as absorber materials for photovoltaic thin-film technologies. The record power conversion efficiency of CZTSSe has reached 12.6%,¹⁵ followed by numerous studies on other promising polynary copper-based sulfides.

Cu₃SbS₄, Cu₁₂Sb₄S₁₃, CuSbS₂ and Cu₃SbS₃, four phases of the ternary Cu–Sb–S (CAS) system, contain inexpensive and abundant elements and are expected to be of interest for a broad range of applications. The chalcostibite CuSbS₂ with a favorable bandgap of 1.55 eV has attracted much attention for research and utilization in thin-film solar cells.^9,14 In our previous work², we synthesized Ag-substituted CuSbS₂ films as the light-absorption layer material, and the efficiency of the device was significantly promoted by 2.48%. The tetrahedrite Cu₁₂Sb₄S₁₃ is being regarded as a competitive p-type thermoelectric material due to its ultralow lattice thermal conductivity of under 0.6 W mK⁻¹ originating from its complex structure with rattler atoms.^16,17 The famatinite Cu₃SbS₄ with a tetragonal structure has a bandgap of around 1.0 eV, deep valence band positions, a high hole carrier mobility (1.7 cm² V⁻¹ s⁻¹) and high absorption coefficient (10⁵ cm⁻¹).^18–20 The electron and hole relaxation dynamics of CAS compounds have been measured, and the carriers in Cu₃SbS₄ possessed a suitable lifetime of ∼5 picoseconds before trapping.²¹ Although these properties and the first-principles calculations²² have made Cu₃SbS₄ an outstanding material for potential thin-film solar cell applications, Cu₃SbS₄-involved devices have only received low efficiencies to date. Several strategies have been proposed to improve the optical absorption and transport of carriers as well as the solar energy conversion performance of devices related to CAS.

There are no materials like nanomaterials, let alone gather into specific structures. Nanocrystals composing hollow structures with interior cavities have been intentionally developed for solar cell, photocatalysis and photo-relevant applications^23–29 due to their high-surface-area, low-weight and distinct hollow features, which greatly enhances light absorption efficiency by multiple light diffraction and reflection effect within the interior cavities.²⁵ Furthermore, the shortened path of electronic or mass transport by hollow structures also leads to rapid separation of electron and hole and effectively decreases their recombination rate.^24,26,28,30 Various types of hollow metal-sulfides structures, such as spheres,^27,31,32 cubes^33,34 and tubes,^30,35 have been synthesized and obtained showing consequential enhanced photoelectric performance. Normally, binary compounds (Sb₂S₃,^36,37 NiS,³⁸ CoS₂,^27,39 CdS,²⁵ CuS⁴⁰), but not complex ternary compounds with a hollow sphere structure (such as CAS compounds), have been synthesized by tedious processes for photovoltaic applications. The detailed evolution processes and reasons for self-driven breakage are still vague, which led us to explore the formation of hollow ternary CAS structures. Meanwhile, Albuquerque has demonstrated that the preferential thin oxidized Cu²⁺ layer on the surface of the Cu₃SbS₄ nanocrystal film significantly limited the photoelectric performance of solar cells.¹¹ The surface treatment or compositing with noble metals, such as silver (Ag), may tackle this problem. The light absorption would be enhanced due to the surface plasmon resonance absorption of Ag and tunable electronic band structure of the host material.⁴¹ Meanwhile, the introduction of noble metals not only brings a high light sensitizing ability and reliable conductivity^42,43 but also the synergistic effect from different energy levels with sulfides.^44,45 The as-synthesized composites have been designed for various applications of optics,⁴⁶ solar cells,⁴⁷ photocatalysis.^30,31 Accordingly, the large surface area with a rough surface of the hollow sphere nanostructure provides ideal adsorption sites for metal–ion source nanoparticles. Can a superior all-around performance for solar energy conversion applications be obtained from the synergistic effect of hollow structures and nanocomposites?

The solvothermal method has been widely adopted to synthesize sulfide nanomaterials.^12,24 It is a simple method and can be used to efficiently produce nanomaterials at a large scale. However, a mix-phase powder was always obtained because of the narrow thermodynamic stability window of CAS in conventional processes. It remains a great challenge to synthesize phase-controlled phase-pure CAS nanocrystals and understand the formation mechanism. In this study, a facile one-pot solvothermal chemical route has been developed to synthesize monodisperse phase-pure CAS nanocrystals (Cu₃SbS₄, Cu₁₂Sb₄S₁₃, CuSbS₂ and Cu₃SbS₃) in a selectable way. The Cu₃SbS₄ nanocrystals were assembled to construct the distinctively special hollow microspheres with an opening. Afterward, in situ reduction of Ag⁺ was conducted and Ag clusters were introduced into the hollow structure to construct a metal-semiconductor nanocomposite. The hybrid exhibited a remarkably enhanced absorption of over 2 × 10⁵ cm⁻¹ covering UV and visible regions of the electromagnetic spectrum, but not a nominal composite. The band structure was systematically investigated by extrapolated optical gaps and theoretical calculations. A reduction in the bandgap energy by the addition of Ag to Cu₃SbS₄ was confirmed. The electrical resistivity decreases drastically from 0.88 Ω cm to 0.08 Ω cm by the coating of Ag. Meanwhile, the inside-outside diffusion solvothermal mechanism was proposed to elucidate the formation of the hollow microspheres with an opening in the solution. This work demonstrated that the famatinite hollow nanospheres with a noble metal coating could potentially solve problems like low efficiency, low open-circuit voltage and low fill factor values of thin-film solar panels.

2. Experimental

2.1 Synthesis of Cu–Sb–S nanocrystals

Copper(II) acetate monohydrate (99%), antimony acetate (99.99%), thiourea (99%), anhydrous acetic acid (99.8%), 2-methoxyethanol (CH₃OCH₂CH₂OH, 2-ME, 99%), diethanolamine (HO(CH₂)₂NH(CH₂)₂OH, DEA, 99%), oleylamine (OLA, 90%), silver nitrate (99%), sodium borohydride (98%) were all ordered from Aladdin Chemicals (Shanghai, China) and used without further purification.

CAS nanocrystals were synthesized by a practical one-pot solvothermal method. In a typical synthesis of Cu₃SbS₄, 9 mmol Cu(CH₃COO)₂·H₂O was loaded into 250 mL of 2-ME solution and completely dissolved with stirring. After that, 0.5 mL DEA, 72 mmol thiourea, and 3 mmol of antimony precursor were dissolved successively. The antimony precursor was prepared by dissolving antimony acetate powder in glacial acetic acid under N₂ gas protection in a glove box. The color of the solution changed from turquoise to bluish violet, then to transparent pale yellow, and finally to transparent white. Then, the prepared solution without precipitation was transferred in a Teflon-lined autoclave with a stainless steel shell. The autoclave was heated and maintained at 240 °C for 12 h before being cooled to room temperature. The solution was centrifuged at 10 000 rpm for 10 min several times by using ethanol and acetone as the dispersing solution, respectively. The upper clear solution was decanted continually, and the isolated particles were finally dried in a vacuum oven at 60 °C for 6 h. For the synthesis of the other three Cu–Sb–S phases, the preparation procedure was the same, but the ratios of elements, temperature and duration time were varied severally.

2.2 Coating of Ag on hollow spheres

The as-synthesized Cu₃SbS₄ hollow spheres were used as hosts for the incorporation of Ag nanoparticles. Firstly, Ag⁺ ions were adsorbed on the surface and interspaces by soaking spheres in the silver nitrate (0.1 mmol L⁻¹) solution for 2 h. Then, sodium borohydride was added to the solution to coat Ag onto the surface. The reduction of Ag produced a vivid yellow solution. The metal-semiconductor composite precipitates were obtained by the above-mentioned precipitation/dispersion cycles at room temperature.

2.3 Measurements and characterization

CAS nanocrystals and composites were analyzed using X-ray diffraction (XRD, D8 Advance, Bruker) with graphite monochromatic Cu Kα radiation (λ = 0.154 nm). The size of particles was calculated roughly according to the Scherrer equation,⁵⁰D = Kλ/βcos θ, where K is a dimensionless constant, λ is the XRD wavelength, β is the line-broadening at half of the maximum peak intensity and θ is the Bragg angle. The microstructural properties were studied using the Carl Zeiss GeminiSEM 300 field emission scanning electron microscope (SEM) with an energy dispersive spectrometer (EDS). Transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HRTEM) images were obtained on a JEM 2100F field-emission transmission electron microscope. The chemical composition and distribution were analyzed by X-ray photoelectron spectroscopy (XPS, Escalab 250xi, Thermo Scientific) operating at constant analyzer energy mode. The binding energies were calibrated by referencing C 1s. The intensity ratios of different binding energy peaks for several elements were set as reported.⁵¹ The optical properties of products were characterized using a UV–VIS–NIR spectrophotometer (Cary 5000, Agilent) equipped with an integrating sphere attachment using the as-prepared dry powers of nanocrystals. The optical bandgaps were estimated from (αhν)^1/2vs. hν plots by extrapolating the linear part of the function.⁵² Layer thicknesses were measured using a Veeco Dektak surface profiler. Electric properties were measured with a Hall measurement system (H5000, MMR technology) with a temperature-control system (K2000). Samples for Hall measurements were fabricated by a hot-pressing procedure under a pressure of 80 MPa at 673 K for 20 minutes.

2.4 Density functional theory calculation

Density functional theory (DFT) calculations were performed using the Vienna Ab-initio Simulation Package (VASP) code.⁵³ The generalized gradient approximation (GGA)⁵⁴ and the projector augmented wave (PAW) approach⁵⁵ were applied to obtain the optimized structures. The kinetic cut-off energy was set as 390 eV. The criterion of energy convergence for atomic positions and lattice parameters was 1 × e⁻⁵ eV. The Hellmann–Feynman forces on each atom were converged within −0.01 eV Å. The output structures and density of states (DOS) were relaxed by the hybrid functional method (HSE06). To avoid the coupling effects triggered by periodic structures, the periodic images of the slab were separated by a vacuum region of 20 Å. Brillouin zones with Γ-centered k grids of 6 × 6 × 3 and 5 × 4 × 1 were used for Cu₃SbS₄ and the one with addition of Ag clusters, respectively.

3. Results and discussions

Copper antimony sulfide nanocrystals were synthesized through a facile solvothermal route. The XRD pattern of as-synthesized powder is shown in Fig. 1(a). The diffraction peaks with high intensity corresponding only to famatinite Cu₃SbS₄ can be clearly observed (JCPDS card No. 35-0581). The sharp diffraction peaks indicate that high purity and crystallinity of the products were obtained, compared to other CAS results by solvothermal,¹² hot-injection,^{11,19,56–58} and hydrothermal methods.⁵⁹ As shown in Fig. 1(b and c), polyhedral Cu₃SbS₄ nanoparticles were generated. The inset in Fig. 1(b) exhibits the curve of the size distribution of the as-synthesized nanocrystals. Interestingly, there are obviously two peaks with high distribution at about 53 nm and 105 nm. However, the estimated calculation diameter from the full-width half maximum of the characteristic peak is about 60 nm. Why were there two sizes of nanocrystals? The HRTEM image and converted FFT image (Fig. 1(d)) of the selected region of nanocrystals in Fig. 1(c) show the perfect single-crystal structure of Cu₃SbS₄ and no amorphous edges were observed. Spacings of 0.31 nm and 0.47 nm with angles of 39.4° and 78.8° corresponding to (11̄2̄) and (101̄) planes, respectively, which is consistent with the XRD result were observed. It is worth noting that all four pure CAS phases were fabricated selectively and successfully through the same solvothermal route; XRD patterns and microstructures of other phases are shown in Fig. S1 (ESI†). The nanocrystals presented different outline features, such as irregular, trigonal-pyramidal and two-dimensional sheets morphology for Cu₃SbS₄, Cu₁₂Sb₄S₁₃ and CuSbS₂, respectively.

Fig. 1 (a) XRD patterns of the as-synthesized Cu₃SbS₄ sample along with the standard reference pattern. (b) SEM image of Cu₃SbS₄ nanocrystals along with the inset of the size distribution histogram of the sample. (c) TEM image of Cu₃SbS₄ monocrystals. (d) HRTEM image as well as FFT image as the inset.

XPS measurements were conducted to confirm the surface bonding state of CAS nanocrystals prepared using the solvothermal method. As shown in Fig. 2, elements Cu, Sb, S, except C, N and O could be observed on survey scans of all CAS samples. In a high-resolution spectrum, the chemical state of Cu 2p distinctly occupied two peaks located at 952.2 and 932.3 eV, which can be attributed to Cu (2p_1/2) and Cu (2p_3/2) binding states, respectively.^10,60,61 The peak splitting with 19.9 eV indicated Cu in Cu₃SbS₄ is in the form of Cu⁺ and no oxidation happened on the surface. As shown in Fig. S2 (ESI†), there are Cu(2P_3/2) satellite peaks located at about 962.0 and 942.0 eV as marked by arrows for the contribution from Cu²⁺ in Cu₁₂Sb₄S₁₃, besides the two peaks for the Cu⁺ oxidation state.¹⁹ The core-level spectrum of Sb 3d shows the line fit excellently to a 3d spin–orbit split doublet at 528.9 (3d_5/2) and 538.3 eV (3d_3/2) with a split orbit of 9.4 eV, which is consistent with that of Sb⁵⁺ in a thiourea-rich environment for Cu₃SbS₄.¹³ However, element Sb in Cu₁₂Sb₄S₁₃ possessed two chemical valence states of Sb⁵⁺ and Sb³⁺ according to two broad peaks.⁶² The splitting two peaks position, separations and relative intensities of S 2p are consistent with S²⁻ in all CAS compounds.¹¹

Fig. 2 XPS spectra of the as-synthesized Cu₃SbS₄ nanocrystals, including the survey spectrum along with the corresponding high-resolution XPS spectra for the Cu 2p, Sb 3d, and S 2p core levels.

A series of low magnification SEM images of the as-synthesized Cu₃SbS₄ nanocrystals are shown in Fig. 3. It can be observed that the nanocrystals are not monodispersed, but closely combined to form the nanoparticles-textured microsphere structure. The microspheres with a diameter of 0.9–1.1 μm were dispersed homogeneously. The higher magnification images in (b) and (c) show that the microspheres are particularly hollow and broken with an opening. Although the proportion of the microspheres with an opening appeared to be only about 10% from the SEM observation, the actual proportion is expected to be larger due to the limited vision. The hollow spherical structure and unique nanoparticles-textured surface would be beneficial for enhancing the light absorption, efficient transport pathways and separation of electrons and holes.^23–26 Meanwhile, the large superficial area with a rough surface provides ideal adsorption sites for metal–ion sources nanoparticles, which would introduce better optical and electric properties.

Fig. 3 (a) Low-magnification SEM image; (b) and (c) high-magnification SEM images of the as-synthesized hollow Cu₃SbS₄ spheres consisting of nanocrystals, especially the ones with an opening. (d) TEM image of the hollow sphere. (e) and (f) HRTEM images of Cu₃SbS₄ nanocrystals and their connection.

The shell of hollow microspheres with a thickness of about 200 nm at a certain stage of growth consists of big particles outside (marked in blue circle) and small nanoparticles inside (marked in red circle), as shown in Fig. 3(c). The two distinctly different sizes of crystals are close to the statistical data in Fig. 1(b). The unique morphology may give some hint on the formation of Cu₃SbS₄ hollow microspheres. TEM and HRTEM results on Cu₃SbS₄ samples are also shown in Fig. 3. The hollow microsphere can be identified by the high contrast in the middle of the volleyball-like structure. The electron penetration could be observed in the microsphere, which suggests the existence of voids or openings. To test the mechanical property and stability of the microsphere, the ultrasonic treatment for a long time was conducted before the centrifugal treatment, but the spheres were not broken in the solution, indicating that the nanocrystals were attached with each other through some connection. The HRTEM image of Fig. 3(e) shows a Cu₃SbS₄ nanoparticle connecting the other one. The lattice fringes of the two domains belong to the Cu₃SbS₄ phase. The zone axis of two planes (marked in a yellow and cyan rectangle) could be all in the [11̄1] direction, as shown in Fig. 3(f). There may be a junction or overlap (marked in ginger ellipse) between the single crystals.

As stated previously, there is no doubt that polycrystalline Cu₃SbS₄ hollow microspheres would be superior to regular particle structures as supports for noble-metal nanoparticles and thus are endowed with the improved optical and electric properties for the following reasons: (i) the distinct hollow structure for light trapping and depressed carriers recombination; (ii) the large rough surface and open interior for loading nanoparticles; (iii) surface plasmon resonance absorption and high electric conductivity of noble-metal; (iv) mitigation of oxidation of the Cu₃SbS₄ surface and guaranteed stability as absorber layers for thin-film solar cells. Fig. 4 shows the results of hollow Cu₃SbS₄ microspheres after incorporation of Ag. Element mapping analysis shows that elements Cu, Sb, S and Ag are uniformly distributed over the whole structure. We can see that the element Ag occupied more interior space with a brighter yellow. The average atomic ratio of the Ag element was about 1.98% for 1 vol% Ag within the Cu₃SbS₄-Ag heterostructure, as shown in Fig. S3 (ESI†). From XRD patterns in Fig. 4(b), the crystalline Ag phase has been synthesized and no other phases could be matched in the composite, but Cu₃SbS₄ phases. The grain size of Ag could be roughly estimated as 5 nm from XRD results (Cu₃SbS₄-Ag). There was no obvious change in the bonding state in Cu₃SbS₄ crystals, as shown in Fig. S4 (ESI†). In the high-resolution XPS spectrum of Ag 3d shown in Fig. 4(c), the two peaks located at about 367.4 and 373.4 eV can be indexed to metallic Ag 3d_5/2 and 3d_3/2, respectively. The splitting of the doublet was 6.0 eV, which also verified the synthesis of metallic silver.^29,42 There was no shift, splitting or satellite peak, and thus no presence of oxidized states.^48,63 We demonstrated that the Ag crystal was successfully decorated on the surface and interior of Cu₃SbS₄ microspheres, and the chemical transformation did not destroy the spherical morphology and composition of Cu₃SbS₄.

Fig. 4 (a) SEM image and EDS elemental mapping images of the Ag-coated hollow Cu₃SbS₄ spheres along with the relative chemical composition. (b) XRD patterns of Ag-coated Cu₃SbS₄ and the as-synthesized Ag nanoparticles along with the standard reference pattern of Ag. (c) High-resolution XPS spectrum for the Ag 2d core level.

After coating Ag, the optical absorption and electric properties for the hollow Cu₃SbS₄ microspheres were surveyed by UV–Vis absorption and Hall measurements. The absorption coefficient α is shown in Fig. 5. The estimation of α could be obtained using the following formula by minimizing the interference effects:⁶⁴

where d is the thickness of powder, T is the total transmittance, R is the total reflectance. Hollow Cu₃SbS₄ microspheres (marked in yellow) had a strong absorption coefficient of over 2 × 10⁵ cm⁻¹ covering UV and visible regions of the electromagnetic spectrum and was significantly enhanced when modified by Ag coating.⁶⁵ It is obvious that the surface plasmon resonance, dipole and quadrupolar resonance excitation from Ag were presented on Cu₃SbS₄@Ag samples, which can be beneficial for solar cells in increasing visible-light absorption and reducing the reflection of light.^66,67 The α of all samples including Ag nanoparticles, Cu₃SbS₄ powder prepared by the melting method (XRD pattern is shown in Fig. S6, ESI†), Cu₃SbS₄ hollow spheres, Cu₃SbS₄@0.5 vol% Ag, and Cu₃SbS₄@1 vol% Ag in 200–2000 nm wavelength are presented in Fig. S5(a) (ESI†). It could be concluded that Ag-coating hollow Cu₃SbS₄ microspheres with immense absorption over a wide wavelength range can be attributed to the synergistic effect of Ag-coating and the hollow structure. The optical bandgaps of these samples were determined from the optical transition by using the Tauc plot, as shown in Fig. S5(b) (ESI†). The estimated indirect bandgap of Cu₃SbS₄ nanocrystals and the Ag-coating one was close to 0.71 eV and 0.62 eV, respectively. The absorption bandgap of Cu₃SbS₄ nanocrystals and theoretical calculation was around 0.7–1.4 eV, as shown in Table S1 (ESI†). A reduction in the bandgap energy for the addition of Ag to Cu₃SbS₄ was observed, as several nanohybrids including oxide and noble metal had been reported with a similar reducing phenomenon.^67,68 Electrical properties of the as-synthesized Cu₃SbS₄ nanocrystals and Cu₃SbS₄@Ag hybrid with different Ag content were measured using the Hall measurement, as shown in Table 1.⁶⁹ The addition of Ag had an obvious impact on the electrical resistivity, carrier concentration and mobility of Cu₃SbS₄ nanocrystals. The electrical resistivity decreases drastically from 0.88 Ω cm to 0.08 Ω cm by the coating of 1 vol% Ag. All samples exhibited p-type behavior with hole concentrations from 3.1 × 10¹⁹ to 8.24 × 10¹⁹ cm⁻³. The hole mobilities were in the range of 0.9–2.4 cm² V⁻¹ S⁻¹ and present a descending trend with increasing Ag content, which would be arisen from more grain boundaries and interfaces.^64,70 The increased conductivity could be resulting from: silver intrinsically has better electric conductivity than inorganic materials, and Ag nanoparticles can fill voids and joints between crystals; the decreased bandgap results in the easier electron transition from valence band; the synergistic effect of Ag and Cu₃SbS₄ leads to the electron transport. It could demonstrate that a superior all-round performance including light absorption and electrical properties qualified hollow Cu₃SbS₄ microspheres with Ag-coating to be applied as absorbers in thin film solar cells by the synthetic effect of hollow structures and nanocomposites.

Fig. 5 UV–Vis spectra of Ag, hollow Cu₃SbS₄ spheres and Cu₃SbS₄@Ag.

Table 1 Electrical resistivity, carrier concentration and carrier mobility of the bulk samples at 300 K

Sample	Electrical resistivity (Ω cm)	Carrier concentration (cm^−3)	Carrier mobility (cm^2 V^−1 s^−1)
Cu3SbS4	8.83 × 10^−1	3.1 × 10^19	2.34
Cu3SbS4 + 0.5 vol% Ag	5.14 × 10^−1	3.35 × 10^19	1.61
Cu3SbS4 + 1 vol% Ag	8.1 × 10^−2	8.24 × 10^19	0.92

---

4. Discussion

All four stable CAS phases have been synthesized selectively by the solvothermal process. Even though there was very narrow thermodynamic stability window for CAS phases, several methods have been adopted to fabricate successfully the selective single-phase, mostly by the hot-injection method. As shown in Table S1 (ESI†), Cu₃SbS₄ nanocrystals or films with various structures and sizes were synthesized for diverse applications. Many variables were also investigated in detail to manipulate phase purity and structure of CAS in those reports and this study, and are the decisive factors in various aspects of nucleation, growth, aggregation and recrystallization, such as precursors, ratio, solvent, ligands, temperature, duration time, degree of filling, and pH as reported sporadically elsewhere.^{12,56,58,59,71,72} The systematic synthesis mechanism and analysis will be announced to elucidate the synthesis of CAS phases in other places.

In short for the Cu₃SbS₄ phase, the 2-methoxyethanol with a low polarity was used as a solvent in the solvothermal method.² Small doses of diethanolamine acted both as a surfactant and as the chemical stabilizer of precursors, especially thiourea.⁷³ When the solvothermal equipment provided temperature and pressure at a certain level, the emergence of the supersaturation of monomers would force the quick nucleation. The convection of solution and more effectively solute transport led to the growth of crystal nuclei (mostly binary Cu-S phases at the early stage, then ternary phases), which would end in the small nanocrystals and uniform in all faces, due to the insufficient source of precursors in the solvent and the low polarity of the solvent.⁶¹ In other words, solution-processed methods prefer producing nanosized nearly spherical uniform crystals, as shown mostly in Table S1 (ESI†). Actually, the high purity of Cu₃SbS₄ nanocrystals is not readily available just by selecting the element-source ratio. Based on the formation energy theory and experimental practice, the Cu₃SbS₄ phase is normally more difficult to be synthesized than Cu₁₂Sb₄S₁₃ and CuSbS₂.^58,71 It is occasionally a subproduct of the Cu₁₂Sb₄S₁₃ syntheses. The requisite for the synthesis of single-phase Cu₃SbS₄ and high purity is a higher content of the sulfide precursor than the elemental ratio in the formula.^19,56,71,74 Most of the new reports in Table S1 (ESI†) concur that element Sb would be oxidized to be Sb⁵⁺ by the created excessive sulfur source environment. Therefore, element Sb presents the V valence state in Cu₃SbS₄, also verified from the XPS measurement in this study. Interestingly, element Cu presents a reduced I valence state. A possible reaction process is proposed here. Copper(II) acetate monohydrate and antimony(III) acetate were used as Cu and Sb precursors, respectively. Thiourea is somewhat reductive, and thus Cu²⁺ would be reduced to Cu⁺ in an excessive thiourea environment, which could be confirmed by the color change of the solution during the experimental part.⁷⁵ With the rise in temperature, thiourea will decompose rapidly, with the sulfur source and generation of gases according to the following chemical reactions. Thus, excessive sulfur would oxidize Sb³⁺ to Sb⁵⁺ by a combination of a sulfidic and basic environment.^35,71,76,77

SC(NH₂)₂ + 3H₂O → (NH₄)₂CO₃ + 2H⁺ + S²⁻ (1)

SC(NH₂)₂ + 2H₂O → 2NH₃ + H₂S + CO₂ (2)

Sb³⁺ + 4S²⁻ → SbS₄³⁻ (3)

Nanocrystals with different sizes and geometries of the Cu₃SbS₄ phase have been prepared by solvothermal, hot-injection and solution-based methods, as summarized in Table S1 (ESI†), but not framing specific structures to enhance optical and electric performances so far. The hollow Cu₃SbS₄ microspheres with enhanced light absorption efficiency and improved electric properties were evolved from ions in the solvent. The evolution process is proposed in Scheme 1 based on the results from multivariate experiments. As many crystal nuclei were growing, they tended to gather into intact, nearly-spherical aggregates with an uneven surface and bonded together, as shown in Fig. S7 (ESI†). The tendency toward spherical morphology can be attributed to lowering the surface free energy of nuclei under high pressure and temperature conditions in the low polar solvent.⁷⁸ The epitaxial crystal growth would be afterward oriented on the specific crystal faces by the continuous growth model due to the rough solid–liquid interface, so the surface of microspheres became rougher.⁷⁹ Chen has synthesized the cauliflower-like spherical Cu₃SbS₄ particles with intersectional nanoflakes using the microwave irradiation method, and it could come down to a similar evolution process, yet with a two-dimensional layer growth model due to the more formative energy.¹³ Given the formation of a hollow structure, it can be deduced that there existed the Ostwald-ripening process and material transport.^23,80 Indeed, this postulation is confirmed surprisingly and directly with our results in Fig. 1 and 3. The inner crystallites (as marked in red in Fig. 3) with a small diameter remained on the inner surface of shells since the dense central crystallites dissolved to minimize the total interfacial energy. The mass transfer must occur and limit between inner solid crystals and outer particles due to the low concentration of ions in the solution after nucleation.^23,40,78,81 These particles were in close contact as the junction existed between crystals, as seen in TEM results, which is beneficial for the diffusion of ions among them. The simultaneous ripening process includes the diffusion and recrystallization, which need a sufficient experimental duration time (at least 2 hours, also as reported earlier).^40,82,83 This is the reason to believe that interactional diffusion was the primary migration route for ions. The usage of autoclave is advantageous to the acceleration of the Ostwald-ripening process, because of its ability to provide a high temperature and pressure.⁸⁴ Finally, the inner small solid crystals gave out and interior cavities were formed. The hollow microspheres with nanoparticles-textured surfaces were fabricated by the one-pot solvothermal approach. Most importantly, the hollow Cu₃SbS₄ microspheres exhibited better light absorption than that of the Cu₃SbS₄ powder prepares using the melting method as shown in Fig. S5 (ESI†). At the same time, the nanoparticles building the hollow structure possess satisfactory structural robustness for providing desorption sites to Ag, compared to that in the normally template-directed synthesis.^26,37

Scheme 1 Illustration of the hollow Cu₃SbS₄ nanocrystal spheres with Ag leading to light absorption.

The self-driven breakage of the hollow Cu₃SbS₄ microspheres resulting in an open structure is the main incentive to composite with Ag, which is beneficial for the better light absorption efficiency.^67,85 Actually, hollow Sb₂S₃ microspheres with an opening were prepared, and one more step for erosion and breaking of the closed shells by acid was needed.³⁶ Several recent reports have reported synthesizing broken hollow spheres of sulfides.^21,37,40,82 However, no detailed discussion was carried out to date. In this study, Cu₃SbS₄ spheres have fractures, but not similar to those in Cu₁₂Sb₄S₁₃, as shown in Fig. S1 (ESI†). Most strikingly, the geometrical morphologies of nanoparticles of the two phases were different, irregular polyhedron for Cu₃SbS₄ and trigonal-pyramid for Cu₁₂Sb₄S₁₃. The empirical predictor of whether the metal oxide layer that forms on a metal surface will be protective or cracking, Pilling–Bedworth ratio, is for one continuous surface, rather than the nanoparticles-textured surface.⁸³ Therefore, it is deduced that each face of the polyhedron with more faces that is close to the spheroid possesses less area than the one with several faces, which decides the area size of overlapped surfaces or junctions between nanoparticles. This is to say, trigonal-pyramid for Cu₁₂Sb₄S₁₃ results in a bigger connecting area than the irregular polyhedron for Cu₃SbS₄. The geometrical morphology of nanoparticles plays an important role in the self-driven breakage mechanism. Furthermore, it was reported that stress buildup during growth would cause the removal of shells from the high end of the outer diameter distribution due to the compressive stress concentration.⁸³

While the synthesis of hollow Cu₃SbS₄ microspheres was well established and light absorption from trapping was beneficial for its application in solar cells, it is still extremely challenging to come up with a solution for the oxidization of the surface Cu₃SbS₄ film.^11,33,40In situ reduction of Ag-ion precursor was conducted and Ag clusters were uniformly deposited on the providential hollow spherical architecture, prior to oxidation of the surface of nanoparticles according to XPS and TEM results. The building of metal-semiconductor nanosystems also brought an intense absorption in both the UV and visible regions and substantial improvement of electric properties. Firstly, the enhancing light absorption should be due to the surface plasmon resonances of Ag clusters with different geometrical morphologies according to UV–Vis results. The specific resonance effect, such as in-plane dipole and out-of-plane quadrupole characters, played an antireflection role on the surface.^66,67 The photogenerated carriers or energy excited and trapped from the resonances of dispersed Ag clusters migrated irreversibly to the direct electronic coupling Cu₃SbS₄ shell through the coherent interfaces, which was in the region of the enhanced accumulation local field since the bottom of the conduction band of Cu₃SbS₄ is close to the overall Fermi level of Ag.^36,85–87 In addition, silver intrinsically has better electric conductivity than inorganic materials. The Ag nanoparticles can fill the voids and joints between crystals, and thus increase the carrier transport. Hence, carrier concentration and electrical conductivity presented a growing tendency as the amount of added Ag increased. There is also an important aspect involved here: the ameliorative band structure by the redistribution of the space charges. It has been demonstrated that the additional states could be generated by the coating of the noble metal, which acted as the scavenger of carriers at the semiconductor interface.^49,68

Many groups have theoretically investigated the energy band structure of CAS compounds.^{56,69,72,88,89} To check the above conjecture, we also calculated for electronic structures of Cu₃SbS₄ and the simplified model for noble metal decorated Cu₃SbS₄ (e.g., Ag) by HSE06. Fig. 6 shows the atomic structures of Cu₃SbS₄ and the one with the addition of the Ag cluster. The lattice parameters of famatinite were determined as a = b = 5.3850 Å and c = 10.7540 Å. The effect of the Ag cluster on the Cu₃SbS₄ (112) face was analyzed. The distance between the one-atom-Ag cluster and S was optimized to get a stable state based on previous reports.^68,90 They bonded through van der Waals bonds with a 2.43 Å length.⁶⁸ The total density of states (TDOS) for Cu₃SbS₄ and partial density of states (PDOS) for the heterostructure computed by HSE06 are shown. There is a disagreement of the bandgap of bulk Cu₃SbS₄ between the computed (2.14 eV) and measured ones (0.71 eV), similar to Ramasamy's report.⁵⁶ An obvious reduction in bandgap energy for the addition of Ag clusters to Cu₃SbS₄ was observed due to the reconstruction of the space charges distribution,⁶⁷ as shown in Fig. 6(b). It was demonstrated that coating of Ag clusters would generate additional states.⁶⁸ The reduced bandgap led to lower electrical resistivity and enhanced light absorption.

Fig. 6 (a) The total density of states for the Cu₃SbS₄ phase and (b) the partial density of states for Cu₃SbS₄/Ag. Insets show the corresponding optimized crystal structures. The green dashed line stands for the Fermi level position. The red dotted line in (b) frames out the Ag cluster and adjacent bilayer of Cu₃SbS₄.

5. Conclusion

To improve all-around performance in thin-film solar cells, we have demonstrated hollow Cu₃SbS₄@Ag spheres composed of mutually bonded nanoparticles. A remarkably enhanced absorption over 2 × 10⁵ cm⁻¹ covering UV and visible regions and electrical resistivity of 0.08 Ω cm could be obtained from the synergistic effect of hollow structural features and surface plasmon resonance of Ag clusters. The bandgap energy was calculated theoretically for the Cu₃SbS₄/Ag system and a slight reduction of bandgap was observed. The detailed evolution process and self-driven breakage mechanism of hollow microspheres, and reinforcing way of noble metal were proposed in detail. Additionally, a facile one-pot solvothermal method was demonstrated to selectively synthesize four CAS nanocrystals of high purity and crystallinity and not only Cu₃SbS₄. The element Cu in Cu₃SbS₄ was in the form of Cu⁺ and not oxidized immediately after the synthesis, which allowed us to decorate the surface. The preparation mechanism and the reaction process of CAS phases was discussed in a chemistry way. It is believed that the detailed mechanism could guide the synthesis of most ternary sulfides and control their structural performance, as well as the optical, electrical and catalytic properties.

Author contributions

Xiang Meng: conceptualization, data curation, writing – original draft; Xihao Chen: software, formal analysis, investigation; Jiang Cheng: methodology, visualization, resources; Fuqiang Zhai: formal analysis, resources; Wen Li: supervision, validation, writing – review & editing; Rui Hu: supervision, validation, writing – review & editing; Lu Li: supervision, validation, visualization; all authors have agreed to their individual contributions in this manuscript.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work was supported by the program of Public Service Platform for the Industrialization of Technological Innovation Achievements in the field of Robot and Intelligent Manufacturing in Chongqing (No. 2019-00900-1-1), the Science and Technology Research Program of Chongqing Municipal Education Commission Grant (No. KJQN201901324, KJQN201901348 and KJCX2020048) and Chongqing University of Arts and Sciences (P2018CL09).

Notes and references

1. T. D. Lee and A. U. Ebong, Renewable Sustainable Energy Rev., 2017, 70, 1286–1297.
2. L. Fu, J. Yu, J. Wang, F. Xie, S. Yao, Y. Zhang, J. Cheng and L. Li, Chem. Eng. J., 2020, 400, 125906.
3. R. Kamada, T. Yagioka, S. Adachi, A. Handa, K. F. Tai, T. Kato and H. Sugimoto, in 2016 IEEE 43rd Photovoltaic Specialists Conference (PVSC), IEEE, Portland, OR, USA, 2016, pp. 1287–1291.
4. W. K. Metzger, S. Grover, D. Lu, E. Colegrove, J. Moseley, C. L. Perkins, X. Li, R. Mallick, W. Zhang, R. Malik, J. Kephart, C.-S. Jiang, D. Kuciauskas, D. S. Albin, M. M. Al-Jassim, G. Xiong and M. Gloeckler, Nat. Energy, 2019, 4, 837–845.
5. F. Jiang, S. Ikeda, T. Harada and M. Matsumura, Adv. Energy Mater., 2014, 4, 1301381.
6. K. Woo, Y. Kim and J. Moon, Energy Env. Sci., 2012, 5, 5340–5345.
7. V. Steinmann, R. Chakraborty, P. H. Rekemeyer, K. Hartman, R. E. Brandt, A. Polizzotti, C. Yang, T. Moriarty, S. Gradečak, R. G. Gordon and T. Buonassisi, ACS Appl. Mater. Interfaces, 2016, 8, 22664–22670.
8. Y. Han, S. Siol, Q. Zhang and A. Zakutayev, Chem. Mater., 2017, 29, 8239–8248.
9. S. C. Riha, A. A. Koegel, J. D. Emery, M. J. Pellin and A. B. F. Martinson, ACS Appl. Mater. Interfaces, 2017, 9, 4667–4673.
10. C. Behera, R. Samal, C. S. Rout, R. S. Dhaka, G. Sahoo and S. L. Samal, Inorg. Chem., 2019, 58, 15291–15302.
11. G. H. Albuquerque, K.-J. Kim, J. I. Lopez, A. Devaraj, S. Manandhar, Y.-S. Liu, J. Guo, C.-H. Chang and G. S. Herman, J. Mater. Chem. A, 2018, 6, 8682–8692.
12. L. Shi, C. Wu, J. Li and J. Ding, J. Alloys Compd., 2017, 694, 132–135.
13. G. Chen, W. Wang, J. Zhao, W. Yang, S. Chen, Z. Huang, R. Jian and H. Ruan, J. Alloys Compd., 2016, 679, 218–224.
14. B. Yang, L. Wang, J. Han, Y. Zhou, H. Song, S. Chen, J. Zhong, L. Lv, D. Niu and J. Tang, Chem. Mater., 2014, 26, 3135–3143.
15. W. Wang, M. T. Winkler, O. Gunawan, T. Gokmen, T. K. Todorov, Y. Zhu and D. B. Mitzi, Adv. Energy Mater., 2014, 4, 1301465.
16. F.-H. Sun, J. Dong, H. Tang, P.-P. Shang, H.-L. Zhuang, H. Hu, C.-F. Wu, Y. Pan and J.-F. Li, Nano Energy, 2019, 57, 835–841.
17. X. Lu, D. T. Morelli, Y. Wang, W. Lai, Y. Xia and V. Ozolins, Chem. Mater., 2016, 28, 1781–1786.
18. Q. Zeng, Y. Di, C. Huang, K. Sun, Y. Zhao, H. Xie, D. Niu, L. Jiang, X. Hao, Y. Lai and F. Liu, J. Mater. Chem. C, 2018, 6, 7989–7993.
19. J. van Embden, K. Latham, N. W. Duffy and Y. Tachibana, J. Am. Chem. Soc., 2013, 135, 11562–11571.
20. N. D. Franzer, N. R. Paudel, C. Xiao and Y. Yan, in 2014 IEEE 40th Photovoltaic Specialist Conference (PVSC), IEEE, Denver, CO, USA, 2014, pp. 2326–2328.
21. F. Zhang, K. Chen, X. Jiang, Y. Wang, Y. Ge, L. Wu, S. Xu, Q. Bao and H. Zhang, J. Mater. Chem. C, 2018, 6, 8977–8983.
22. C. T. Crespo, J. Phys. Chem. C, 2016, 120, 7959–7965.
23. L. Zhou, Z. Zhuang, H. Zhao, M. Lin, D. Zhao and L. Mai, Adv. Mater., 2017, 29, 1602914.
24. J. Wang, Y. Cui and D. Wang, Adv. Mater., 2019, 31, 1801993.
25. M. Luo, Y. Liu, J. Hu, H. Liu and J. Li, ACS Appl. Mater. Interfaces, 2012, 4, 1813–1821.
26. L. Zhang, J. Colloid Interface Sci., 2018, 514, 396–406.
27. H. Sun, X. Zhang, F. Yan, H. Yuan, C. Zhu and Y. Chen, Appl. Surf. Sci., 2021, 554, 149238.
28. J. Wang, Y. Quan, G. Wang, D. Wang, J. Xiao, S. Gao, H. Xu, S. Liu and L. Cui, CrystEngComm, 2021, 23, 7385–7396.
29. W. Zuo, L. Liang, F. Ye and S. Zhao, Chemosphere, 2021, 283, 131167.
30. G. Zhang, D. Chen, N. Li, Q. Xu, H. Li, J. He and J. Lu, Angew. Chem., 2020, 132, 8332–8338.
31. W.-N. Wang, C.-Y. Zhang, B.-J. Chen, Z.-Y. Chu, Z.-B. Zha and H.-S. Qian, Appl. Surf. Sci., 2020, 515, 146130.
32. X. Feng, Q. Jiao, H. Cui, M. Yin, Q. Li, Y. Zhao, H. Li, W. Zhou and C. Feng, ACS Appl. Mater. Interfaces, 2018, 10, 29521–29531.
33. L. Su, Y. Hu, Z. Ma, L. Miao, J. Zhou, Y. Ning, Z. Chang, B. Wu, M. Cao, R. Xia and J. Qian, Sol. Energy Mater. Sol. Cells, 2020, 210, 110484.
34. T. Duan, W. Lou, X. Wang and Q. Xue, Colloids Surf., A, 2007, 310, 86–93.
35. Y.-H. Zhang, L. Guo, L. He, K. Liu, C. Chen, Q. Zhang and Z. Wu, Nanotechnology, 2007, 18, 485609.
36. X. B. Cao, L. Gu, L. J. Zhuge, W. J. Gao, W. C. Wang and S. F. Wu, Adv. Funct. Mater., 2006, 16, 896–902.
37. S. Ding, X. Liu, Y. Shi, Y. Liu, T. Zhou, Z. Guo and J. Hu, ACS Appl. Mater. Interfaces, 2018, 10, 17911–17922.
38. Y. Wang, Q. Zhu, L. Tao and X. Su, J. Mater. Chem., 2011, 21, 9248.
39. J. Yang, Y. Niu, J. Huang, L. Liu and X. Qian, Electrochim. Acta, 2020, 330, 135333.
40. B. Zhao, G. Shao, B. Fan, W. Zhao, Y. Xie and R. Zhang, J. Mater. Chem. A, 2015, 3, 10345–10352.
41. Y. Sun, Nanoscale, 2010, 2, 1626.
42. S. Saeed, H. Siddique, R. Dai, J. Wang, A. Ali, C. Gao, X. Tao, Z. Wang, Z. Ding and Z. Zhang, J. Phys. Chem. C, 2021, 125, 7542–7551.
43. W. Ren, Y. Yang, J. Yang, H. Duan, G. Zhao and Y. Liu, Chem. Eng. J., 2021, 415, 129052.
44. Z. Li, W. Huang, J. Liu, K. Lv and Q. Li, ACS Catal., 2021, 11, 8510–8520.
45. B. Vedhanarayanan, C.-C. Chen and T.-W. Lin, J. Power Sources, 2020, 477, 229033.
46. M. W. Knight, H. Sobhani, P. Nordlander and N. J. Halas, Science, 2011, 332, 702–704.
47. H. Li, Y. Hu, H. Wang, Q. Tao, Y. Zhu and Y. Yang, Sol. RRL, 2021, 5, 2000524.
48. N. Zhang, M. Li, C. F. Tan, C. K. Nuo Peh, T. C. Sum and G. W. Ho, J. Mater. Chem. A, 2017, 5, 21570–21578.
49. W. Gao, Q. Liu, S. Zhang, Y. Yang, X. Zhang, H. Zhao, W. Qin, W. Zhou, X. Wang, H. Liu and Y. Sang, Nano Energy, 2020, 71, 104624.
50. P. Scherrer, Nachr. Ges. Wiss. Gottingn, 1918, 2, 89–93.
51. J. F. Moulder and J. Chastain, ed., Handbook of X-ray photoelectron spectroscopy: a reference book of standard spectra for identification and interpretation of XPS data, PerkinElmer Corporation, Eden Prairie, Minn, Update., 1992.
52. E. A. Davis and N. F. Mott, Philos. Mag., 1970, 22, 0903–0922.
53. G. Kresse and J. Furthmüller, Phys. Rev. B: Condens. Matter Mater. Phys., 1996, 54, 11169–11186.
54. J. P. Perdew and Y. Wang, Phys. Rev. B: Condens. Matter Mater. Phys., 1992, 46, 12947–12954.
55. P. E. Blöchl, Phys. Rev. B: Condens. Matter Mater. Phys., 1994, 50, 17953–17979.
56. K. Ramasamy, H. Sims, W. H. Butler and A. Gupta, Chem. Mater., 2014, 26, 2891–2899.
57. Q. Liang, K. Huang, X. Ren, W. Zhang, R. Xie and S. Feng, CrystEngComm, 2016, 18, 3703–3710.
58. K. Chen, J. Zhou, W. Chen, Q. Chen, P. Zhou and Y. Liu, Nanoscale, 2016, 8, 5146–5152.
59. C. An, Y. Jin, K. Tang and Y. Qian, J. Mater. Chem., 2003, 13, 301–303.
60. V. K. Mariappan, K. Krishnamoorthy, P. Pazhamalai, S. Sahoo, S. S. Nardekar and S.-J. Kim, Nano Energy, 2019, 57, 307–316.
61. S. Yang, K. Liu, W. Han, L. Li, F. Wang, X. Zhou, H. Li and T. Zhai, Adv. Funct. Mater., 2020, 30, 1908382.
62. L. Wang, B. Yang, Z. Xia, M. Leng, Y. Zhou, D.-J. Xue, J. Zhong, L. Gao, H. Song and J. Tang, Sol. Energy Mater. Sol. Cells, 2016, 144, 33–39.
63. Q. Dong, Z. Jiao, H. Yu, J. Ye and Y. Bi, CrystEngComm, 2014, 16, 8317.
64. P. Sinsermsuksakul, J. Heo, W. Noh, A. S. Hock and R. G. Gordon, Adv. Energy Mater., 2011, 1, 1116–1125.
65. P. A. Fernandes, A. Shongalova, A. F. da Cunha, J. P. Teixeira, J. P. Leitão, J. M. V. Cunha, S. Bose, P. M. P. Salomé and M. R. Correia, J. Alloys Compd., 2019, 797, 1359–1366.
66. E. Thouti, N. Chander, V. Dutta and V. K. Komarala, J. Opt., 2013, 15, 035005.
67. Ö. A. Yıldırım, H. E. Unalan and C. Durucan, J. Am. Ceram. Soc., 2013, 96, 766–773.
68. C. A. Huerta-Aguilar, V. P. Barba, P. Thangarasu and R. T. Koodali, Chemosphere, 2018, 213, 481–497.
69. L. Yu, R. S. Kokenyesi, D. A. Keszler and A. Zunger, Adv. Energy Mater., 2013, 3, 43–48.
70. Q. Zhang, X. Ai, L. Wang, Y. Chang, W. Luo, W. Jiang and L. Chen, Adv. Funct. Mater., 2015, 25, 966–976.
71. F. Baum, T. Pretto, A. G. Brolo and M. J. L. Santos, Cryst. Growth Des., 2018, 18, 6521–6527.
72. U. V. Ghorpade, M. P. Suryawanshi, S. W. Shin, X. Wang, E. Jo, H. Bae, K. Park, J.-S. Ha, S. S. Kolekar and J. H. Kim, J. Mater. Chem. A, 2018, 6, 19798–19809.
73. J. Cheng, Z. Dai, B. Chen, R. Ji, X. Yang, R. Hu, J. Zhu and L. Li, Nanoscale Res. Lett., 2016, 11, 550.
74. J. van Embden and Y. Tachibana, J. Mater. Chem., 2012, 22, 11466.
75. S. Ariponnammal, S. Chandrasekaran and C. Sanjeeviraja, Cryst. Res. Technol., 2012, 47, 145–150.
76. S. Li, Z. Wu, W. Li, Y. Liu, R. Zhuo, D. Yan, W. Jun and P. Yan, CrystEngComm, 2013, 15, 1571.
77. X. Yu, J. Yu, B. Cheng and B. Huang, Chem. – Eur. J., 2009, 15, 6731–6739.
78. Y. Zhu, T. Liu, W. Ma, D. Zhang, N. Zhang, L. Chen, X. Liu and G. Qiu, Chem. Lett., 2015, 44, 1321–1323.
79. M. Aliofkhazraei, ed., Handbook of Nanoparticles, Springer International Publishing, Cham, 2016.
80. W. Ostwald, Z. Phys. Chem., 1900, 34U, 495–503.
81. H. G. Yang and H. C. Zeng, J. Phys. Chem. B, 2004, 108, 3492–3495.
82. L. Shen, L. Lu, H. Wu, X.-Y. Yu, X. Zhang and X. W. Lou, Nat. Commun., 2015, 6, 1–8.
83. Y. Yin, C. K. Erdonmez, A. Cabot, S. Hughes and A. P. Alivisatos, Adv. Funct. Mater., 2006, 16, 1389–1399.
84. X. Cao, L. Gu, L. Zhuge, W. Qian, C. Zhao, X. Lan, W. Sheng and D. Yao, Colloids Surf., A, 2007, 297, 183–190.
85. T. A. Estrada-Mendoza, D. Willett and G. Chumanov, J. Phys. Chem. C, 2020, 124, 27024–27031.
86. J. Li, X. Han, J. Li, Y. Zhao and C. Fan, Phys. Status Solidi B, 2017, 254, 1600608.
87. Y. Sun, Y. Wang, J. Y. C. Chen, K. Fujisawa, C. F. Holder, J. T. Miller, V. H. Crespi, M. Terrones and R. E. Schaak, Nat. Chem., 2020, 12, 284–293.
88. T. Shi, W.-J. Yin, M. Al-Jassim and Y. Yan, Appl. Phys. Lett., 2013, 103, 152105.
89. K. Chen, C. Di Paola, B. Du, R. Zhang, S. Laricchia, N. Bonini, C. Weber, I. Abrahams, H. Yan and M. Reece, J. Mater. Chem. C, 2018, 6, 8546–8552.
90. V. Bonačic-Koutecky, V. Veyret and R. Mitric, J. Chem. Phys., 2001, 115, 12.

---

Footnote

† Electronic supplementary information (ESI) available: XRD patterns and SEM images for Cu₁₂Sb₄S₁₃, Cu₃SbS₃ and CuSbS₂ phase by the solvothermal method; XPS spectra of as-synthesized Cu₁₂Sb₄S₁₃ nanocrystals; Calculation method of Ag content from XRD results; EDS result of the as-synthesized Ag-coated Cu₃SbS₄ nanocrystals; XPS spectra of the as-synthesized Ag-coated Cu₃SbS₄ nanocrystal; UV–Vis–NIR spectra of Ag nanoparticles, Cu₃SbS₄ powder by melting method, Cu₃SbS₄ nanocrystals hollow spheres, 0.5 vol% Ag-coated Cu₃SbS₄ and 1 vol% Ag-coated Cu₃SbS₄, the corresponding Tauc plot of (αhν)^1/2versus (hν) of the as-synthesized Cu₃SbS₄ nanocrystals hollow spheres and Ag-coated ones; XRD pattern of the as-synthesized bulk Cu₃SbS₄ by melting method; SEM images of the smaller spherical aggregation of crystals by only 5 min duration time at high temperature; The representative preparation method, structural properties and optical bandgap (indirect) of Cu₃SbS₄. See DOI: 10.1039/d1tc04098g

---