Nanoarray heterojunction and its efficient solar cells without negative impact of photogenerated electric field

Abstract

Efficient, stable and low-cost solar cells are being desired for the photovoltaic conversion of solar energy into electricity for sustainable energy production. Nanorod/nanowire arrays of narrow-bandgap semiconductors are the promising light-harvesters for photovoltaics because of their excellent optoelectrical properties. Here, the array of preferentially oriented antimony trisulfide (Sb₂S₃) single-crystalline nanorods is grown on polycrystalline titania (TiO₂) film by a tiny-seed-assisted solution-processing strategy, offering an Sb₂S₃/TiO₂ nanoarray heterojunction system on a large scale. It is demonstrated that the Sb₂S₃ nanorod growth follows a tiny-seed-governed orientation-competing-epitaxial nucleation/growth mechanism. Using a conjugated polymer hole transporting layer on the heterojunction, we achieve a power conversion efficiency of 5.70% in the stable hybrid solar cell with a preferred p-type/intrinsic/n-type architecture featuring effectively straightforward charge transport channels and no negative impact of photogenerated electric field on device performance. An architecture-dependent charge distribution model is proposed to understand the unique photovoltaic behavior.

Introduction

Photovoltaic technology is one of the promising ways to solve the problems of energy demand, and efficient, stable, low-cost, and non-toxic solar cells are being desired for the conversion of solar energy into electricity. With the excellent properties, such as photon absorption through coupling effect, straightforward pathways for charge carrier transport, and defined optoelectrical property by tailored structure, the nanorods/nanowires of narrow-bandgap semiconductors have attracted intensive attention to their photovoltaic applications^1,2,3,4. Many narrow-bandgap nanorod/nanowire arrays have been used as light-harvesting materials in solar cells, including silicon^1,2 and III−V compounds (e.g., GaAs, GaN, InP)^5,6,7. Amongst the III−V compounds, axial and radial p–n^5,7,8,9 or p-intrinsic-n (p–i–n)^5,6,7,10,11 junction nanorods/nanowires are generally designed for those nanorod/nanowire arrays in order to obtain high photovoltaic performance. An axial p–n or p–i–n junction nanorod/nanowire is normally created by changing doping states along the axis direction of the nanorod/nanowire, while a radial p–n or p–i–n junction nanorod/nanowire is often obtained by forming a core-shell structure along the radial direction of the nanorod/nanowire due to the different doping states for example. However, the preparation of those p–n and p–i–n junction nanorod/nanowire arrays usually involves sophisticated chemical vapor deposition and molecular beam epitaxy technologies.

Growing the nanorod/nanowire array of intrinsic narrow-bandgap semiconductor (A) on the film of n-type semiconductor (B) renders a material system, which is here referred to as a nanoarray heterojunction (NHJ) since individual A/B heterojunctions are actually resulted upon the contact between each of A nanorods and the beneath B film and the A/B-NHJ property correlates tightly to the discontinuous and nanosized A/B heterojunctions (i.e., A/B interfaces). On one hand, the NHJ is quite different from the p–n and p–i–n junction nanorod/nanowire arrays in that the properties of the latter two arrays are mainly determined by the p–n and p–i–n junction nanorod/nanowires. On the other hand, the A/B-NHJ is also essentially different from the planner heterojunction (PHJ) formed by sequentially depositing A compact film onto B film (A/B-PHJ). The A/B-NHJ has many discontinuous and nanosized A/B interfaces for interfacial charge generation and transportation due to the interspaces in between nanorods, but the A/B-PHJ has only one two-dimensional and micro-/macro-sized A/B interface for interfacial charge behaviors. Moreover, when combining with a third material C, the A/B-PHJ only offers an A surface to form another new PHJ with a two-dimensional interface of A and C components and results in a ternary C/A/B-PHJ derivative system, while the A/B-NHJ has many side-faces and top-faces of A nanorods/nanowires to form a bulk heterojunction (BHJ) with three-dimensionally distributed A/C interfaces and yields an integrated C/A/B derivative of two-dimensional PHJ and three-dimensional BHJ structures. Therefore, the distinct structural characteristics of the NHJ and PHJ will inevitably lead to the different optoelectrical and photophysical properties of themselves and their derivatives. In short, the NHJ represents a kind of heterojunction formation and is quite different from the PHJ and the p–n or p–i–n junction nanorod/nanowire array. The NHJs formed by the nanorods/nanowires of narrow-bandgap semiconductors on wide-bandgap semiconductors (e.g., TiO₂, ZnO) films should be of great interest for photovoltaic application in that they combine together the good absorption and straightforward charge transport channels, but their preparation still remains a challenge.

Antimony trisulfide (Sb₂S₃) is a promising photon-harvesting candidate for solar cells because of its narrow and direct bandgap (~1.4−1.8 eV) matching well the solar spectrum, high absorption coefficient (~10⁵cm⁻¹), excellent air/moisture stability, non-toxicity, earth-abundant elemental composition, and easy preparation¹². The Sb₂S₃-based single-junction solar cells are theoretically expected to achieve the efficiency of η = 27–33% according to Shockley–Queisser limit¹³. Sb₂S₃ nanoparticle-sensitized solar cells have exhibited a record efficiency of η = 7.5%¹⁴. Moreover, the Sb₂S₃ nanoparticle^15,16,17 or bulk^18,19 films as a photon-harvesting layer in planar heterojunction solar cells have achieved the efficiencies of 5.69−6.78% and 3.50−5.12%, respectively. Sb₂S₃ crystals comprise the (Sb₄S₆)_n ribbons that grow along [001] direction via strong covalent Sb–S bonds and are packed together alone [100] and [010] directions by van der Waals forces (Inset to Fig. 1), and therefore Sb₂S₃ often features a one-dimensional crystalline structure and highly anisotropic properties along different axes. While the one-dimensional nanostructures of Sb₂S₃ nanorods/nanowires have been prepared by different methods²⁰, no reports on the growth of Sb₂S₃ nanorod/nanowire arrays are available up to now.

Fig. 1: Tiny-seed assisted repetition of spin-coating and annealing (RSCA) method, structures of materials, and device.

The schematic illustration shows the preparation of the Sb₂S₃/TiO₂-nanoarray heterojunction (Sb₂S₃/TiO₂-NHJ) by a tiny-seed-assisted RSCA method on fluorine-doped tin oxide (FTO) substrate and the structure of p–i–n solar cell, where p means p-type PTB7-th, i means intrinsic Sb₂S₃ and n means n-type TiO₂. The insets show the photographs of seed growth (bottle 1, colorless) and nanorod growth (bottle 2, yellow) solutions, the molecular structures of PTB7-th, the crystal structure of orthorhombic Sb₂S₃, and the cross-sectional scanning electron microscopy (SEM) characterization of a practical p–i–n cell. Normally, the seed growth solution concentration of C_sg = 0.1 M and the nanorod growth solution concentration of C_ng = 0.8 M were, respectively, used for growing the Sb₂S₃ seeds prior to the RSCA cycles and the Sb₂S₃ nanorods in the RSCA cycles, unless otherwise specified. During the RSCA method, each of the precursor film formations by spin-coating followed by once annealing is referred to as one RSCA cycle for the nanorod growth, and the nanorod length (i.e., nanorod array thickness) is mainly controlled by the RSCA cycle number (n).

In this work, we develop a tiny-seed-assisted repetition of spin-coating and annealing (RSCA) strategy for preparing preferentially [211]-oriented and vertically aligned Sb₂S₃ single-crystalline nanorods on polycrystalline TiO₂ nanoparticle film, yielding an Sb₂S₃/TiO₂-NHJ, for which an epitaxial nucleation/growth mechanism gets revealed for the Sb₂S₃ nanorod growth. After depositing a conjugated polymer layer onto the Sb₂S₃/TiO₂-NHJ, a stable hybrid solar cell with p–i–n architecture is obtained and exhibits a considerable efficiency of 5.70%. It is found that, furthermore, there is no negative impact of photogenerated electric field (E_ph) on photovoltaic performance in such p–i–n solar cells, and an electric field distribution model is suggested for understanding the distinguishing photovoltaic feature.

Noticeably, we previously developed a seed-assisted RSCA method for preparing the monolayer of [221]-oriented bulk Sb₂S₃ single-crystalline cuboids on polycrystalline TiO₂ nanoparticle film¹⁹, where [211]/[221]-oriented Sb₂S₃ crystalline seeds laterally of 60−400 nm in size were deposited on the TiO₂ film prior to the RSCA cycles for Sb₂S₃ crystal growth, and an orientation-competing-epitaxial (OCE) nucleation/growth mechanism was proposed for the [221]-oriented Sb₂S₃ cuboid growth. In this work, we use the tiny [211]/[221]-oriented Sb₂S₃ crystalline seeds of laterally <10 nm in size to perform the seed-assisted RSCA process and obtain the [211]-oriented Sb₂S₃ single-crystalline nanorod array on the polycrystalline TiO₂ nanoparticle film. Our results demonstrate that the Sb₂S₃ nanorod growth follows a tiny-seed-governed OCE nucleation/growth mechanism, where the tiny size of Sb₂S₃ seeds is exclusively important to govern the orientation and morphology of resulting Sb₂S₃ nanorods. Undoubtedly, in regard to the heterojunction preparation, the insights in this work make significant advances in expanding the knowledge on the seed-assisted RSCA method and the OCE nucleation/growth mechanism.

Results

Preparation strategy

For simplicity, the Sb₂S₃ precursor solutions for Sb₂S₃ seed growth (SG) and nanorod growth (NG) are referred to as SG and NG solutions, respectively. The SG and NG solutions have the same composition and their concentrations (C_sg and C_ng) are defined by the molecular concentration of SbCl₃ therein. For the growth of the Sb₂S₃ nanorod array, C_ng must be much higher than C_sg, as indicated by solution color (Bottles 1 and 2, inset to Fig. 1). As shown in Fig. 1, an Sb₂S₃ precursor film is spin-coated from an NG solution onto the TiO₂ nanoparticle film decorated with oriented Sb₂S₃ crystalline seeds and then annealed in N₂ atmosphere at 300 ^°C, which is referred to as one RSCA cycle number (n); the Sb₂S₃/TiO₂-NHJ is produced after a certain RSCA cycle numbers (i.e., n ≥ 1). By coating a poly[4,8-bis(5-(2-ethylhexyl) thiophen-2-yl)benzo[1,2-b;4,5-b’]dithiophene-2,6-diyl-alt-(4-(2-ethylhexyl)-3-fluorothieno[3,4-b]thiophene-)-2-carboxylate-2-6-diyl)] (PTB7-th, inset to Fig. 1) layer, we obtain a p–i–n hybrid solar cell with the spatially separated intrinsic Sb₂S₃ nanorods sandwiched between p-type PTB7-th hole transporting layer (HTL) and n-type TiO₂ electron transporting layer (ETL), and scanning electron microscopy (SEM) image (Inset to Fig. 1) clearly confirms the p–i–n architecture.

Characterization of nanorod array

SEM images (Fig. 2a, b) show that a large-area Sb₂S₃ layer is grown on the TiO₂ film and consists of the Sb₂S₃ nanorods that are of a right square prism shape, 50−150 nm in width, ca. 320 nm in length, and a number density (i.e., the number of the different nanorods per unit area) of ca. 36 µm⁻² on the TiO₂ nanoparticle (ca. 10−20 nm in diameter) film of ca. 110 nm in thickness. Clearly, the Sb₂S₃ nanorods are grown along the directions out of substrate planes (i.e., out-of-plane growth) with the growth direction tilting some angles from the normal of TiO₂ film, resulting in an Sb₂S₃ nanorod array on the TiO₂ film. Fig. 2c shows the transmission electron microscopy (TEM) image of a single Sb₂S₃ nanorod of ca. 50 nm in width. Both selected area electron diffraction pattern and high-resolution TEM (HRTEM) image (Insets to Fig. 2c) clearly show that the nanorod is single-crystalline in nature and grows along [001] direction, where the HRTEM image exhibits the lattice fringes of 0.35 nm and 0.38 nm in spacing corresponding, respectively, to the $\left( {111} \right)/\left( {\bar 1\bar 11} \right)$) and (001) planes of orthorhombic Sb₂S₃ under $\left[ {1\bar 10} \right]$ viewing direction. X-ray diffraction (XRD) results (Fig. 2d) show that the Sb₂S₃ nanorods exhibit the patterns of orthorhombic Sb₂S₃ (stibnite) (JCPDS No.42-1393), and are preferentially oriented with [211] direction normal to substrate plane (i.e., [211]-oriented) with the intensity (I) ratio of I₍₂₁₁₎/I₍₁₃₀₎ = 2.20, much larger than the I₍₂₁₁₎/I₍₁₃₀₎ = 0.82 for the powder sample, indicating that most of the nanorods have the out-of-plane growths along [001], [010], and [100] directions with (001), (100), and (010) planes tilting the average angles of ca. 37^o, 57^o, and 74^o on substrate plane, respectively (Inset to Fig. 2d). Moreover, the TiO₂ film on fluorine-doped tin oxide (FTO) substrate exhibits the strongest XRD peak of (101) plane of anatase TiO₂ (101) planes (Fig. 2d), inferring that the (101) crystal planes of TiO₂ nanoparticles exposed on the film surface orient in the direction parallel to substrate plane^19,21.

Fig. 2: Characterization of the Sb₂S₃ nanorod array grown on TiO₂ film.

a Bird-view and (b) cross-sectional scanning electron microscopy images, (c) transmission electron microscopy (TEM) image, and (d) X-ray diffraction patterns. The insets to c are the selected area electron diffraction pattern (top-right) of the nanorod and the high-resolution TEM image from the marked part of the nanorod (bottom-left). Shown by the inset to d, the angle between a specific Sb₂S₃ crystal plane and the substrate is derived from the angle between the crystal plane and the Sb₂S₃ [211] direction that is normal to the substrate plane. The Sb₂S₃ nanorod array is prepared at the seed growth solution concentration of C_sg = 0.1 M, the nanorod growth solution concentration of C_ng = 0.8 M, and the repetition of spin-coating and annealing cycle number of n = 2.

Growth mechanism of nanorod array

When Sb₂S₃ seeds are absent from the TiO₂ film, only the large and irregular Sb₂S₃ grains without a preferential orientation are obtained (Supplementary Fig. 1). Clearly, the Sb₂S₃ crystalline seeds are very crucial for the growth of the Sb₂S₃ nanorod array. Our results show that the TiO₂ film surface undoubtedly is decorated by the preferentially [211]/[221]-oriented Sb₂S₃ crystalline seeds that are of a C_sg-dependent size in the range of C_sg = 0.05−0.4 M (Supplementary Fig. 2 and Supplementary Note 1), where the [211]/[221]-oriented Sb₂S₃ seeds of irregular shapes (laterally ca. 60−450 nm) derived from C_sg = 0.4 M are clearly observed and similar to our previous observations¹⁹, but the [211]/[221]-oriented Sb₂S₃ seeds from C_sg = 0.05−0.2 M are not observable under SEM. Hence, the [211]/[221]-oriented seeds formed at C_sg ≤ 0.2 M are very tiny and must be <10 nm in size as evaluated from the SEM resolution.

In order to get insights into the growth mechanism for the Sb₂S₃ nanorod array on TiO₂ film, the dependences of the Sb₂S₃ nanorod growth on the conditions (C_sg, C_ng, and n) for the seed-assisted RSCA method are investigated. We find that the formation of [211]-oriented Sb₂S₃ nanorod array depends exclusively on the C_sg concentration, and the tiny [211]/[221]-oriented Sb₂S₃ seeds derived from a dilute SG solution of C_sg ≤ 0.1 M are indispensable (Supplementary Fig. 3 and Supplementary Note 2); moreover, both C_ng and n in the ranges of C_ng < 1.1 M and n < 3 imposes no remarkable influences on the lateral size of Sb₂S₃ nanorods, but the lateral size of Sb₂S₃ nanorods gets dramatically increased to make an exception when both C_ng = 1.1 M and n = 3 (Supplementary Fig. 4 and Supplementary Note 3). Here, we mainly show the length and number density of Sb₂S₃ nanorods dependent on C_ng and n values. The increase in either C_ng or n leads to the Sb₂S₃ nanorod growth along c-axis (i.e., a larger array thickness) (Fig. 3a). However, for C_ng < 1.1 M and n < 3, the number density is mainly affected by C_ng, and an increased C_ng results in a significantly reduced number density (Fig. 3b). Actually, the number density of Sb₂S₃ nanorods in most cases is not remarkably changed with n value, strongly indicating that the Sb₂S₃ nanorods mainly take a self-nucleated growth similar to the Sb₂S₃ cuboid growth¹⁹.

Fig. 3: Influencing factors and epitaxial growth model.

a Length and (b) number density of Sb₂S₃ nanorods dependent on the nanorod growth solution concentration (C_ng) and the  repetition of spin-coating and annealing (RSCA) cycle number (n) at the seed growth solution concentration of C_sg = 0.1 M. c Schematic illustration of the tiny-seed-governed orientation-competing-epitaxial nucleation/growth model for the growth of [211]-oriented Sb₂S₃ single-crystalline nanorods on polycrystalline TiO₂ nanoparticle film, where few other [221]-oriented ones, for example, might be possible. The inset to c shows the Sb₂S₃ crystal growth during the first RSCA cycle (n = 1), where the most competitive TiO₂ nanoparticle orientation for the Sb₂S₃ nanorod to nucleate/grow features its (101) plane tilting α = 22^o (Particle 2), but the nucleation/growth at other sites around (e.g., Particle 1 of α = 0^o and Particle 3 of α = 33^o) is relatively much slower or suppressed due to their unfavorable TiO₂ (101) plane orientations for the [211]-oriented Sb₂S₃ crystal growth.

The growth of Sb₂Se₃ nanorod arrays on selenized Mo substrate by sublimation technique was proposed to follow a split growth model on the basis of the morphology transition from thin film to nanorod array when the lateral stress of the Sb₂Se₃ grains was beyond the tolerance of the van der Waals forces between (Sb₄Se₆)_n ribbons during Sb₂Se₃ vapor evaporates²². Besides, quasi-epitaxial growth and dual-selection growth models were proposed, respectively, to elucidate the formation of [hk1]-oriented Sb₂S₃ compact films on TiO₂²³ and CdS²⁴ nanoparticle films. Unfortunately, there is no growth mechanism for the formation of Sb₂S₃ single-crystalline nanorods on polycrystalline surfaces. We previously proposed an OCE nucleation/growth mechanism to describe the growth of the Sb₂S₃ cuboids on polycrystalline TiO₂ film¹⁹. In the OCE nucleation/growth mechanism, the epitaxial relationship of Sb₂S₃ (101)//TiO₂ (101) exists for the heterogeneous nucleation of orthorhombic Sb₂S₃ crystal on favorably oriented anatase TiO₂ nanoparticle surface as the result of the lattice match between Sb₂S₃ (101) and TiO₂ (101) planes.

Known from the experimental evidence (Supplementary Fig. 5 and Supplementary Note 4), the growth of the [211]-oriented Sb₂S₃ nanorods essentially follows an OCE nucleation/growth mechanism that is similar to that proposed for the growth of [221]-oriented Sb₂S₃ cuboids¹⁹. As shown by calculated results (Supplementary Table 1, Supplementary Fig. 6 and Supplementary Note 5), a TiO₂ nanoparticle theoretically needs to orientate with its (101) plane tilting in an angle range of 0^° ≤ α < 45^° on TiO₂ film plane in order to act as an effective epitaxial nucleation site for the out-of-plane growths of [211]-oriented Sb₂S₃ crystals along [001], [100], and [010] directions; moreover, considering the fact that the [211]-oriented Sb₂S₃ nanorods normally have the (001) planes tilting an angle of ca. 37^° on substrate plane (Inset to Fig. 2d), it is revealed that the most competitive TiO₂ nanoparticle orientation for the heterogeneously epitaxial nucleation/growth of the [211]-oriented Sb₂S₃ single-crystalline nanorods features the tilting angle of α ≈ 22^o (Supplementary Table 1, Supplementary Fig. 6 and Supplementary Note 5).

Actually, the intrinsic distinction between the growths of [211]-oriented Sb₂S₃ nanorods and [221]-oriented Sb₂S₃ cuboids originates from the sizes of [211]/[221]-oriented Sb₂S₃ seeds, that is, tiny Sb₂S₃ seeds of laterally <10 nm in size for nanorods and large seeds of laterally 60−400 nm in size for cuboids (Supplementary Fig. 5 and Supplementary Note 4). Since the orientation and morphology of resulting Sb₂S₃ nanorods are exclusively governed by the tiny size of [211]/[221]-oriented Sb₂S₃ seeds (Supplementary Figs. 2−5 and Supplementary Notes 1−4), we here propose a tiny-seed-governed OCE nucleation/growth mechanism for the growth of [211]-oriented Sb₂S₃ nanorods and consequently the nanorod array on polycrystalline TiO₂ nanoparticle film (Supplementary Table 1, Supplementary Fig. 6 and Supplementary Note 5), as schematically illustrated in Fig. 3c. First, to form the [211]/[221]-oriented Sb₂S₃ seeds, the heterogeneously epitaxial nucleation of orthorhombic Sb₂S₃ crystals occurs on favorably oriented anatase TiO₂ nanoparticles, followed by the self-nucleated growth of the Sb₂S₃ nuclei into the Sb₂S₃ tiny crystalline seeds when the orientations of the Sb₂S₃ nuclei and the resulting seeds are governed by TiO₂ (101) plane orientations. Since the TiO₂ nanoparticles are relatively much larger (10−20 nm in diameter), each of the tiny Sb₂S₃ seeds (laterally <10 nm in size) has a high possibility to be formed on an individual TiO₂ nanoparticle. Those tiny Sb₂S₃ seeds inevitably have more chances to nucleate [211]-oriented Sb₂S₃ crystals, because more TiO₂ nanoparticles tending to orient with (101) face parallel to the film plane (Fig. 2d) inevitably result in more TiO₂ nanoparticle orientations in the range of 0^o ≤ α < 45^o (Supplementary Table 1, Supplementary Fig. 6 and Supplementary Note 5). During the first RSCA process (i.e., n = 1), the Sb₂S₃ seeds on the TiO₂ nanoparticles oriented with α ≈ 22^o (Particle 2) have the fastest rate to take a self-nucleated growth into laterally enlarged Sb₂S₃ nanorods without changing the seed orientations, where the growth rate in c-axis direction is much larger than in a-axis and b-axis directions, but the Sb₂S₃ crystal growths at around sites with the less-competitive/unfavorable orientations (Particles 1 and 3) are merged or suppressed (Inset to Fig. 3c). Furthermore, with increasing the RSCA cycle numbers (n > 1), the growing nanorods still take a self-nucleated growth mainly along c-axis direction into longer ones without remarkably increasing lateral size (Supplementary Fig. 4 and Supplementary Note 3) nor changing the orientation (Supplementary Fig. 7). During the nanorod growth, the sites, where there are no seeds and/or suppressed/merged crystal growths, create the empty interspaces in between nanorods, leading to the array of [211]-oriented Sb₂S₃ nanorods on the polycrystalline TiO₂ nanoparticle film. On the other hand, the tiny Sb₂S₃ seeds have a weak capability to merge and suppress the far-off nucleation at around sites due to their small size, leading to the number density (ca. 36 µm⁻²) of Sb₂S₃ nanorods much higher than that (ca. 3 µm⁻²) of the  Sb₂S₃ cuboids derived from large Sb₂S₃ seeds (Supplementary Figs. 3, 5 and Supplementary Notes 2, 4).

Solar cell architecture and working principle

The Sb₂S₃ nanorod array in the Sb₂S₃/TiO₂-NHJ shows an absorption onset at about 737 nm corresponding to the optical band gap (E_g) of 1.68 eV (Fig. 4a), very close to the E_g value of Sb₂S₃ bulk crystal²⁵. With the valence band level (E_VB) of E_VB = −5.29 eV for the Sb₂S₃ bulk crystals²⁵, the conduction band level (E_CB) of the Sb₂S₃ nanorod array is calculated to be −3.61 eV from the E_g value; moreover, the PTB7-th exhibits the absorption mainly within 550−750 nm with an absorption onset at ca. 772 nm, corresponding to E_g = 1.61 eV, the highest occupied molecular orbital (HOMO) energy level of −5.20 eV and the lowest unoccupied molecular orbital (LUMO) energy level of −3.59 eV²⁶. With the E_CB = −4.20 eV for TiO₂²⁷, the E_CB = −2.30 eV and E_VB = −5.30 eV for MoO₃²⁸ and the work-functions (W_f) of −4.40 eV for FTO and −5.10 eV for Au²⁹, the combination of Sb₂S₃/TiO₂-NHJ and PTB7-th renders a ternary materials system, where TiO₂, Sb₂S₃ and PTB7-th components form type II hybrid heterojunctions with staggered band alignments to facilitate the charge transport in solar cell (Fig. 4b). The Sb₂S₃ nanorod has a faceted top of (001) plane and the faceted sides of (100) and (010) planes (Fig. 2c). Density functional theory (DFT) calculations reveal that the S atoms of PTB7-th interact weakly with the Sb atoms on Sb₂S₃ (001) and (100) planes, but no evident interfacial electric field is formed at the Sb₂S₃/PTB7-th interfaces (Fig. 4c); moreover, the Sb₂S₃/PTB7-th interfaces formed by other Sb₂S₃ planes (e.g., (010) and (110) ones) have the same interfacial interaction features (Supplementary Fig. 8). The wavefunctions of electronic states near the conduction band minimum (CBM) and valence band maximum (VBM) of the Sb₂S₃/PTB7-th interfaces formed by Sb₂S₃ (100) plane (Fig. 4d), and Sb₂S₃ (001), (010), and (110) planes as well (Supplementary Fig. 8), are also calculated using DFT. It is found that the CBM wavefunction only locates in the Sb₂S₃ material, but the VBM wavefunction predominantly distributes within the PTB7-th phase and slightly overlaps the Sb₂S₃ material, which is similar to the well-known efficient poly(3-hexylthiopene)/[6,6]-phenyl-C61-butyric acid methyl ester (i.e., P3HT/PCBM) organic photovoltaic system³⁰. Those unique separate electronic states inevitably favor the dissociation and injection of photogenerated electron-hole pairs and free charge carriers at the Sb₂S₃/PTB7-th interface upon the excitation of Sb₂S₃ and PTB7-th, with the photogenerated electrons remaining in or being injected into Sb₂S₃ crystal and the photogenerated holes being injected into or remaining in PTB7-th phase.

Fig. 4: Material property and architecture-dependent device performance.

a Absorption spectra of Sb₂S₃/TiO₂-nanoarray heterojunction (Sb₂S₃/TiO₂-NHJ) and PTB7-th film. b Energy level diagram and charge transfers amongst the materials in the solar cells of this work, where the photogenerated electrons injected into the conduction band of TiO₂ may be trapped (k_tr) by surface states (SS) and the trapped electrons can be thermally released (k_dt) into the conduction band for transport. c Density functional theory (DFT) calculated interfacial interaction at the Sb₂S₃/PTB7-th interfaces formed by PTB7-th on Sb₂S₃ (001) plane (above) and Sb₂S₃ (100) plan (below), where red color denotes charge accumulation and green color stands for charge depletion, and the electron density isosurface is set to 0.0004 e Å⁻³. d DFT calculated wavefunctions of electronic states near the conduction band minimum (CBM) (above) and valence band maximum (VBM) (below) of Sb₂S₃(100)/PTB7-th interface system, where the isosurface of wavefunction square is set to 4 × 10⁻¹¹ e Å⁻³. e Structure of p-(p + i)-n solar cell (up-left: cross-sectional scanning electron microscopy image of a practical device; below-right: schematic illustration), where p means p-type PTB7-th, i means intrinsic Sb₂S₃ and n means n-type TiO₂. f Current-voltage curves under AM1.5 G illumination, (g) incident photon-to-current conversion efficiency (IPCE) spectra (the right-hand axis shows the integrated photocurrent from the IPCE spectra) and (h) electrochemical impedance spectroscopic spectra (scatter: experimental data measured in the dark; solid line: fitting data) of the p–i–n and p-(p + i)-n solar cells with the highest efficiency. Inset to (h) shows the equivalent circuit used for fitting, which consists of series resistance (R_s), a hole transport and extraction resistance (R_t) at selective electrode, a recombination resistance (R_rec), and two constant phase angle elements (CPE).

The p–i–n solar cells (Fig. 1) are fabricated by spin-coating PTB7-th onto the Sb₂S₃/TiO₂-NHJ. Optimizing experiments demonstrate that the p–i–n solar cell delivers a peak efficiency when the Sb₂S₃/TiO₂-NHJ is derived either at C_sg = 0.1 M (Supplementary Table 2) or at C_ng = 0.8 M and n = 2 (Supplementary Table 3). In the following sections, therefore, we focus on the optimized Sb₂S₃/TiO₂-NHJ with a length of 320 nm prepared at C_sg = 0.1 M, C_ng = 0.8 M, and n = 2, to investigate the architecture-dependent performance of solar cells. In order to reveal the photovoltaic characteristics of the p–i–n architecture, we also fabricate the p-(p + i)-n counterpart solar cells for comparison, which features the polymer infiltration into the interspaces between nanorods resulting in the (p + i) blend of intrinsic Sb₂S₃ nanorods and p-type polymer sandwiched between p-type HTL and n-type ETL (Fig. 4e). To further visualize the PTB7-th infiltration into the Sb₂S₃ nanorod array, the SEM image with a larger view field is shown in Supplementary Fig. 9. As illustrated by Fig. 4b, in both p–i–n and p-(p + i)-n devices, free charge carriers are generated in the Sb₂S₃ nanorods upon illumination as the result of the small exciton binding energy (E_b) in Sb₂S₃ (E_b < 10 meV) that is much smaller than the thermal energy at ambient temperature (k_BT ∼ 26 meV, where k_B is Boltzmann constant and T is temperature)³¹; however, the excitons generated in the PTB7-th phase upon illumination need to diffuse to the Sb₂S₃/PTB7-th interface for dissociation into free electrons (e) and holes (h), with the electrons being injected into Sb₂S₃ and the holes remaining in the polymer, because the E_b in conjugated polymers is normally about 0.4−0.5 eV^32,33, much higher than k_BT ~ 26 meV. Due to the high electron (μ_e ≈ 10 cm² V⁻¹ s⁻¹) and hole (μ_h ≈ 2.6 cm² V⁻¹ s⁻¹) mobilities in the Sb₂S₃ crystals¹⁹, the electrons generated in the Sb₂S₃ nanorods are quickly transported into the TiO₂ layer under the built-in electric field (E_bi) due to the W_f difference between the FTO and Au asymmetric collection electrodes^34,35, while the holes generated in the Sb₂S₃ nanorods are injected into the PTB7-th layer on nanorod top in the p–i–n device or the PTB7-th layers on nanorod top and amongst nanorods in the p-(p + i)-n device. As the result of quite a low hole mobility in the PTB7-th (μ_h = 2.8 × 10⁻⁴ cm² V⁻¹ s⁻¹)³⁶, the transport of holes in the polymer phase proceeds very slowly and they finally tunnel through the MoO₃ layer acting as an optical spacer/Au diffusion barrier³⁷ to reach the Au electrode (Fig. 4b).

Architecture-dependent photovoltaic performance

The current-voltage (J−V) curves for the best p–i–n and p-(p + i)-n solar cells are presented in Fig. 4f, and the averaged performance of seven independent devices for each sample are summarized in Table 1. We achieve a champion efficiency of η = 5.70% in the p–i–n solar cells, which is comparable to those of the planar heterojunction solar cells based on the Sb₂S₃ nanoparticle films prepared by atomic layer deposition (e.g., η = 5.77%)¹⁵ and solution-processing method (e.g., η = 5.69%)¹⁷. Furthermore, the p–i–n solar cell exhibits good stability and retains 88% of the initial efficiency after storage in the N₂ atmosphere for 31 days (Supplementary Fig. 10). Generally, the p–i–n solar cells are much more efficient than the p-(p + i)-n ones. The averaged fill factor (FF) of the p–i–n devices is ca. 54%, much higher than that (ca. 38%) of the p-(p + i)-n ones, and the increased FF in the p–i–n devices is accompanied by a larger open-circuit voltage (V_oc), a higher short-circuit current (J_sc), a smaller series resistance (R_s), and a larger shunt resistance (R_sh) (Table 1). Moreover, incident photon-to-current conversion efficiency (IPCE) spectra show that both p–i–n and p-(p + i)-n solar cells have a similar and broad spectral profile in 300−750 nm, and the integrated J_sc from the IPCE data (i.e., 14.78 mA cm⁻² for p–i–n device and 12.51 mA cm⁻² for p-(p + i)-n) one) agree with the observed J_sc values (Fig. 4g). While the polymer amount accumulated in the p-(p + i)-n architecture is higher than that in the p–i–n one, the IPCE of the former device in the whole absorption range is significantly lower than that of the latter one, inferring that much more photogenerated charge carriers contribute to photocurrent generation in the p–i–n device.

Table 1 Architecture-dependent photovoltaic performance of the p–i–n and p-(p + i)-n solar cells^a.

The electrochemical impedance spectroscopy (EIS) spectra in the form of Nyquist plots for the p–i–n and p-(p + i)-n solar cells measured in the dark are shown in Fig. 4h. The EIS data are fitted with a transmission line model³⁸ as depicted by a simplified equivalent circuit (Inset to Fig. 4f), and the fitting results are summarized in Supplementary Table 4. On the Nyquist plots, the small arc at high frequency is for the hole transport and extraction resistance (R_t) at the selective electrode, while the main arc at low frequency is attributed to recombination resistance (R_rec) in the devices³⁹. The p–i–n solar cell has R_t = 32.08 Ω cm², much smaller than that (R_t = 74.58 Ω cm²) of the p-(p + i)-n device, inferring that the p–i–n architecture has a much easier hole the transport and extraction at the MoO₃/Au cathode. Moreover, the R_rec value (204.66 Ω cm²) in the p-(p + i)-n device is lower than that (260.10 Ω cm²) in the p–i–n one, suggesting much more serious charge recombination in the p-(p + i)-n device. Furthermore, the series resistance (R_s) in the p-(p + i)-n device (4.42 Ω cm²) is also much higher than that (1.73 Ω cm²) in the p–i–n device, in agreement with the change in the R_s data measured by J−V characterization (Table 1).

Photogenerated electric field effects on photovoltaic performance

Obviously, the p–i–n and p-(p + i)-n hybrid solar cells exhibit a significantly architecture-dependent performance. A similar phenomenon has also been observed in the hybrid solar cells based on CdS nanorod array and conjugated polymer MEH-PPV, that is, a partial infiltration of MEH-PPV into CdS nanorod array results in better device performance, but for which the reason is not clear yet⁴⁰. Since materials in the p–i–n and p-(p + i)-n solar cells are the same, the observed different photovoltaic performances of them, in particular, the change in V_oc, must be governed by an architecture-dependent factor other than the materials property (e.g., band energy level difference).

Here, an architecture-dependent charge distribution (ADCD) model is proposed on the basis of light intensity attenuation in the polymer phase to describe the photovoltaic behaviors in the p–i–n and p-(p + i)-n devices, as schematically illustrated in Fig. 5a. In the p-(p + i)-n solar cells, the polymer phase in between the interspaces of Sb₂S₃ nanorods in the Sb₂S₃/TiO₂-NHJ has a thickness comparable to the Sb₂S₃ nanorod length (i.e., 320 nm) and a remarkable absorption contribution to the generation of charge carriers. Actually, the light intensity (I) is exponentially attenuated when passing a polymer film according to the relation of I = I₀ exp(−αx), where I₀ is the incident light intensity, I is the intensity at the illumination depth of x nm and α is the absorption coefficient of PTB7-th. Given α = 10⁵cm⁻¹ for PTB7-th, the illumination intensity at the illumination depth of 320 nm is attenuated to 4% of the incident intensity (Fig. 5b). As the result of the light intensity attenuation, therefore, in the p-(p + i)-n solar cells, there is an original photogenerated hole density gradient from a higher density at the stem region of Sb₂S₃ nanorods to a lower one at the nanorod top region, after the photogenerated excitons in PTB7-th phase dissociate at the nanorod-side-face-formed Sb₂S₃/PTB7-th interface with electrons being injected into the Sb₂S₃ nanorods. However, the conductivity of a conjugated polymer in pristine form is very low, but it increases under the photo-doped state due to illumination. Hence, subjected to the light intensity attenuation, the conductivity in the PTB7-th film inevitably becomes gradually reduced along the illumination depth direction. Consequently, during the photovoltaic process, with the diffusion of the photogenerated holes in between the interspaces amongst the Sb₂S₃ nanorods towards the MoO₃/Au collection electrode, a reversed hole density gradient is formed from a higher hole density at the nanorod top region to a lower one at the Sb₂S₃ nanorod stem region, to render a photogenerated electric field (E_ph) opposite to the E_bi in p-(p + i)-n architecture (right, Fig. 5a). In contrast, the p–i–n solar cells are actually an E_ph-free device (left, Fig. 5a), due to the absence of the polymer infiltration into the interspaces in between the Sb₂S₃ nanorods and thereby no photogenerated hole accumulation around each of Sb₂S₃ nanorods.

Fig. 5: Architecture-related photogenerated electric field effects.

a Architecture-dependent charge distribution model showing the absence and presence of photogenerated electric field (E_ph) effects in the p–i–n (left) and p-(p + i)-n (right) solar cells during the photovoltaic process, where p means p-type PTB7-th to form hole transporting layer (HTL), i means intrinsic Sb₂S₃ and n means n-type TiO₂ to form electron transporting layer (ETL). b Light attenuation in the polymer film, where the inset illustrates the distribution of photogenerated excitons in the PTB7-th film along the illumination depth direction with a deeper purple for a smaller exciton population. c, d Simulated electric potential distributions in (c) p–i–n and (d) p-(p + i)-n solar cells at 1 V forward bias, where, E_v stands for the bias electrostatic field, the color gradient bars depict the electric potential, and the substrate and electrode refer to the FTO/TiO₂ and MoO₃/Au layers in solar cells, respectively. e The capacitance-voltage data in the dark for the p–i–n and p-(p + i)-n solar cells with the highest efficiency.

We performed the finite element simulation of the bias electrostatic filed (E_v) distribution in the p–i–n and p-(i + p)-n solar cells at a forward bias of 1 V, to mimic the case of selectively accumulated charge carriers in the devices, using the AC/DC module on commercial COMSOL Multiphysics software. The structural models for simulation (Supplementary Table 5) were built by assuming that all the nanorods of a right square prism shape are vertically grown on the substrate and the materials dimensions are given by referring to the SEM results (Figs. 1 and 4e). All the parameters of materials property and dimensions for simulation are provided in Supplementary Table 5. In the p–i–n solar cell, where the FTO/TiO₂ and MoO₃/Au layers are defined, respectively, as substrate and electrode for a simplified simulation, the highest electrostatic potential dominantly at the electrode gets gradually reduced along normal of HTL (i.e., PTB7-th film) until the lowest potential at the top of Sb₂S₃ nanorods, suggesting the presence of a very narrow bias E_v due to the forward bias in the HTL to drive electrons to move towards the electrode (Fig. 5c). As for the p-(p + i)-n solar cell, however, the highest electrostatic potential at the electrode gradually decreases along the axis of Sb₂S₃ nanorods until the lowest potential at the substrate, informing a bias E_v across the whole device to drive electrons to reach the electrode (Fig. 5d). Clearly, the simulations inform that, as the results of the hole accumulation near the nanorod top region and the electron accumulation in the nanorod stem region, an electric field (e.g., E_v) across the whole device and opposite to the built-in field E_bi is theoretically expected to reduce the collection efficiency of electrons photogenerated inside both Sb₂S₃ nanorods and PTB7-th phase in the p-(p + i)-n solar cells, while only an electric field across the HTL negatively affects the very limited population of electrons at the top of Sb₂S₃ nanorods and in the HTL in the p–i–n devices. The finite element simulations theoretically agree with the expectations from the ADCD model (Fig. 5a). Interestingly, further simulated results inform that once the PTB7-th infiltration into the Sb₂S₃ nanorod array is present either partially or completely, a similar electric field (e.g., E_v) distribution across the whole device and opposite to the built-in field E_bi is theoretically expected (Supplementary Fig. 11 and Supplementary Note 6).

The ADCD model in Fig. 5a is also experimentally supported by the capacitance-voltage (C−V) measurements of solar cells in the dark. The C−V results (Fig. 5e) clearly show that the p-(i + p)-n device has a much higher capacitance (C) than the p–i–n one in the whole bias voltage (V) range, indicating a higher capability of being charged by injected charge carriers in the former device. The relation C⁻² ∝ V in Mott-Schottky (MS) equation expresses the relationship amongst the C, built-in potential (V_bi), and V of the solar cell, and the intersection point of the linear slope with voltage axis on the C⁻²−V plot stands for the V_bi across the device^41,42. The V_bi of the p–i–n device is 0.68 V (Inset to Fig. 5e), comparable to the maximum theoretical V_bi (0.70 eV) in the solar cell due to the W_f difference between Au (−5.10 eV) and FTO (−4.40 eV) electrodes²⁹. Reasonably, the V_bi obtained from the C⁻²−V plots originates from the W_f offset of the asymmetric electrodes⁴³. However, the p-(p + i)-n device exhibits a V_bi of 0.57 V, much lower than the maximum theoretical V_bi value. The V_bi values of the p–i–n and p-(p + i)-n devices are almost identical to their V_oc data from J−V measurements (Table 1).

The smaller V_bi accompanied by a higher capacitance at V < V_bi in the p-(i + p)-n cell suggests that a negative impact on the V_bi results from the charge accumulation therein^44,45. Since there is a greatly unbalance amongst the hole mobility in PTB7-th phase (μ_h = 2.8 × 10⁻⁴cm² V⁻¹ s⁻¹) and the electron mobilities in TiO₂ (μ_e ≈ 1.98 × 10⁻³ cm² V⁻¹ s⁻¹)²⁵ and Sb₂S₃ (μ_e ≈ 10 cm² V⁻¹ s⁻¹), the charge accumulation predominantly correlates with the low-mobility holes in the polymer phase in between Sb₂S₃ nanorods under forwarding bias⁴⁶, resulting in a distribution state of a high hole density toward the MoO₃/Au cathode and a high electron density in the TiO₂ layer. Clearly, the C−V measurements reveal that an opposite electric field opposite to the E_bi across the device is generated due to the hole accumulation in the polymer phase in between the Sb₂S₃ nanorods with a distribution of a high density toward the MoO₃/Au cathode to impose a significant negative impact on the device V_oc in the p-(i + p)-n solar cells, but it is not the case in the p–i–n devices, in agreement with the expectations from the ADCD model.

Both the finite element simulation and C−V data demonstrate that the E_ph negative effect on charge transport is much more noteworthy in p-(p + i)-n devices than in p–i–n ones. From the ADCD model (Fig. 5a), the following points can be expected: (i) the p–i–n device only have nanorod-top-face-formed Sb₂S₃/PTB7-th interface, but the p-(p + i)-n device has additionally a much larger area of nanorod-side-face-formed Sb₂S₃/PTB7-th interface besides the nanorod-top-face-formed one and a heavy photogenerated hole accumulation around each of nanorods, hence, the p-(p + i)-n device should have a much shorter electron lifetime (τ_e1) due to the charge recombination at the Sb₂S₃/PTB7-th interface; (ii) as the TiO₂/Sb₂S₃ interface in the p–i–n and p-(p + i)-n solar cells are the same, the electron lifetime (τ_e2) subjected to the charge recombination at TiO₂/Sb₂S₃ interface in two solar cells should be comparable to each other, but, a more serious charge recombination at the TiO₂/Sb₂S₃ interface for a shorter τ_e2 should practically exist in the p-(p + i)-n device because the photogenerated holes in Sb₂S₃ nanorods do not easily escape from the TiO₂/Sb₂S₃ interface under the E_ph effect; (iii) the transit time (τ_D) for the photogenerated electrons accommodating in the TiO₂ layer to reach the FTO collection electrode in the p-(p + i)-n device should be much longer than that in the p–i–n one, because the electron diffusion rate greatly depends on the electron concentration in conduction band⁴⁷, and the electron density injected into the TiO₂ layer in the p-(p + i)-n device is much lower due to the serious loss by interfacial charge recombination at nanorod-side-face-formed Sb₂S₃/PTB7-th interface. Those expectations are confirmed by dynamic spectroscopic investigates of solar cells (Supplementary Fig. 12 and Supplementary Note 7). That is, the τ_e1 = 31.53 ms of the p-(p + i)-n device is much shorter than the τ_e1 = 895.44 ms in the p–i–n device; both cells have comparable τ_e2 values, but the τ_e2 = 38.74 μs for p-(p + i)-n cell is shorter than the τ_e2 = 61.22 μs for p–i–n cell; moreover, the τ_D = 19.55 μs in p-(p + i)-n device is much longer than τ_D = 6.25 μs in the p–i–n device (Table 1). Surprisingly, the p–i–n device has a very long electron lifetime τ_e1 related to the charge recombination at Sb₂S₃/PTB7-th interface, which is attributed to a very low electron population at the top of Sb₂S₃ nanorods in such device as the result of the high electron transport in the Sb₂S₃ nanorods toward FTO collection electrode. Moreover, the very low electron population in the Sb₂S₃ nanorod top region inevitably results in a limited negative impact of the E_ph across the HTL on the transport of photogenerated electrons in the p–i–n device, as indicated by the finite element simulation (Fig. 5c).

Advantages of the p–i–n solar cells

The observed J−V characteristic difference between the p–i–n and p-(p + i)-n devices (Fig. 4f and Table 1) gets well accounted for by the proposed ADCD mode (Fig. 5a). First, after ruling out the band energy level influence on the device V_oc, it is reasonable to believe that the V_oc of the Sb₂S₃/TiO₂-NHJ solar cells strongly correlates with the net built-in field in device^34,35. Since the p–i–n solar cell under illumination has no negative E_ph effects as the result of lacking the photogenerated hole accumulation around Sb₂S₃ nanorods, the device has a higher V_oc resulting from the E_bi across the device due to the asymmetric W_f values between Au and FTO electrodes, while there are the negative E_ph effects to weaken the device E_bi under illumination for a reduced V_oc in the p-(i + p)-n device. Moreover, a larger nanorod-side-face-formed Sb₂S₃/PTB7-th interface area and the negative E_ph effects on charge transport are existing in the p-(i + p)-n device, which deteriorate the charge collection efficiency to reduce reduced J_sc and FF by causing a great electrical coupling between the electrons inside Sb₂S₃ nanorods and the holes in PTB7-th phase with a small R_sh for serious interfacial charge recombination at the Sb₂S₃/PTB7-th interface and a high R_s for a retarded charge transport.

Obviously, the p–i–n solar cells based on Sb₂S₃/TiO₂-NHJ have at least two unique advantages: one is that no negative E_ph effects on device performance thanks to lacking the photogenerated hole accumulation around Sb₂S₃ nanorods; another is the [211]-oriented single-crystalline Sb₂S₃ nanorods therein provide the straightforward and highly effective charge transport channels formed by the covalently bounded (Sb₄S₆)_n ribbons tilting ca. 53^° on the substrate for the photogenerated charge carriers inside the Sb₂S₃ layer to reach collection electrode. Indeed, those advantages make the Sb₂S₃/TiO₂-NHJ have a good application potential to efficient solar cells. However, while the efficiency (η = 5.70%) of our p–i–n solar cells is comparable to the hybrid planar heterojunction solar cell based on the Sb₂S₃ nanoparticle film (i.e., FTO/TiO₂/Sb₂S₃/P3HT/Au (η = 5.77%)¹⁵, but it is still much lower than those of all-inorganic solar cells of the Sb₂Se₃ nanorod array on Mo-substrate (i.e., Mo/MoSe₂/Sb₂Se₃/CdS/ZnO/ZnO:Al/Ag, η = 9.2%)²² and the InP p–n junction nanowire array on InP substrate (i.e., InP/InP(p-n)/ITO, η = 17.8%)⁹. For a simple preparation of an HTL over the Sb₂S₃/TiO₂-NHJ without infiltration into the interspaces in-between nanorods, we chose the good-performing polymer PTB7-th to prepare organic HTL. The not high efficiency in our p–i–n solar cells reasonably originates from the low hole mobility in PTB7-th and the week interaction between the PTB7-th and Sb₂S₃ as indicated by DFT calculation (Fig. 4c) for a not very efficient Sb₂S₃/PTB7-th interface for charge separation and injection⁴⁸. We envision that in order to get the more efficient p–i–n solar cells based on the Sb₂S₃/TiO₂-NHJ, it is necessary to replace the polymer HTL with one of high hole mobility, where the HTL material needs to be prevented from infiltrating into the interspaces in-between nanorods, for example, by planarizing the nanoarray with an isolating substance infiltrated into the interspaces among nanorods^8,9,11.

Discussion

In summary, the Sb₂S₃ nanorod array on polycrystalline TiO₂ film is realized by a tiny-seed-assisted RSCA method, offering a nanoarray heterojunction system (i.e., Sb₂S₃/TiO₂-NHJ) with the preferentially [211]-orientated Sb₂S₃ nanorods of lateral sizes much larger than the beneath TiO₂ nanoparticles. The Sb₂S₃ nanorods in the Sb₂S₃/TiO₂-NHJ have a bandgap of E_g = 1.68 eV and provide straightforward and highly effective charge transport channels formed by covalently bounded (Sb₄S₆)_n ribbons. Those make the Sb₂S₃/TiO₂-NHJ have a good potential for efficient solar cells. By depositing a conjugated polymer PTB7-th film as a hole transporting layer onto the top of the Sb₂S₃/TiO₂-NHJ, a hybrid solar cell with a p–i–n architecture is obtained. The p–i–n solar cells exhibit a V_oc strongly correlating with the net built-in field across devices and no negative E_ph impact on device performance. With the two unique advantages of the straightforward and effective charge transport channels and the absence of negative E_ph impact on device performance, a power conversion efficiency of 5.70% is achieved in the p–i–n solar cells for an optimized nanoarray structure (i.e., 320-nm-thick Sb₂S₃ nanoarray prepared at C_sg = 0.1 M, C_ng = 0.8 M and n = 2). Due to the absence of the negative E_ph impact, the p–i–n solar cells exhibit a long electron lifetime on the sub-second scale (τ_e1 ≈ 0.9 s) related to the charge recombination at Sb₂S₃/PTB7-th interface.

The orientation and morphology of the resulting Sb₂S₃ nanorods on the TiO₂ film are exclusively governed by the tiny [211]/[221]-oriented Sb₂S₃ seeds of laterally <10 nm in size. The growth of the [211]-oriented Sb₂S₃ nanorods follows a tiny-seed-governed OCE nucleation/growth mechanism according to the epitaxial relationship of Sb₂S₃ (101)//TiO₂ (101), where the Sb₂S₃ seeds on the TiO₂ nanoparticles oriented with α ≈ 22^° are most competitive during the followed RSCA cycles (n ≥ 1) to take a self-nucleated growth into much larger [211]-oriented Sb₂S₃ nanorods with a much larger growth rate in c-axis direction than in a-axis and b-axis directions. During the tiny-seed-governed OCE nucleation/growth process, the sites there are no seeds and/or suppressed/merged crystal growths create the empty interspaces in between nanorods for the formation of the Sb₂S₃ nanorod array. Moreover, the length of nanorods (i.e., nanoarray thickness) strongly depends on both C_ng and n values; the lateral size and the number density of the Sb₂S₃ nanorods are mainly obtained at the first RSCA cycle (n = 1) and not n-dependent, while they are affected by C_ng and increasing C_ng leads to increased lateral size and reduced number density.

In the Sb₂S₃/PTB7-th hybrid system, no evident interfacial electric field is formed at the Sb₂S₃/PTB7-th interfaces due to the weak interaction between the organic and inorganic components; however, the wavefunctions of electronic states near the CBM and VBM of the Sb₂S₃/PTB7-th interface mainly locates in the Sb₂S₃ and PTB7-th phases, respectively, resulting in the unique separate electronic states favorable to the dissociation and injection of photogenerated electron-hole pairs and free charge carriers at the Sb₂S₃/PTB7-th interface upon the excitation of Sb₂S₃ and PTB7-th, with the electrons remaining in or being injected into Sb₂S₃ crystal and the holes being injected into or remaining in PTB7-th phase. The solar cells with p–i–n architecture, consisting of Sb₂S₃/TiO₂-NHJ and PTB7-th, exhibits no negative E_ph effects on device performance, which is well accounted for by the ADCD model proposed on the basis of light intensity attenuation in the polymer phase and verified by the finite element simulations and C−V measurements. In contrast, in the p-(p + i)-n counterpart solar cell, the negative E_ph effects mainly manifest themselves to reduce the opposite built-in filed E_bi across the device, to retard the charge transport and to cause the serious recombination for a seriously deteriorated overall device performance (V_oc, J_sc, FF, and IPCE).

Our results provide a facile solution-processing method to prepare the high-quality Sb₂S₃/TiO₂ nanoarray heterojunction system with preferentially orientated, straightforward, and highly efficient charge transport channels for photovoltaic and other optoelectronic applications, as well as an understanding of the heterogeneously epitaxial growth of single-crystalline nanorods on the polycrystalline surface. Moreover, the NHJ material system and the p–i–n solar cell without the negative impact of E_ph on device performance offer the guides to the conceptual design of the efficient photoactive layer and device configuration in the solar cells based on nanorod/nanowire arrays, in particular, the devices configurated using a partner material with a greatly unbalanced charge transport property.

Methods

Sb₂S₃ precursor solution

The Sb₂S₃ precursor solution was prepared by a method similar to our previous report except for the addition of glycerin¹⁹. Briefly, 0.8 mmol of SbCl₃ (≥99.0%, Alfa Aesar) and 1.44 mmol of thiourea (≥99.0%, Alfa Aesar) were sequentially dissolved in 1 ml of N,N-dimethylformamide (DMF) under stirring at room temperature, resulting in a clear yellow precursor solution with 0.8 M SbCl₃ and an SbCl₃/thiourea molar ratio of 1:1.8, which is stable for weeks in ambient conditions. The different SG and NG solutions were derived from this precursor solution by DMF dilution. The concentrations of the SG (i.e., C_sg) and NG (i.e., C_ng) solutions were defined by the molecular concentration of SbCl₃ dissolved therein.

Preparation of Sb₂S₃/TiO₂-NHJ

The FTO glass substrates (25 × 25 mm², ca. 14 Ω sq⁻¹) patterned into stripes (20 × 4 mm²) in the middle part of the substrate by laser were commercially obtained from the Advanced Election Technology Co. Ltd., Liaoning Province, China. The FTO substrates were ultrasonically cleaned with acetone, isopropanol, and deionized water for 20 min, respectively, and then further cleaned with UV-Ozone for 15 min and washed with ethanol before use. A compact and polycrystalline TiO₂ nanoparticle film of ca. 110 nm in thickness was prepared on the FTO substrate as described elsewhere¹⁹.

Sb₂S₃/TiO₂-NHJ was prepared by a tiny-seed assisted RSCA method. Typically, the Sb₂S₃ precursor deposited onto the TiO₂ film by spin-coating (3000 rpm, 40 s) the SG solution (120 μl; normally C_sg = 0.1 M, unless otherwise specified) was sequentially annealed at 150 ^°C for 5 min and 300 ^°C for 10 min on a hotplate in an N₂-protected glovebox, resulting in tiny Sb₂S₃ seeds on the TiO₂ film. A transparent Sb₂S₃ precursor film was prepared on a tiny Sb₂S₃ seeds-decorated TiO₂ film by spin-coating (3000 rpm, 40 s) the NG solution (120 μl; normally C_ng = 0.8 M, unless otherwise specified) and thermally annealed by the route identical with that for the Sb₂S₃ seed formation, which was referred to as one RSCA cycle. After the seed-decorated TiO₂ film had been subjected to several RSCA cycle numbers (n), the Sb₂S₃/TiO₂-NHJ with a desirable length was in situ formed on the FTO substrate.

Solar cell fabrication

PTB7-th (M_w ~180,000; M_w/M_n ~2.5, 1-Material) was dissolved in chlorobenzene and stirred overnight at 60 °C; after the solution was naturally cooled down to room temperature, the PTB7-th solution (15 mg ml⁻¹) for spin-coating was obtained. To prepare p–i–n solar cells (Fig. 1), a PTB7-th film was spin-coated (2000 rpm, 60 s) onto the top of the Sb₂S₃/TiO₂-NHJ from the PTB7-th solution (70 μl) and annealed at 100 °C for 15 min to eliminate the solvent residue, where the spin-coating was immediately carried out once the PTB7-th solution was loaded. In order to prepare p-(p+i)-n solar cells (Fig. 4e), the PTB7-th solution droplet (70 μl) was first loaded onto the Sb₂S₃/TiO₂-NHJ and kept in the chlorobenzene atmosphere overnight for the infiltration of the polymer into the interspaces amongst the Sb₂S₃ nanorods; then, the sample was dried by rotating (2000 rpm, 60 s) and further annealed at 100 °C for 15 min to remove the solvent. Finally, an 8-nm-thick molybdenum oxide (MoO₃) film and a 100-nm-thick gold electrode were sequentially evaporated onto the PTB7-th layer under a pressure of 5.0 × 10⁻⁴ Pa. The optimized Sb₂S₃ nanorod array for efficient solar cells was prepared at C_sg = 0.1 M, C_ng = 0.8 M and n = 2.

Sample characterization and DFT calculation

The J−V curves of solar cells were measured with forwarding scan and dwell time of 50 s on a 94023 A Oriel^®Sol3A solar simulator (Newport, Inc.) under AM 1.5 G illumination (100 mW cm⁻²) at room temperature in air, and the light intensity from a 450 W xenon lamp was calibrated with a standard crystalline silicon solar cell. The J−V curves were collected with an Oriel® I-V test station (PVIV-1A, Keithley 2400 Source Meter, Labview 2009 SP1 GUI Software, Newport, Inc.). During the J−V measurements, the illumination area was limited by a mask with an aperture (3 × 3 mm²), which actually defined the effective area for getting the photovoltaic parameters of the device. The characterizations of film samples, other characterizations of solar cells, and the DFT calculation method are provided in Supplementary Methods.

Reporting summary

Further information on research design is available in the Nature Research Reporting Summary linked to this article.