Inorganic Ions Assisted Anisotropic Growth of CsPbCl3 Nanowires

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Inorganic Ions Assisted Anisotropic Growth of CsPbCl3 Nanowires with Surface Passivation Effect Yingying Tang, Xianyi Cao, Alireza Honarfar, Mohamed Abdellah, Chaoyu Chen, José Avila, Maria C. Asensio, Leif Hammarström, Jacinto Sa, Sophie E. Canton, Kaibo Zheng, Tõnu Pullerits, and Qijin Chi ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b09113 • Publication Date (Web): 08 Aug 2018 Downloaded from http://pubs.acs.org on August 9, 2018

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ACS Applied Materials & Interfaces

Inorganic Ions Assisted Anisotropic Growth of CsPbCl3 Nanowires with Surface Passivation Effect Yingying Tang,a Xianyi Cao,a Alireza Honarfar,b Mohamed Abdellah,c,d Chaoyu Chen,e José Avila,e Maria-Carmen Asensio,e Leif Hammarström,c Jacinto Sa,c Sophie E Canton,f,g Kaibo Zheng,*a,b Tõnu Pullerits,b and Qijin Chi*a a. Department of Chemistry, Technical University of Denmark, DK-2800 Kongens Lyngby, Denmark b. Department of Chemical Physics and NanoLund, Lund University, Box 124, 22100, Lund, Sweden c. Ångstrom Laboratory, Department of Chemistry, Uppsala University, Box 523, 75120 Uppsala, Sweden d. Department of Chemistry, Qena Faculty of Science, South Valley University, 83523 Qena, Egypt e. Synchrotron SOLEIL, L’Orme des Mérisiers, Saint Aubin-BP 48, 91192 Gif sur Yvette Cedex, France f. ELI-ALPS, ELI-HU Non-Profit Ltd., Dugonicster 13, Szeged 6720, Hungary g. Attoscience Group, Deutsche Elektronen Synchrotron (DESY), Notkestrasse 85, D-22607 Hamburg, Germany

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ABSTRACT. All-inorganic halide perovskite nanowires (NWs) exhibit improved thermal and hydrolysis stability and could thus play a vital role in nanoscale optoelectronics. Among them, blue-light based devices are extremely limited due to the lack of facile method to obtain high-purity CsPbCl3 NWs. Herein, we report a direct and facile method for the synthesis of CsPbCl3 NWs assisted by inorganic ions that served both as a morphology controlling agent for the anisotropic growth of nanomaterials and a surface passivation species modulating the surface of nanomaterials. This new approach allows us to obtain high-purity and size-uniform NWs as long as 500 nm in length and 20 nm in diameter with high reproducibility. XPS and ultrafast spectroscopic measurements confirmed that a reduced bandgap caused by the surface species of NWs relative to nanocubes (NCs) was achieved at the photon energy of 160 eV, due to the hybrid surface passivation contributed by adsorbed inorganic ions. The resulting NWs demonstrate significantly enhanced photoelectrochemical performances, 3.5-fold increase in the photocurrent generation, and notably improved stability compared to their nanocube counterparts. Our results suggest that the newly designed NWs could be a promising material for the development of nanoscale optoelectronic devices.

KEYWORDS:

inorganic

perovskite,

CsPbCl3

nanowires,

surface

passivation,

photoelectrochemical cell, electron and hole injection, ultrafast spectroscopy


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INTRODUCTION Metal halide perovskites (MHPs) with unique physicochemical properties have emerged at the forefront of light absorber materials, because of their great potential in design and fabrication of new-generation optoelectronic devices, including solar cells,1-5 photodetectors,6-9 light-emitting devices,10 field-effect transistors11 and lasers.12 Tunable wavelength, large diffusion length and long carrier lifetime represent their most striking characteristics. These properties are, however, morphology and size dependent. To date, researchers have successfully engineered MHPs in the forms of nanocubes (NCs)13 and nanoplatelets (NPLs)14 with sizes from several to hundreds of nanometers. Perovskite NWs with a high length-to-width ratio and well-defined size and morphology are highly desirable for both fundamental studies and their practical applications, but low-cost and facile synthesis has rarely been successful. The recently developed synthesis methods for NWs include vapour-phase method, hotinjection method, and room-temperature processes,15-18 but they have suffered from some disadvantages, for instance, restricted reaction conditions such as temperature, costly reaction process, large amounts of defects and low-purity products. In fact, these are essential for the development of optoelectronic devices and therefore limit their wide applications. In general, a facile and efficient synthesis of perovskite NWs remains a challenge. In particular, the blue-light devices are still in their infancy due to the lack of a facile method for the synthesis of high purity Cl-based perovskite NWs. So far, several synthetic routes have been attempted to prepare high purity CsPbCl3 NWs. However, unlike the synthesis of CsPbCl3 NCs the direct synthesis method can only produce a small portion of CsPbCl3 NWs in the product mixture that contains NCs, nanosheets (NSs) and NWs.19 To overcome this challenge, a few attempts have been made by adding organic agents such as tri-n-octylphosphine (TOP) or trioctylphosphine oxide (TOPO) in 3 ACS Paragon Plus Environment


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order to improve the selectivity of the synthesis,20 but only mixed products dominated by CsPbCl3 NCs were obtained. An alternative route involves a halide-exchange procedure using CsPbBr3 NWs as the precursors.21,22 This method is complicated, time-consuming and costly. Therefore, a more efficient way is thus needed and pursued in the present work. To this end, previous studies on the synthesis of metallic nanostructures offer the key clue that inorganic ions could have a notable impact on the formation of nanoscale crystals. For example, Sun et al. synthesized Au nanostructures with controlled morphology of rod, cuboid and decahedron through regulating the concentration of Cu2+.23 Inspired by these studies, we have tested the impacts of inorganic ions on the formation of MHP nanostructures during solution-processed synthesis. In addition, it is noticed that hybrid passivation from both organic and inorganic ligands have been shown as an effective way to passivate trap states, improve conductivity and enhance carrier transport for the nanomaterials.24, 25 In the present work, we have demonstrated the development of a facile approach to the direct synthesis of 20 nm (in diameter) and 500 nm (in length) single-crystalline and high-purity CsPbCl3 NWs with a product yield of 95%. This is achieved by the direct hot-injection method, facile and reproducible, in the presence of specific inorganic ions. We find that the Cu2+ ions are among the most efficient species and play two complementary roles. One is to act as a morphology controlling agent by modulating the nanocrystal growth; the other is to passivate the perovskite NWs via surface adsorption. The latter modulates the valence band maximum (VBM) of NWs compared to NCs, further facilitating the efficient hole transfer from perovskite to the electrolyte. As a result, we have achieved a significant enhancement in photocurrent generation and notably improved stability at ambient conditions. These NWs might be a promising candidate material for fabrication of relevant optoelectronic devices.


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MATERIALS AND METHODS Chemicals and Materials. All the chemical reagents were at least of analytical grade and used as received without further purification. Lead (II) chloride (PbCl2, 99%), copper (II) chloride dihydrate (CuCl2·2H2O, 99%), cesium carbonate (Cs2CO3, 99%), 1-octadecene (ODE, technical grade 90%), oleic acid (OA, technical grade 90%), oleylamine (OAm, 90%), octylamine (OCT, 99%), titanium (IV) isopropoxide (97%), acetic acid and lithium perchlorate (99.99%) were all purchased from Sigma Aldrich. Hexane was obtained from Honeywell. Ethanol and HNO3 was obtained from Fluka. For simplicity, we use CuCl2 instead of CuCl2.2H2O in the rest text. Preparation of CsPbCl3 nanostructures. CsPbCl3 NCs was synthesized according to a modified synthetic approach previously reported.21 In brief, 0.2 g Cs2CO3 and 0.6 mL OA were loaded into a 3-neck flask along with 7.5 mL ODE, degassed at 120 °C for 20 min, and then heated under Ar flow to 150 °C until all Cs2CO3 reacted with OA. After that, 0.2 mmol of PbCl2 was first added into a 50 mL three-neck flask containing 5 mL ODE and 1 mL OA, degassing under Ar flow at room temperature for 20 min and then heated at 120 °C under Ar flow with constant stirring for 20 min. Thereafter, 0.8 mL OCT and 0.8 mL OAm were successively injected at 120 °C under Ar. The temperature was then elevated to 160 °C, and the solution was further stirred for 20 min. The solution was kept at 160 °C, and 0.7 mL as-prepared Cs-oleate was injected quickly. After 2h reaction, the mixture was cooled in an ice-water bath. CsPbCl3 products were collected by centrifugation and re-dispersed in hexane for further use. For the preparation of perovskites in the presence of inorganic salts, a mixture of 0.2 mmol of PbCl2 and CuCl2 was used, into which certain amount of CuCl2 was added (molar ratio). Synthesis of TiO2 nanoparticles. The anatase-TiO2 nanoparticles were synthesized by a hydrothermal method. 11.72 g of titanium (IV) iso-propoxide was mixed with 2.4 g of acetic acid 5 ACS Paragon Plus Environment


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and stirred at room temperature for 15 min. This was followed by adding 58 ml deionized water and stirring for 60 min with a magnetic stirrer at 200 rpm. After that, 0.8 ml concentrated nitric acid (HNO3, 70%) was added to the solution, which was heated to 80 °C and stirred for another 2 h. 75 ml deionized water was then added, and the solution was transferred to Teflon cups in an autoclave held at 180 °C for 17 hours. After cooling to room temperature, 0.48 ml HNO3 was added, followed by centrifugation at 7000 rpm and rinsing with water three times. Centrifugation and rinsing were repeated twice more before a final rinse in the absolute ethanol to obtain a final dispersion concentration of 13.6 mg/ml in ethanol. Materials characterization. TEM imaging was conducted on Tecnai G2 T20 TEM from FEI Company. The grids made from Cu were used as support materials for TEM imaging. Crosssection SEM and EDX characterizations were carried out on an FEI Quanta FEG 200 ESEM. Al foils and PELCO TablsTM carbon conductive tabs were used as substrates in SEM tests. Samples were tested with gold coating. A 5500 AFM system (Agilent Technologies) was used for all AFM imaging. The AFM scanning size error was measured on TGX calibration gratings. The phase and purity of all samples were also characterized by XRD with Cu Kα1 (λ = 1.5406 Å) radiation. XPS (Thermo Scientific) was performed to analyze the compositions of samples, with Al-Kα (1486 eV) as the excitation X-ray source. The pressure of the analysis chamber maintained at 2 × 10-10 mbar during measurements. All characterizations were carried out at room temperature. The absorption behaviors of the samples were studied by the UV-vis spectrometer from Agilent Technologies (Santa Clara, USA). Absorption (α/S) data were converted from diffuse reflectance spectra using the Kubelka−Munk function, α/S = (1 − R)2/2R, where R is the reflectance coefficient, and α, S are the absorption and scattering coefficient, respectively.26 Photoluminescence (PL) spectra were scanned on a PL spectrometer (FLS980,


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Edinburgh Instruments Ltd.) from 380 to 900 nm with 1 nm increments and 1 s integration time, under an excitation wavelength of 360 nm. Fabrication of photoelectrodes. Prior to use, ITO substrates were cleaned by sonicating sequentially in deionized water, alcohol and acetone, each for 10 min. The TiO2 nanoparticles layer was deposited by drop-casting onto the clean ITO electrodes at a rotation rate of 600 rpm, dried under 70 °C for 1 h. The perovskite film was deposited by drop-casting method onto TiO2/ITO (600 rpm) and dried under 70 °C for 15 min. Photoelectrochemical measurements. All the photoelectrochemical (PEC) characterizations were performed on the EC-lab workstation in a 3-electrode configuration with the assembled photoelectrodes (CsPbCl3 NCs or NWs on TiO2/ITO glass) as the working electrode, the Pt wire as the counter electrode and the saturated calomel electrode (SCE) as the reference electrode. The ethyl acetate solution containing 0.5 M lithium perchlorate was used as electrolyte throughout, which was saturated with Ar gas before measurements. Photocurrent stability tests were carried out by measuring the photocurrent under light illumination (33 mW/cm2, Ocean Optics). Ultrafast transient visible absorption spectroscopy. Transient absorption (TA) experiments were performed by using a femtosecond pump probe setup. Laser pulses (800 nm, 80 fs pulse length, 1 kHz repetition rate) were generated by a regenerative amplifier (Spitfire XP Pro) seeded by a femtosecond oscillator (Mai Tai SP, both Spectra Physics). The pump pulses at 350 nm were generated by an optical parametric amplifier (Topas, Light Conversion). The used excitation photon fluxes are 3 × 1012 photons/cm2/pulse. For the probe, we used the supercontinuum generation from a thin CaF2 plate. The mutual polarization between pump and probe beams was set to the magic angle (54.7°) by placing a Berek compensator in the pump



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beam. The probe pulse and the reference pulse were dispersed in a spectrograph and detected by a diode array (Pascher Instruments). Ultrafast transient mid-IR absorption spectroscopy. The 1 mJ, 45 fs output of a 1 kHz Ti:Sapphire amplifier (Spitfire Pro, Spectra-Physics) was split into two separate commercial optical parametric amplifiers (TOPAS-C, Light Conversion), which generate the visible pump 410 nm and the mid-IR probe (1850 – 2200 cm-1) pulses. Prior to reaching the sample, the probe beam was split into equal intensity probe and reference beams using a wedged ZnSe window. Both beams pass through the sample, but only the probe beam interacts with the photoexcited volume of the sample. All beams are focused with a single f = 10 cm off axis parabolic mirror to a ~70 µm spot size in the sample. The pump intensity changed on the sample via density filter. The probe and reference beams were dispersed by a commercial monochromator (Triax 190, HORIBA Jobin Yvon) equipped with a 75 groove/mm grating and detected on a dual array, 2 × 64 pixel mercury cadmium telluride detector (InfraRed Associated, Inc). The instrument response function for the experiments was approximately 100 fs. The sample was mounted in a Harrick flow cell. The setup had been used before in our previous work.27 X-ray photoelectron spectroscopy. X-ray photoelectron spectroscopy (XPS) were performed at beamline Antares at Soleil Synchrotron Laboratory. The photon energies used for XPS were 700 eV and 160 eV, with normal emission geometry. The energy scale of the spectra was calibrated relative to the Fermi level of the gold foil that was used as the electrical contact with the sample. Multiple positions on the samples were measured to prove the homogeneity of the films. No charging or beam-induced sample decomposition effects were detected in the spectra.

RESULTS AND DISCUSSION 8 ACS Paragon Plus Environment

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To reveal the structural characteristics of the synthesized materials, we performed systematic characterization using a range of microscopy and spectroscopy techniques. The phase and purity of the as-synthesized perovskites were characterized by powder X-ray diffraction (XRD). Figure 1a compares the XRD patterns of CsPbCl3 NWs and NCs. Their crystal structure is similar with both adopting the tetragonal phase (space group P4mm, No. 99).28 The finger-printing diffraction peaks at 15.8°, 22.4°, 31.8°, 31.9°, 35.6°, 39.4°, 45.7° and 56.9° are indexed to the (010), (011), (002), (020), (012), (121), (022) and (222) crystal planes, respectively. The pattern is consistent with those of related standard samples with a collection code (109294-ICSD) in the Inorganic Crystal Structure Database (ICSD) (Figure S1a). These sharp peaks in the XRD patterns indicate the high-crystalline nature of these perovskite nanostructures. The peak splitting observed at 31.8° and 31.9° (Figure S1a and S1b) confirms the lower symmetry tetragonal phase rather than the cubic phase. Moreover, no distinct Bragg diffraction peaks at 2θ = 13°, 25.1°, 39.9° for the orthorhombic phase are detected, further supporting that the present CsPbCl3 nanomaterials adopt the tetragonal crystalline phase. The morphology and size distribution of the CsPbCl3 nanomaterials were analyzed by transmission electron microscopy (TEM) and scanning TEM (STEM). As imaged by TEM (Figure 1b), the CsPbCl3 NCs have a cubic shape with an average size of ~15 nm. The presence of CuCl2 resulted in the morphology change from NCs to NWs with the length of ~ 500 nm and the diameter of ~ 20 nm. Figure 1c displays the TEM image of the CsPbCl3 NWs in a wide range (also see Figure S2a and Figure S3h), indicating a high purity of the NWs with a yield of 95 %. Besides, the TEM images of other batches of NWs are displayed in Figure S2, indicating the reproducibility of this method. As shown in the upper inset of Figure 1d, the lattice fringe with an interplanar spacing of 4.0 Å coincides with the (011) crystal plane of tetragonal crystal



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Figure 1. Structural characteristics of as-synthesized perovskite nanomaterials. (a) XRD patterns of CsPbCl3 NWs, CsPbCl3 NCs and reference sample of orthorhombic (ICSD, #109294); (b) TEM image of CsPbCl3 NCs; (c and d) TEM images of CsPbCl3 NWs; (e) STEM image of CsPbCl3 NWs; (f) AFM image of CsPbCl3 NWs; (g) Cross-sectional height profile of single NWs. Insets in (d) are a HRTEM image (upper) and a FFT image (lower).


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structures, showing the single-crystalline nature of these NWs. This is further supported by the FFT patterns that exhibit the characteristics of (011) and (001) crystal planes (the lower inset in Figure 1d). These observations suggest that the growth of NWs is directed by the (011) plane, consistent with the previous report.29 Furthermore, the strong intensity of the (011) diffraction peaks pinpoints the anisotropic growth of CsPbCl3 NWs. A detailed observation of the NWs was performed by using STEM. The STEM image confirms that uniform NWs are obtained (Figure 1e). In order to reveal the detailed morphology of these perovskite NWs, we performed atomic force microscopy (AFM) measurements. The AFM image shows that the NWs have a smooth surface (Figure 1f). The diameter of treated CsPbCl3 NWs is around 15 nm, which is consistent with the TEM and STEM results. Meanwhile, energy-dispersive X-ray diffraction spectroscopy (EDX) mapping by SEM and STEM was conducted to confirm the chemical compositions and the elemental distribution (Figure S3 and S4). All the key elements, Cs, Pb, Cu, and Cl are detected and uniformly distributed throughout the NWs. The X-ray photoelectron spectroscopy (XPS) analysis indicates that Cu2+ ions are adsorbed on the surface of the NWs (Figure S5 and Table S1). In the absence of Cu2+, NCs are formed from unit cells (Scheme 1a) predominantly driven by the quasi-uniform adsorption of oleylammonium cations (RNH3+) and oleic acid anions (RCOO-) on each of the six surface sides (Scheme 1b). From the structural point of view, CsPbCl3 NCs can be viewed as an extension of crystal units in a three-dimensional fashion (Scheme 1c). As previously reported, however, Cu2+ was evidenced to play a key role in the formation of hexoctahedral Au-Pd alloy nanocrystals.23 In the presence of Cu2+ ions, nanocrystals thus tend to



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grow anisotropically along the [011] direction. According to thermodynamics, high-energy surfaces favor faster growth rate than low energy surfaces. Therefore, the (011) crystal plane

Scheme 1. Schematic illustration of nanostructure formation. (a) unit cell of CsPbCl3 (tetragonal phase); (b) crystallization of the CsPbCl3 NCs; (c) structure of the CsPbCl3 NCs extended in three-dimension; (d) adsorption of Cu2+ on the seeds of the CsPbCl3 NWs; (e) formation of the NWs along the [011] direction; (f) structure of the CsPbCl3 NWs extended in one-dimension. with high-energy surfaces selectively adsorb ions from the reaction solution in order to reduce its surface energy and block the growth (Scheme 1d). It is known that RNH3+ ions behave like Cs+ and compete with Cs+ to suppress the NWs growth.30 Meanwhile, the Cu2+ ions also compete with Cs+ to be adsorbed from the reaction solution onto the surfaces of the nanocrystals. Due to abundant negative charges on the surface of the nanocrystals, the positively charged ions such as RNH3+ and Cu2+ are predominantly adsorbed. This is confirmed from the crystallographic view, in which the high density of Cl- ions are exposed on the surface (Figure S6). In addition, the Cu2+ ions also compete with RNH3+ to occupy more active surface sites due to their stronger electronegativity. Since a single Cu2+ ion can bind to two active sites, Cu2+ ions are more


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efficient than RNH3+ in decreasing the surface energy. Surface sites occupied by RNH3+ are thus too few to suppress the growth of NWs. The energy barrier for formation of these particular MHPs nanostructures is ultralow, and the formation reaction is thus very fast. Compared to Cu-based perovskites, Pb-based perovskites are more stable. For example, CsCuCl3 was chosen as a chemo-dosimeter for the selective sensing of Pb2+ due to its fast conversion into CsPbCl3.31 Pb2+ behaves similarly to Cu2+ in either charge or bonding configuration, so that Cu2+ ions on the surface become nucleation sites and are replaced by Pb2+ ions, thereby facilitating the NWs growth along the [011] direction. In order to ascertain the controllable-morphology ability of CuCl2, control experiments were conducted, in which much lower (3 %) and much larger (10 %) amount of CuCl2 were added. The results show that only the mixture of NCs (3 %, CuCl2) and NWs (10 %, CuCl2) was obtained (Figure S7), in contrast to the sole formation of pure NWs when 5 % CuCl2 was used in the synthesis. Therefore, controlling the precise amount of CuCl2 is critical for the formation of size-uniform and highpurity NWs. In addition, we also tried to explore this method for the synthesis of CsPbBr3 NWs, but only nanosheets were obtained instead of NWs (Figure S7c and 7d). To analyze the optical properties of newly prepared NCs and NWs, their colloidal solutions were characterized by UV-vis absorption and photoluminescence (PL) emission spectroscopies. As shown in Figure 2a, untreated CsPbCl3 NCs show the first absorption onset at ~405 nm with a bandgap of 2.92 eV (estimated from the Tauc’s plots, Figure S8). The absorption of the NWs starts at ~408 nm with a bandgap of 2.80 eV (Figure 2b). The UV-vis spectroscopic analysis thus indicate that the absorption onset and the optical bandgap are only slightly affected by the presence of inorganic ions. Compared to the NCs, however, the emission peak for the NWs observed at ~420 nm is red-shifted by ca 15 nm. The red-shift might be attributed to the quantum



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confinement effect, because larger sized crystalline domains are observed in NWs by TEM compared to NCs. In addition, the full width at half-maximum (FWHM) is narrowed from 15 nm for the NCs to 12 nm for the NWs, which suggests NWs having fewer surface defects.32

Figure 2. UV-vis and photoluminescence spectroscopic analysis. UV-vis absorption spectra and photoluminescence spectra of CsPbCl3 NCs (a) and NWs (b) excited at 360 nm, respectively.

To perform photoelectrochemical (PEC) measurements (Figure 3a), we fabricated the photoactive electrodes by depositing the TiO2 and perovskite films on indium tin oxide (ITO). The SEM image reveals that both TiO2 and perovskite films are uniformly distributed on the ITO surface and the morphology of NWs are well preserved after annealing (Figure 3b and Figure S9). The thickness of each layer measured by cross-sectional SEM images is ~45 nm for the perovskite layer and ~420 nm for the TiO2 layer, respectively. Figure 3c shows the photocurrent responses from the CsPbCl3 NCs and NWs based photoelectrodes. The amperometric i-t curves demonstrated that the electrodes can repeatedly produce stable photocurrent signals in response to periodical light ON and OFF. The photocurrent of the CsPbCl3 NWs (1.1 µA/cm2) displays a 3.5-fold enhancement compared to that of the NCs (0.3 µA/cm2). In control experiments, the


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photocurrent measured at blank ITO was negligible. Increasing the thickness of perovskite films displayed negligible effect on the photocurrent (Figure S10a). Meanwhile, a sharp current spike

Figure 3. Photoelectrochemical performances of the prepared nanomaterials. (a) Schematic representation of the PEC cell setup (electrolyte: 0.5 M LiClO4 in ethyl acetate); (b) Crosssectional SEM micrograph of the photoelectrode; (c) Temporal photoresponse of CsPbCl3 NWs, CsPbCl3 NCs and blank ITO under 30 mW/cm2 AM 1.5 G illumination; (d) Stability of the photocurrent generated for CsPbCl3 NWs and NCs under ambient conditions at room temperature. is induced by the pyro-phototronic effect after increasing the thickness of the perovskite layer.33 The thickness of perovskite layers has been an optimal one in our system. The recorded J-V characteristic for the above photoelectrochemical cell was shown in Figure S10b. A comparison



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of numerous NWs-based optoelectronic properties is summarized in Table S2. The present NWs show enhanced photocurrents compared to the reported CsPbCl3 NCs, NPLs, and quantum dots (QDs). Moreover, our NWs also show comparable photocurrent generation with most of Br and I based perovskite NWs. From the enlarged part of the typical on/off single cycle, it is observed that the rise and decay time of the CsPbCl3 NWs is significantly reduced to 0.88 and 1.3 s (Figure S11), respectively. In contrast, the NCs show much slower responsive time upon their exposure to light illumination, especially with the rise time of ~1.6 s. Furthermore, the photocurrent of the CsPbCl3 NWs is largely constant, while the photocurrent of the CsPbCl3 NCs keeps decreasing from 0.4 to 0.3 µA/cm2. In addition, we measured the stability of the CsPbCl3 NWs and NCs in ambient environment at room temperature (Figure 3d). The photocurrent generated by the CsPbCl3 NCs decayed rapidly, but kept constant for the NWs.34 The results indicate superior moisture stability of the NWs over the NCs. This can be explained partially by the hybrid passivation effects of Cu2+ and organic ions on the NWs, which is known to be an effective way for improving the stability of MHPs. In addition, different crystal faces and density of organic ligands in NWs and NCs could also contribute to the stability difference for two nanomaterials.35 In order to rationalize the difference in PEC performances exhibited by the two types of samples (i.e. CsPbCl3 NCs and NWs), we conducted several photophysical studies using various spectroscopic techniques. In a PEC cell with perovskite sensitized TiO2 as photoanode, the photogenerated electron-hole pairs in the perovskites need to be separated before recombination to generate photocurrents. This involves electron-injection to TiO2 and hole-injection to the electrolyte. We first monitored the electron injection process by ultrafast transient absorption (TA) spectroscopy probed in both the visible (400 - 800 nm) and the mid-IR (4000 - 6000 nm) 16 ACS Paragon Plus Environment

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regions. The former is sensitive to the inter-band optical transition of the photogenerated charges, while the latter can characterize the free carriers in the conduction band of semiconductors. TA spectra of the NCs- and NWs-TiO2 films after 400 nm fs-pulse excitation (delay time = 0.5 ps) are shown in Figure 4a. The negative ground state bleach for the NCs and NWs around 420 nm and the two excited state absorption bands on both the red and blue sides of the bleach bands are observed clearly, providing the evidence for population of band edge excited states after thermalization of the photogenerated hot charges (< 0.5 ps).36 Meanwhile, no other trap-related red shift bleaching band was found in both NWs and NCs, indicating that there is no trap-related emission in our samples.37 After 2 ps, a featureless absorption increasing monotonously with the wavelength is observed in both the NCs-TiO2 and NWs-TiO2 samples (Figure 4b). This is a signature indication of intra-band transition of the free charges at the conduction band of semiconductors or of states to the conduction band.27 In order to confirm the optical transition of trapped electrons from shallow trap electron injection from nanocrystals to TiO2, we investigated the TA kinetics in both the visible and mid-IR probe ranges. Figure 4c and 4d show the comparison of the TA-vis kinetics at the maximum bleach of neat NCs and NWs attached to TiO2. When the nanoparticles are attached to TiO2, the decay of the bleach becomes much faster than that for the neat particles, indicating that new fast excited state depopulation pathways are opened. We can obtain the kinetics of such depopulation processes by normalizing the TA kinetics at a long time decay and extracting the differential curves (the insets in Figure 4c and 4d). Moreover, if we mirror this differential decay trace and overlap it with the TA-IR spectrum at early time scale (red curve in Figure 4e), we can see they undergo almost the same time evolution (Figure 4e). This suggests that the depopulation of the excited charges in nanoparticles



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and the population of charges in TiO2 happened simultaneously. In other words, this is a clear fingerprint indication of electron injection from nanoparticle to TiO2. In addition, we note that the photoexcitation of neat nanoparticles could also induce excited state absorption in TA-IR.

Figure 4. Transient optical spectra of CsPbCl3 NCs and NWs. (a) TA spectra for NCs-TiO2 and NWs-TiO2 at a 0.5 ps time delay after 400 nm fs-pulse excitation; (b) Transient infrared spectra for NC-TiO2 and NW-TiO2 at a time delay of 2 ps excited at 400 nm; (c and d) Normalized bleaching kinetics of NCs-TiO2, NCs, NWs and NWs-TiO2 probed at 400 nm excitation. The insets show the differential curves of bleaching kinetics between NCs-TiO2 (NWs-TiO2) and NCs (NWs); (e) TA kinetics in the visible and mid-IR range of NCs-TiO2 and NCs, the black curve is mirrored from the inset of (c); (f) Schematic illustration of photoinduced charge generation and transfer process in NCs and NWs. This is because some of photogenerated excitons would quickly dissociate into free charges to reach the equilibrium condition according to the Saha-Langmuir theory.38 However, the TA-IR kinetics in this case shows only an instantaneous instrument limited rise after excitation, which is 18 ACS Paragon Plus Environment

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much faster than that in the NCs-TiO2 sample. This further confirms the population of charges in TiO2 after electron injection process. Moreover, the electron injection is faster for the NCs than for the NWs, as evaluated by the lifetime fitting of the differential decay traces in Figure 4c and d, where the injection in the NCs takes 0.4 ps while in the NWs, two components with lifetimes of 0.4 ps and 8 ps are found. They can be attributed to the two-step electron injection involving an intermediate charge transfer state (CTS), which was previously identified in QD-sensitized ZnO.39 Here, the electrons are initially transferred to the CTS remaining bound to the residual holes and CTS dissociates afterwards with slower rates to release the electrons as injected free charges. The CTS formation depends on the Coulombic interaction between the injected charges in the acceptors and residual charges in the perovskites for this system. Such Coulombic interaction in QD-acceptor systems is determined by the dielectric constant and dimension of the QDs. In this scenario, one possible reason could be the effective dielectric constant for NWs with CuCl2 as a shell is smaller than NCs, resulting in larger Coulombic energy. In addition, the dissociation of the CTS states is also dominated by the exciton binding energy of the nanoparticles. As observed in the absorption spectra of both NCs and NWs, more pronounced excitonic peak is seen in NWs, indicating larger exciton binding energy. This could induce the slow dissociation of the CTS states, which will enhance the probability for the occurrence of the CTS in NWs. Although the electron-injection pathways for the two samples are different (Figure 4f), the overall injection processes are all much faster than the intrinsic recombination of the band edge charges as indicated in the TA kinetics of neat NCs and NWs with both average lifetimes of the free charges longer than 1 ns. Consequently, most of the photogenerated electrons in nanoparticles should be effectively injected into the TiO2 film as long as they are



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directly attached to the surface. In such a scenario, therefore, it is unlikely that the difference in the PEC performances of NCs and NWs is caused by the variation of electron-injection rates. Based on the above analyses, we turn our attention to the hole-injection dynamics into ClO4- in the electrolyte.40 Upon illumination, electron-hole pairs are produced. Electrons transport to the Pt counter electrode, where Li+ cations are reduced. ClO4- ions will recombine with holes and are oxidized to ClO4. Unfortunately, the light scattering at the liquid/solid interface restricts direct spectroscopic measurements if the electrolyte layer is present. In addition, the large effective mass of the hole as well as the closely-lying energy states packed at the valence band of nanoparticles makes the ground state bleach in TA insensitive to the excited hole population.41 As an alternative, the difference between the energy positions of the intrinsic valence bands of the two samples elucidates the possible origin for their different hole injection capabilities. Figure 5a shows the X-ray photoelectron VB spectra of NWs and NCs acquired at two different X-ray photon energies (PE) (i.e. 700 eV and 160 eV). We obtained the VBM position from the intersection of the band edge slope to zero. At the photon energy of 700 eV, the NCs and NWs have the same VBM of 2.85 eV below the Fermi level, Ef. When the PE is reduced to 160 eV, however, a large blue shift of the VBM is observed for both samples. The VBM of NCs (5.81 eV vs. Ef) is significantly lower than that of NWs (5.28 eV vs. Ef) by 0.53 eV. According to DFT calculations,42 the VB of CsPbCl3 perovskite is mainly constituted by Cl 3p and Pb 6s orbitals in which the former is predominant. When Cu2+ replaces Pb2+, Cu-Cl bonds can also be formed. We first note that the photoionization cross-sections of atomic Cl 3p, Pb 6s and Cu 3d orbitals keep a similar ratio (5 : 1 : 50) for both PE (Figure S12).43 This means that the possible factor of variable elemental sensitivity through different photoionization cross-section can be excluded in first approximation. Another possibility arises from the different escape depth (DE) of the 20 ACS Paragon Plus Environment

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photoelectrons for the two cases of PE. For the case of PE =700 eV, DE is larger than 1.5 nm; while it is smaller than 0.5 nm for PE =160 eV.44 Given the fact that the lattice constant of

Figure 5. (a) XPS spectra at incident photon energy of 700 eV (upper) and 160 eV (lower), respectively; (b) Schematic energy level diagram of valence band in NCs (upper) and NWs (lower).

CsPbCl3 in the cubic phase is 5.6 Å, only the very first lattice layer of the surface can be probed with PE = 160 eV.45 In this scenario, the different VBM obtained in Figure 5a can be directly interpreted as an increased energy bandgap (Eg) at the surface components compared to the bulk for both samples. Based on the above discussion, we could clearly confirm the existence of Cu2+ ions on the surface rather than in the intrinsic body of the NWs due to the different bandgap of surface species. This is illustrated in Figure 5b. The origin for such distinct surface band structures could rest on the two factors: 1) Halide salts PbCl2 with a typical bandgap (Eg) of 4.73 eV are located preferentially at the surface in the absence of Cu2+ cations and formed the type-I 21 ACS Paragon Plus Environment


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core-shell structure in the present case. In the existence of Cu2+, however, Cu2+ ions can partially replace the Pb2+ ions in the NWs to decrease the surface bandgap since CuCl2 has a bandgap of 3.67 eV (Figure S13). 2) Smaller crystalline domains causing stronger quantum confinement are formed on the surface of the NCs leading to a larger bandgap. Such process has much less impact in the NWs due to lower surface energy and intercalation of the Cu2+. Although further characterization is desirable in ongoing studies especially to quantify various contributions, we can nevertheless conclude that the notable reduction in the surface Eg of the NWs weakens the hole-injection energy barrier from the inside NWs to the electrolyte (Figure 5b). We believe this finding is the main factor responsible for the enhancement of photocurrents observed in the NWs-based PEC cells.

CONCLUSIONS In conclusion, a facile solution-phase processing strategy enables direct synthesis of highcrystalline, uniform CsPbCl3 NWs via anisotropic growth along the [011] direction assisted by inorganic ions. PEC measurements have verified that CsPbCl3 NWs enable a 3.5-fold enhancement in photocurrent generation. This is due to efficient injection of holes to the electrolyte induced by the adsorption of Cu2+ on the surface of NWs, which is evidenced by a reduced bandgap in the surface species of NWs revealed by XPS measurements. Furthermore, the NWs exhibit notably improved stability at ambient conditions, which benefits their application as a promising photonic material in nanoscale optoelectronic devices. This study makes an important step toward the direct synthesis of single-crystalline and high-yield CsPbCl3 NWs via simple and low-cost solution-phase route. This newly developed method could be


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useful as well for exploiting other types of inorganic ions in order to control the morphology and optoelectronic properties of metal halide perovskite nanostructures. ASSOCIATED CONTENT The Supporting Information is available free of charge on the ACS Publications website. Powder XRD patterns, TEM and SEM images, XPS spectra, photoionization cross-sections, Tauc plots, current-time curves. AUTHOR INFORMATION Corresponding Author [email protected] and [email protected] Notes The authors declare no competing financial interests. ACKNOWLEDGMENT This work was supported by the Villum Foundation, Independent Research Fund DenmarkNature Sciences (DFF-FNU, Project No DFF-7014-00302), Independent Research Fund Denmark-Sapere Aude starting grant (No. 7026-00037A) and Swedish Research Council VR starting grant (No. 2017-05337), the Helmoltz Recognition Award, the ELI-ALPS project (GINOP-2.3.6-15-2015-00001) and the Chinese Scholarship Council for the PhD scholarship to X.C. (No. 201406170040). We acknowledge SOLEIL for provision of synchrotron radiation facilities. The authors thank the beam-line "Antares" program (No. 20171424). We thank Dr. Shiyu Gan for assistance to the photoelectrochemical measurements and Dr. Hongyu Sun for HAADF STEM imaging. REFERENCES 23 ACS Paragon Plus Environment


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Inorganic Ions Assisted Anisotropic Growth of CsPbCl3 Nanowires

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