Niobium doped TiO2 nanorod arrays as efficient electron transport materials in photovoltaic

Journal of Power Sources 450 (2020) 227715

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Niobium doped TiO2 nanorod arrays as efficient electron transport materials in photovoltaic Peng Zhong a, b, *, Xinpeng Chen a, Bingqiang Niu a, Cong Li a, Yucheng Wang d, He Xi a, b, Yimin Lei a, b, Zhenni Wang a, Xiaohua Ma b, c, ** a

School of Advanced Materials and Nanotechnology, Xidian University, 266 Xinglong Section of Xifeng Road, Xi’an, 710126, Shaanxi, People’s Republic of China Key Lab of Wide Band-Gap Semiconductor Materials and Devices, Xidian University, Xi’an, 710071, Shaanxi, People’s Republic of China School of Microelectronics, Xidian University, Xi’an, 710071, Shaanxi, People’s Republic of China d School of Physics and Optoelectronic Engineering, Xidian University, Xi’an, 710071, Shaanxi, People’s Republic of China b c

H I G H L I G H T S

G R A P H I C A L A B S T R A C T

� Ln is substantially improved from ~14 μm to ~50 μm after 1mol% Nb doping. � Charge lifetime at perovskite/TiO2 interface is prolonged from 62μs to 107μs. � PCEs of PSCs and DSSCs are enhanced by ~16% and ~33% after Nb doping, respectively. � Heavy Nb doping induces substantial changes of properties of TiO2 NRAs.

A R T I C L E I N F O

A B S T R A C T

Keywords: Niobium doping TiO2 nanorod array Perovskite solar cell Dye-sensitized solar cell Charge transport Recombination

One-dimensional (1-D) rutile TiO2 nanorod arrays (NRAs) synthesized by a hydrothermal method suffer from low electrical conductivity and large amounts of surface defects, hindering their further applications. Nb doping is thus introduced to modify their electronic properties. Results indicate that light Nb doping reduces rod nano­ sizes, increases electron concentrations, decreases surface defective oxides and lowers conduction band of the TiO2 NRAs, while heavy doping induces transformations of morphologies and crystalline orientations as well as occurrences of compositional deviations and low oxidative states of Ti3þ. After 0.1 mol% and 1 mol% Nb in­ corporations, device efficiencies are substantially improved by ~16% and ~33% for the model perovskite and dye-sensitized solar cells, respectively, which are ascribed to reduced recombination at the perovskite/TiO2 interfaces (e.g. charge lifetime increasing from 62 μs to 107 μs) and improved electron transport through the photoanode of TiO2 NRAs (e.g. electron diffusion length increasing from ~14 μm to ~50 μm). Our study verifies that Nb doped 1-D TiO2 NRAs are versatile electron transporting materials in different kinds of emerging solar cells, and are also potential for other fields including photocatalysis, sensors and batteries etc.

* Corresponding author. School of Advanced Materials and Nanotechnology, Xidian University, 266 Xinglong Section of Xifeng Road, Xi’an, 710126, Shaanxi, People’s Republic of China. ** Corresponding author. Key Lab of Wide Band-Gap Semiconductor Materials and Devices, Xidian University, Xi’an, 710071, Shaanxi, People’s Republic of China. E-mail addresses: [email protected] (P. Zhong), [email protected] (X. Ma). https://doi.org/10.1016/j.jpowsour.2020.227715 Received 28 September 2019; Received in revised form 20 December 2019; Accepted 3 January 2020 Available online 29 January 2020 0378-7753/© 2020 Elsevier B.V. All rights reserved.

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Journal of Power Sources 450 (2020) 227715

1. Introduction

three types of emerging PV devices as the ETLs, including perovskite, dye-sensitized and Cu2O/TiO2 NRAs heterojunction solar cells. Furthermore, an in-depth study is conducted to understand the in­ fluences of Nb incorporation on charge transport and recombination as well as device performances in the model PV cells.

Utilization solar energy is one of the most effective solutions to solve the critical energy and environmental issues globally [1]. Nowadays, photovoltaic (PV) technologies attract more and more attentions both in industries and in research communities. However, conventional solar cells based on Si and GaAs etc. have an intrinsic drawback of high costs as compared to other energy sources, which would hinder their further developments [2,3]. In the last twenty years, some emerging PV can­ didates with low costs are potential alternatives to traditional solar technologies, including perovskite solar cells (PSCs) [4–8], dye-sensitized solar cells (DSSCs) [9–11], quantum-dot solar cells (QDSCs) [12,13] and organic solar cells (OSCs) [14–16] etc. Typically in emerging solar cells, a light-active layer (LAL) absorbs photons of sun­ light, subsequently converted into excitons or electron-hole pairs, which are further separated, and then the photo-generated charge carriers transport through an electron transport layer (ETL) or a hole transport layer (HTL), before being collected by the respective electrodes [17,18]. To achieve a high-performance PV device, the aforementioned compo­ nents and their interfaces are equally important. The TiO2 nanoparticular mesoporous film is regarded as one of the most important ETLs in emerging solar cells [19–25], due to their high specific surface areas for creation of large interfaces, interconnected pathways for charge transport, suitable energy band structures, and flexibilities for surface modification etc. However, this class of nano­ materials suffers from low electrical conductivities and large amounts of defects on their surfaces and at grain boundaries, which would decrease device performances [26–28]. Hence, one-dimensional (1-D) TiO2 nanostructure arrays were fabricated perpendicular to transparent conducting oxide (TCO) substrates, in order to improve charge transport in emerging solar cells, including TiO2 nanorod arrays (NRAs) [29–34], nanowire arrays (NWAs) [35–37], nanotube arrays (NTAs) [38–40] and so on. On one hand these vertically-aligned nanostructures provide direct pathways for electron transport, on the other hand they preserve large surface areas for charge separation. Among the various 1-D nanostructures, TiO2 nanorod arrays synthesized from a hydrothermal method have appealed great attentions in the PV and other photo­ electrochemical (PEC) fields, due to their single-crystalline feature as well as easy processing [29–34]. Generally, the most stable rutile phase of TiO2 NRAs can be obtained in these hydrothermal reactions [41]. Generally, rutile is less active than anatase for optoelectronic applica­ tions. Thus, it is still challenging to improve the microscopic electrical property of rutile TiO2 NRAs derived from a hydrothermal way, and thus to enhance the performances of solar cells incorporated with such nanostructures [42]. Except for morphological control of nanostructures, another approach to increase the conductivity of TiO2 is doping with guest ions such as Li, W, Bi, Nb, Y, Ta, Sn, Mg and so on [5,43–49]. These atomic doping could not only change the mobility and electron concentration of TiO2, but also greatly affects other properties like shifting conduction band (CB) positions, passivating surface defects and influencing mor­ phologies of TiO2 nanostructures as well as the top LALs. Recently, Nb doped TiO2 was reported to substantially increase the conductivity of pristine TiO2 regardless of anatase or rutile phase, which might find widespread applications in TCO electrodes, solar cells, gas sensors etc [46,50–55]. In most cases, TiO2 dense films or mesoporous films are employed as the scaffolds for Nb doping, because in these systems doping effects are simplified. Until now, it is seldom reported about the influences of Nb doping on the comprehensive properties of rutile TiO2 NRAs prepared by a hydrothermal synthesis, because Nb incorporation might simultaneously affect complex characteristics of NRAs, which makes the system complicated [56]. In this work, a systematic study is carried out to investigate the ef­ fects of Nb doping on the morphologies, crystallinity, chemical com­ positions, optical and electrical properties of the TiO2 NRAs synthesized by a hydrothermal reaction. Then the TiO2 NRAs are incorporated into

2. Experimental 2.1. Synthesis of TiO2 NRAs [29] FTO (F doped SnO2) glass substrates (7 Ω/□, Wuhan Geo, China) were cleaned in deionized (DI) water, acetone and ethanol in an ultra­ sonic bath for 10 min, respectively, followed by N2 blow drying. A TiO2 seed film was first prepared by soaking the FTO substrate into 0.2 M TiCl4 aqueous solution at 70 � C for 30 min, before calcinations in an oven at 550 � C for 1 h. The reaction solution for hydrothermal synthesis consisted of 15 ml DI water, 15 ml HCl (36.5–38 wt%) and 0.25–0.5 g Butyl titanate (TBOT) (or 1.5 ml TiCl4 for grow long TiO2 NRAs). Nb doping was realized by adding NbCl5 into the above solution. After stirring for 10 min, the mixture solution was moved to a 50 ml Teflonliner stainless steel autoclave. The TiO2 seed film/FTO substrates were placed against the Teflon-liner wall at ~70� with their conducting side facing down. The hydrothermal reaction was conducted at 150–180 � C in an oven for 0–6 h. After synthesis, the autoclave was cooled by flowing cold water, immediately. Finally, the samples were collected after rinsing by DI water extensively and drying by N2 flow. The Nb doping levels were defined as the atomic ratios of Nb/Ti in the whole context. 2.2. Fabrication of perovskite solar cells Short TiO2 NRAs (<800 nm) were synthesized by the above pro­ cedure using TBOT as the precursor. Before device preparation, the TiO2 NRAs/TiO2 seed film/FTO substrates were calcined in an oven at 400 � C for 1 h. Then, a UV-ozone treatment was conducted for the substrates for 15 min. The perovskite films were deposited by a two-step method. Firstly, 1.3 M PbI2 dissolved in a mixture solution of DMF and DMSO (V: V ¼ 95:5) was spin-coated onto the TiO2 NRAs/TiO2 seed film/FTO substrates at 1500 rpm for 30 s, subsequently followed by heat treatment on a hot plate at 70 � C for 1 min. Secondly, a mixed amine salt solution consisting of 60 mg FAI, 6 mg MACl, 6 mg MABr and 1 ml isopropanol was further spin-coated onto the PbI2 film at 1300 rpm for 30 s. Then, the samples were annealed on a hot plate at 150 � C for 15 min in air with humidity of 20–40% to form perovskite films. It should be stressed here all the device fabrication steps were carried out in a groove box filled by N2 except the aforementioned annealing step for the perovskite films. Thirdly, a HTL was deposited onto the perovskite film by spin coating at 4000 rpm for 30 s using a mixture solution of 72.3 mg spiro-OMeTAD, 17.5 μL Li-TFSi dissolved in acetonitrile (520 mgml 1), 28.8 μl 4-tertbutylpyridine and 1 ml chlorobenzene. Finally, an Ag top electrode with a thickness of ~100 nm was evaporated onto the HTL, with device areas of 0.04 cm 2 defined by a shadow mask. The devices should experience an oxidation process for spiro-OMeTAD for about 12 h in air before measurements. 2.3. Assembly of dye-sensitized solar cells Long TiO2 NRAs (>6 μm) were synthesized using TiCl4 as the pre­ cursor. An additional etching step was employed in a mixed bath of DI water and HCl, hydrothermally treated at 150 � C for 3 h, followed by an annealing step in an oven at 400 � C for 1 h. The DSSC encapsulation was similar to that reported previously [29]. Briefly, the TiO2 NRAs/TiO2 seed film/FTO substrates were first annealed at 80 � C for 15 min, before soaking into 0.3 mM ethanol solution of ruthenium complex N719 at room temperature for 24 h. An ultrathin layer of platinum (Pt) was deposited onto FTO substrates by DC sputtering as the counter electrode. 2

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Then, the Pt-FTO counter electrode and the dye-sensitized photoanode were heat sealed by using a hot-melt polymer (25 μm thick, Solaronix SX1170-25). Finally, the liquid electrolyte (Wuhan Geo, China), mainly consisting of 1-propyl-3-methyl-imidazolium iodide, LiI, I2, Guanidine thiocyanateand 4-tert-butylpyridine in acetonitrile was injected into the device through a drilled small hole on the counter electrode. The active area was 0.19 cm 2 defined by the hot-melt polymer.

electrode mode with 1 M NaOH solution (PH ¼ 13.9) as the electro­ lyte. The TiO2 NRAs/TiO2 seed film/FTO substrate was employed as the working electrode; a Pt sheet was used as the counter electrode; a saturated calomel electrode (SCE, saturated KCl) served as the reference electrode. The PV performance of solar cells was measured by recording the current density-voltage (J-V) curves under an illumination of an AM 1.5G solar simulator (100 mWcm 2, Zolix, SS150). The measurements of transient photocurrent/photovoltage spectroscopy (TPC/TPV), intensity-modulated photocurrent/photovoltage spectroscopy (IMPS/ IMVS), and electrochemical impedance spectroscopy (EIS) were all performed on a PEC test system (CIMPS-II, Zahner). The detailed pa­ rameters were described in supporting information.

2.4. Characterization and measurements The morphologies were observed by a field-emission scanning elec­ tron microscope (FESEM, FEI ApreoHiVac). The crystalline properties were determined by X-ray Diffraction (XRD, Bruker D8Advance).The chemical compositions were measured by X-ray Photoelectron Spec­ troscopy (XPS, Thermo Fisher Scientific, K-Alpha) and Energy Disper­ sive Spectrometer (EDS, Oxford X-Max 50). The UV–Vis absorption spectra were recorded by using a PE Lambda 950 UV–Vis–NIR spectro­ scope. Mott-Schottky (M S) measurements were performed with an electrochemical workstation (Zahner, Zennium) in a standard three-

3. Results and discussion 3.1. Properties of Nb doped TiO2 NRAs Fig. 1(a1-d1) show the SEM top views of the as-prepared TiO2 NRAs synthesized using TBOT as the precursor with different Nb doping levels,

Fig. 1. (a1-d1) top views, (a2-d2) cross-sectional views of the TiO2 NRAs synthesized by a hydrothermal method using TBOT as the precusor at 180 � C for 4 h as a function of the Nb doping levels; (e) XRD patterns, (f–i) high-resolution XPS spectra, (j) UV–Vis spectra with the tauc plots in the inset, (k) M S plots, (l) energy band structures of the undoped and Nb doped TiO2 NRAs. 3

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while Fig. 1(a2-d2) show the corresponding cross-sectional images. The reaction conditions are identical for all samples at 180 � C for 4 h. Although the Nb incorporated levels are low, the morphologies of TiO2 NRAs changes obviously with increasing the doping levels. The nanorod shape of the undoped sample (0 mol%) is square, with an average diameter of ~150 nm. The rod tips have rough secondary nanostructures [29]. From the cross-sectional view in Fig. 1(a2), the undoped nanorods pack closely with a length of ~3.5 μm. With 0.1 mol % Nb doping, it can be observed from Fig. 1(b1), the TiO2 nanorods become slimmer with a rod diameter of ~40 nm, and the rod density becomes smaller than the undoped sample. The length of the NRAs decreased to ~1.3 μm sharply at 0.1 mol% as shown in Fig. 1(b2). Increasing the doping level to 0.2 mol%, as shown in Fig. 1(c1,c2), the profiles of the nanorods are better defined, with their diameter and length decreased to ~ 20 nm and ~1 μm, respectively. However, further increasing the Nb doping level to 0.5 mol%, the nanorod diameters increase to ~50 nm, while the length keeps on decreasing to ~0.9 μm, as shown in Fig. 1(d1,d2). The corre­ lation plots between sizes, growth rate of nanorods and Nb doping levels are displayed in Fig. S1. Fig. 1(e) shows the XRD patterns of the as-prepared TiO2 NRAs with different Nb doping levels ranging from 0–5 mol%. For the undoped sample (0 mol%), two main diffraction peaks located at ~ 62.8� and ~36.1� are assigned to (002) and (101) crystal planes of the tetragonal rutile TiO2 (P42/mnm), respectively, in terms of JCPDS card No. 88–1175. Besides, some other peaks also appear such as (110), (111), and (211). The significantly-enhanced (002) peak as compared to other peaks suggests that the TiO2 film is highly oriented, and the TiO2 nanorods grow along the [001] direction with the growth axis perpen­ dicular to the FTO substrate. It can be observed that with increasing the Nb doping level (<0.5 mol%), the intensities of both of the (002) and (101) peaks become more and more weaker, which might be due to the more and more thinner films with the Nb doping levels, while there are no observed peak shifts. With increasing the doping level to 2 mol%, the (002) peak disappears, while the (110) peak located at ~ 27.4� becomes the dominated one. Besides, a list of other peaks assigned to rutile ap­ pears at 2 mol% including (310), (220), (211), (210), (111), etc. With further increasing the Nb doping level to 5 mol%, only some small peaks exist probably due to the fact that the film is very thin. Fig. 1(f–i) show the high-resolution XPS patterns of the Nb3d, Ti2p, O1s and valence band (VB), respectively. The undoped sample and the Nb doped one with 0.5 mol% are compared. All the peaks are calibrated by the C1s peak (adventitious carbon) centered at 284.8 eV as shown in Fig. S2. As shown in Fig. 1(f), there exist symmetric spin-orbit doublets with Nb 3d5/2 binding energy of ~207 eV and Nb 3d3/2 binding energy of ~209.7 eV for the Nb doped sample, indicating of the presence of Nb5þ [57], while no signals are observed for the undoped one. The Ti 2p scans as shown in Fig. 1(g) exhibit asymmetric doublet peak resulting from spin-orbit splitting, with Ti 2p3/2 peak located at ~ 458.4 eV and the splitting energy of ~5.7 eV, indicating that Ti exists in the form of Ti4þ. No peak shift and appearance of Ti3þ are observed for the Nb doped sample as compared to the undoped one due to slight doping. As to the O1s signals in Fig. 1(h), the broad and asymmetric O1s peak can be deconvoluted into three peaks, centered at 529.6 eV, 531.1–531.3 eV and 531.8 eV, which are assigned to crystal lattice oxygen, defective oxides, and the chemisorbed oxygen species. As shown in Table S1, the defective oxide content is significantly reduced from 18% for the undoped TiO2 to 7% for the 0.5 mol% Nb doped sample, indicating that the Nb doping can reduce oxygen vacancies, which might influence the electronic structures of the TiO2 NRAs. Such a phenomenon was also found in a case of Ta doped TiO2 nanorods [58]. VB XPS are carried out to evaluate the density of states of VB of TiO2 as shown in Fig. 1(i). The edge of the VB maximum is measured to be ~2.5 eV below the Fermi energy for the undoped sample, while this value is increased to ~2.75 eV for the Nb doped one. Fig. 1(j) shows the UV–Vis absorption spectra of the TiO2 NRAs with different Nb doping levels. All the samples are only active in the UV

region. Negligible absorption are detected for all samples in the visible range (400–800 nm). The optical band gap (Eg) can be measured to be ~2.88 eV for the undoped sample as shown in the inset of Fig. 1(j), and the values are slightly reduced to close to 2.80 eV for the Nb doped ones. Fig. 1(k) shows the M S curves. The flatband potentials (Efb) can be calculated by the M S equation [59]: 1=C2 ¼ ð2 =e0 εε0 Nd ÞðEapp Efb kT =e0 Þ, where C is the differential capacitance of the Helmholtz layer, e0 is the electron charge, ε is the relative dielectric constant of TiO2 (~173 [60]), ε0 is the vacuum permittivity, Nd is the donor density, Eapp is the applied bias at the electrode, and k is the Boltzmann constant. Efb are measured to be 1.046 V and 0.886 V vs SCE in 1 M NaOH aqueous solution for the undoped and Nb doped samples, respectively. Both of the samples display a positive slope in the M S plots, a char­ acteristic of n-type semiconductors. The Nb doped sample shows a smaller slope than the undoped sample, indicating of an increase of donor density. The Nd can be calculated using the following formula

[59]: Nd ¼ ð2=e0 εε0 Þ½dð1=C2 Þ=dEapp � 1 . The Nd can be determined to be 2.1 � 1018 cm 3 and 5.8 � 1018 cm 3 for the undoped and Nb doped samples, respectively. Then the conduction band (CB) edge (ECB) can be obtained according to the equation: ​ ECB ¼ Efb þ kTlnðNc =Nd Þ, here Nc is the effective density of states with typical value of 1020 cm 3 [61–63]. In combination with the Eg and VB results, the energy band structures can be depicted as shown in Fig. 1(l). The detailed calculation process is described in the supporting information. It can be observed that the ECB shows a positive shift from 4.42eV down to 4.61eV (vs vacuum level) after Nb doping. Besides, the Fermi energy level (EF) moves towards the CB after doping. Such a change of electronic structures might facilitate charge injection at the hetero-interface between TiO2 and another donor, as well as improve electron transport in TiO2. To grow TiO2 NRAs with longer length (>6 μm), TiCl4 can be chosen as the precursor in place of TBOT to speed up the hydrothermal reaction [29]. Nb doping with a wide range from 0–10 mol% is carried out. Fig. 2 (a1-d1) show the SEM top views of the morphologies of the TiO2 NRAs with some different Nb doping levels with their enlarged images in the insets, and cases with other doping levels are exhibited in Fig. S3. The reaction conditions are the same for all samples at 150 � C for 6 h. It can be observed that the surface morphologies are similar at low doping levels (<0.5 mol%). The nanorods with an average diameter of ~100 nm packed closely in large areas. Some cracks divide the whole surface into small regions with sizes of ~10 μm, probably due to surface tension in the drying process. With increasing the Nb doping level to 1 mol%, the nanorods show a tendency of aggregation in each region, while there is no obvious change for feature sizes. However, when increasing the doping level to 5 mol%, it is surprisingly to discover that the rod-shape features disappear in the whole sample, while the big cracks still exist. With further increasing the doping level to 10 mol%, neither the rod shapes nor the cracks exist, remaining a smooth and compact film on the FTO substrate as shown in Fig. S3. From the cross-sectional images in Fig. 2(a2-d2) and in Fig. S3, it can be observed that nanorods stand vertically on substrates when the Nb doping level is below 1 mol%, and the nanostructures disappear at high doping levels (i.e. > 5 mol%). The film thickness is also influenced by the Nb doping level as shown in Fig. S3. To the best of our knowledge, such a phenomenon related to morphological transformation from nanostructures into compact films is first reported in the situation of Nb doped TiO2. The formed compact film of Nb doped TiO2 might find applications in transparent conducting electrodes etc. EDS are employed to probe the Nb contents in the TiO2 films with different doping levels as shown in Fig. S4 and Table S2. Fig. 2(e) shows the comparison plots of Nb/Ti ratios between the predicted (in pre­ cursors) and the measured results. At low doping levels (<1 mol%), the predicted and the measured data coincide well. However, at high doping levels (>5 mol%), the measured Nb/Ti ratios are as two times high as the predicted results. XRD are carried out to study the crystalline structures of the TiO2 NRAs with different Nb doping as shown in Fig. S5. Below 4

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Fig. 2. (a1-d1) top views with the enlarged images in the insets, (a2-d2) cross-sectional views of the TiO2 NRAs using TiCl4 as the precusor at 150 � C for 6 h as a function of the Nb doping levels; (e) comparsion plots of the predicted (in precusor) and measured atomic ratios of Nb/Ti of the TiO2 NRAs as a function of the Nb doping levels; (f) the (002) diffraction peaks of the TiO2 NRAs with different Nb doping levels; (g) high-resolution Ti2p XPS spectra of the undoped and 5 mol% Nb doped TiO2 NRAs.

mol%, the (002) crystal plane of rutile is the preferred orientation for all samples, which is in good agreement with the growth direction of TiO2 nanorods. However, at 5 mol% it can be observed that the preferred orientation is transformed into the (101) crystal plane of rutile. At the range of 6–8 mol%, there are no dominated diffraction peaks. Further increasing the Nb doping levels to 9 mol% and 10 mol%, the (110) crystal plane becomes the preferred orientation. Combined with the aforementioned SEM results, it is inferred that the shapes of TiO2 NRAs are well preserved for the samples with Nb doping levels below 4 mol%, and high doping levels over 5 mol% would lead to great changes of the morphologies and the crystalline properties. Fig. 2(f) shows the highresolution XRD patterns of the (002) diffraction peak of rutile TiO2 as a function of the Nb doping levels. Slight Nb doping (<0.5 mol%) does not influence the peak position, while doping over 1 mol% would induce a peak shift towards lower two theta values. Specifically, the value of the 2 mol% is observed to be ~0.1� lower than that of the undoped sample. This phenomenon is ascribed to the larger ionic radius of Nb5þ (0.64 Å) than that of the Ti4þ (0.61 Å), in term of the Brigg equation: 2dsinθ ¼ λ. Fig. 2(g) shows the high-resolution XPS spectra of Ti2p of the undoped TiO2 NRAs and the sample with 5 mol% Nb doping. Interestingly, a peak shift of ~0.3 eV towards the higher binding energy is apparently observed upon 5 mol% Nb doping, which might result from the bigger electronegativity of Nb (1.6) than that of Ti (1.54). Besides, minor amounts of Ti3þ are observed in the sample with 5 mol% Nb doping, evidenced by appearance of additional peak located at ~457.97 eV. In the process of substituting Ti4þ in TiO2 by Nb5þ, charge compensation is generally achieved in two ways: I) creation of one Ti cation vacancy per four Nb [Equation (1)]; II) reduction of Ti4þ to Ti3þ per Nb[Equation (2)] [46,55]:

Ⅳ 2Nb2 O5 þ 4TixTi →4Nb● Ti þ VTi þ 4TiO2 þ O2

(1)

1 5 ’ Nb2 O5 þ TixTi ¼ Nb● Ti þ TiTi þ O2 2 4

(2)

Low Nb doping level favors the process in Equation (1), while high doping content selects the process in Equation (2), which is confirmed in our experiments as well. As shown in Fig. 1(g) and Fig. S6, no signals related to Ti with low oxidation states appear by a light doping level of 0.5 mol% Nb in the samples synthesized using TBOT or TiCl4 as the precursors, and Ti vacancies might happen in this situation. However, as show in Fig. 2(g), the occurrence of Ti3þ suggests that heavy Nb doping as high as 5 mol% might follow another approach. Based on the above results, Nb doping levels would complicatedly influence the compre­ hensive features of the rutile TiO2 NRAs prepared by a hydrothermal synthesis, and thus the electronic properties as well as some potential applications. In order to provide well-defined charge transport pathways, we choose relatively low Nb doping levels (<2 mol%) to preserve the nanorod feature in our samples in the following studies on photovoltaic devices. 3.2. Interfacial charge transfer in perovskite solar cells Since a controlled thickness of the LAL is strictly demanded in PSCs (generally < 1 μm), the lengths of the TiO2 NRAs should be adjusted in this range. Through a complex modulation of the reaction conditions, the nanorod lengths can be controlled below 800 nm as shown in Fig. S7. To simplify the research, two types of comparison devices are studied incorporated with undoped TiO2 NRAs and 0.1 mol% Nb doped TiO2. 5

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Fig. 3(a) shows the top views of the SEM morphologies of the perovskite films deposited on the substrates of undoped and Nb doped TiO2 NRAs by a two-step way, and the insets are the cross-sectional views. Both films have a similar thickness of 600–700 nm. However, their mor­ phologies are slightly different. The perovskite film deposited on undoped TiO2 NRAs is observed to be a little more densely-packed than that on Nb doped TiO2 NRAs. Besides, the grain shapes of the perovskite films are also different. The perovskite film on the Nb doped substrate has elongated grains as compared to that on undoped substrate. Fig. 3(b) shows the UV–vis–NIR absorption spectra of the perovskite films, and the inset is the XRD patterns. Both samples exhibit high absorptions in the whole visible range. As a comparison, the absorption intensity of the Nb doped one is slightly decreased, which might be due to its relativelyloose film as observed in Fig. 3(a). The XRD patterns suggest that both perovskite films have polycrystalline properties with several main diffraction peaks of (110), (211) and (220) etc. The intensity of the (211) peak for the undoped sample is observed to be a little higher than that for the Nb doped one, which might be closely related to the morphol­ ogies of the perovskite films. Above results suggest that the perovskite films prepared on the undoped and the Nb doped TiO2 NRAs based substrates only have very small differences. The PSCs are fabricated based on the TiO2 NRAs with the cell structure as shown in Fig. 3(c). The J-V curves in Fig. 3(d) shows that the power conversion efficiency (PCE) is greatly increased from 13.5% for the undoped device to 15.7% for the Nb doped cell, with an enhance­ ment of ~16%. The PCE enhancement is mainly due to the increase of the open-circuit voltage (Voc) (from 1.051 V to 1.111 V) and the fill factor (FF) (from 64.5% to 70.4%), while the short-circuit current density (Jsc) (from 20.47 mAcm 2 to 20.46 mAcm 2) remains stable regardless of the Nb doping. In addition, the hysteresis is apparently

reduced after Nb doping, with the hysteresis index decreased from 0.268 to 0.077 (figure not shown).The values of Voc and FF are closely related to two main recombination processes in PSCs, namely trap-assisted nonradiative recombination in perovskite films and interface recombi­ nation at perovskite/charge transporting layer interfaces. The correla­ tion between Voc and recombination can be described as follows [64]: μ ¼ μoc;rad þ kTlnðEQYPL Þ, where μ is the free energy of photo-exited electron-hole pairs, equal to the energy splitting of the quasi-Fermi levels of electrons in the conduction and valence bands. EQYPL is referred to the external electroluminescence quantum yield, which is inversely proportional to nonradiative recombination. μoc,rad is the free energy only with radiative recombination (EQYPL ¼ 1). In fact, Voc closely associated with μ is proportional to EQYPL, and is thus inversely proportional to nonradiative recombination. In the present study, similar perovskite films are obtained on both of the undoped and Nb doped TiO2 substrates, which should contribute almost equally to the nonradiative recombination and device Voc. Herein, interfacial recom­ bination induced by imperfect interfacial band alignments and interface defects would have dominant impact on Voc. The maximum FFmax of PSCs can be determined by using the equation [65]: FFmax ¼ voc lnðvoc þ0:72Þ , voc þ1

here voc is proportional to the open-circuit voltage (Voc).

Consequently, the FF of PSCs is closely correlated with the cell Voc. As discussed above, Voc is greatly affected by the interface recombination. As a result, this factor would also have great influence on FF. Baena et al. discovered that a CB mismatch existed between the mixed perovskite (FAPbI3)0.85(MAPbBr3)0.15 and TiO2, that is, the CB of perovskite is 300 meV below that of TiO2. They pointed out that this band misalignment would cause accumulation of photo-generated charge carriers, leading to undesirable consequences including deteri­ orative device performance as well as pronounced hysteretic effect.

Fig. 3. (a) SEM morphological images, (b) UV–Vis–NIR spectra and XRD patterns (inset) of the perovskite films deposited on the undoped and 0.1 mol% Nb doped TiO2 NRAs; (c) device architecture of the PSCs using the (Nb doped) TiO2 NRAs as the ETLs; (d) J-V curves of the PSCs (inset is the derived PV parameters); (e–h) correlation plots between detailed PV parameters and lengths of undoped TiO2 NRAs; (i–l) correlation plots between detailed PV parameters and lengths of Nb doped TiO2 NRAs. 6

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Journal of Power Sources 450 (2020) 227715

Furthermore, they figured out a desirable solution by using SnO2 with a deeper CB than TiO2 as the electron transporting material, which is now widely employed in present ultra-high efficiency PSCs [66,67]. In the present study, light Nb doping is verified to lower the CB of rutile TiO2 NRAs by ~0.2 eV as shown in Fig. 1(l), which should be beneficial for better band alignment between perovskite and TiO2, and thus improved charge transfer and reduced recombination at the interface. In addition, light Nb doping can reduce the oxygen defects on the surface of TiO2 NRAs, which might be another reason to reduce interfacial recombina­ tion. Fig. 3(e–l) show the comparison plots of the PV parameters of the undoped and Nb doped devices as a function of the lengths of TiO2 NRAs, and the detailed PV parameters are also exhibited as shown in Table S3. It can be observed that the PV performances of the Nb doped devices show less dependence on the lengths of TiO2 NRAs than those of the undoped PSCs. For the undoped devices, all parameters (Voc, Jsc, FF and PCE) drop sharply when using relatively-long nanorods (240 nm) as compared to using the shortest nanorods (140 nm), indicating that the charge transport is inert through long nanorods. However, for the Nb doped PSCs, an obvious decrease of PV parameters is only observed by using the longest nanorods (780 nm), suggesting that Nb doping could greatly improve the charge transport in TiO2 nanorods as well as elec­ tron collection at the corresponding electrode, which would be studied further in the next part of this paper. The enhanced charge transport in

the TiO2 NRAs due to Nb doping might also be additional reason for reducing interfacial recombination. Based on the aforementioned dis­ cussions, it is inferred that by using Nb doped TiO2 NRAs as the ETL in PSCs, the PCE enhancement (mainly contributed by the improvements of Voc and FF) is mainly due to decreased interfacial recombination. A series of techniques are employed to probe the charge dynamics in PSCs as shown in Fig. 4. It can be observed from Fig. 4(a) that Jsc re­ covers rapidly from dark to light every circle for the Nb doped device as compared to the undoped one, indicating that an efficient charge transfer occurs at the perovskite/Nb doped TiO2 interface, due to wellaligned band levels and passivation of surface defects on TiO2. From Fig. 4(b), Voc of the Nb doped PSC shows prolonged periods in both the recovering and decay processes, suggesting that charge carriers survive long time before recombination with a comparison to that of the undoped device. Furthermore, the TPC and TPV spectra as shown in Fig. 4(c and d) provide direct evidences on improved charge transport properties upon Nb doping. The charge extraction time of the Nb doped PSC is determined to be 14 μs, which is faster than that of the undoped device (18 μs). This facilitated charge extraction originates from the modified interface after Nb doping. The charge lifetime of the Nb doped device is measured to be 107 μs, almost twice as that of the undoped one (62 μs), which should be ascribed to reduced interfacial recombination. EIS is also carried out to study the charge recombination properties in

Fig. 4. chopped (a) Jsc, (b) Voc of PSCs based on (0.1 mol% Nb doped) TiO2 NRAs with interval illuminations; (c) TPC, (d) TPV plots of PSCs based on (0.1 mol% Nb doped) TiO2 NRAs; correlation plots between (e) Rs, (f) Rct and light intensity for PSCs based on undoped and 0.1 mol% Nb doped TiO2 NRAs, derived from EIS measurements. 7

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Journal of Power Sources 450 (2020) 227715

PSCs. The Nyquist plots and the equivalent circuit are displayed in Fig. S8. The fitted electrical parameters including the series resistance (Rs) and recombination resistance (Rct) are plotted against the light in­ tensity as shown in Fig. 4(e and f). After Nb doping, the Rs are decreased, while the Rct are increased obviously regardless of the light intensity, further evidencing substantially-enhanced charge transport and greatlyreduced recombination exist in the PSCs using Nb doped TiO2 NRAs as compared to those using undoped NRAs. Except for perovkiste solar cells, the performance of the Cu2O/TiO2 heterojunction PV device is also found to be greatly improved by using Nb doped TiO2 NRAs in place of undoped ones, demonstrating the versatility of our Nb doped TiO2 nanorods in other thin-film PV systems (Fig. S9).

Table 1 PV parameters of DSSCs based on TiO2 NRAs with the Nb doping levels. Device

Voc/V

Jsc/mAcm

0 mol% 0.2mol% 1mol% 2mol%

0.655 0.645 0.615 0.585

9.83 10.35 14.43 10.99

2

FF/%

PCE/%

60.2 61.5 58.3 57.2

3.88 4.10 5.17 3.68

those of the device PCE. The values of Jsc first increases from 9.83 mAcm 2 (0 mol%) to 10.35 mAcm 2 (0.2 mol%), and attains a highest value of 14.43 mAcm 2 at 1 mol% before decreasing to 10.99 mAcm 2 at 2 mol%. The electron transport properties in DSSCs are further investigated by the IMPS/IMVS techniques. The complex plane plots of IMPS/IMVS are displayed in Fig. S10. A series of parameters can be obtained including electron transport time (τd), electron lifetime (τr), electron diffusion coefficient (Dn), electron diffusion length (Ln) and electron collection efficiency (ηcc). The τd and τr stand for the time constants of the injected electrons diffusing through the TiO2 photoanode and recombining with electrolyte as well as oxidized states of dye molecules, respectively. They can be attained by the equations [36]: τd ¼ 1=2πf d , τr ¼ 1=2πf r , where fd and fr are the characteristic frequency minimum of the imaginary components of the IMPS/IMVS semicircles, respectively. As shown in Fig. 5(c), τd that is on the order of milliseconds first de­ creases apparently at 0.2 mol% and 1 mol%, and then increases at 2 mol %, indicating that moderate Nb doping could speed up the electron transit in TiO2 nanorods, while overdoping would slower the electron migration. As to τr [Fig. 5(d)], a small decrease is firstly observed after 0.2 mol% Nb doping. However, increasing the doping level to 1 mol%, it is surprisingly to discover that τr boosts by 3–4 folds as compared to the undoped device, with the maximum value of τr approaching 1 s in the range of low light intensities. With further increasing the Nb doping level to 2 mol%, τr shows obvious decrease. Above results suggest that a moderate Nb doping level could lead to substantially-prolonged lifetime for photo-generated electrons. Taking the thicknesses of used TiO2 NRAs (d) into account (shown in

3.3. Electron transport and collection in dye-sensitized solar cells In the above studies of PSCs, relatively-short TiO2 NRAs (<800 nm) are employed. The Nb doping of nanorods has substantial influence on the perovskite/TiO2 interface, resulting in enhanced interfacial charge transfer as well as reduced recombination. However, the effect of Nb doping on charge transport and collection through TiO2 nanorods is hardly identified. To tackle this issue, relatively-long TiO2 NRAs (6–11 μm) as shown in Fig. 2 are used to assemble DSSCs, with the device architecture as exhibited in Fig. 5(a). Since the bottom NRAs would not affect the properties of the absorbed dye molecules and the liquid electrolyte, opportunities are given to investigate the electron transport and collection in the photoanode using TiO2 nanorods, and to evaluate the effect of Nb doping. Fig. 5(b) shows the J-V curves of the DSSCs using TiO2 NRAs as a function of the Nb doping levels, and the detailed PV parameters are obtained in Table 1. It can be observed that the device PCE first increases from 3.88% for the undoped DSSC (0 mol%) to 4.10% for the 0.2 mol% device, and then continues to increase until gets a maximum value of 5.17% at 1 mol%, and at last decrease to 3.68% at 2 mol%. A substantial PCE enhancement of ~33% is achieved for the 1 mol% Nb doped DSSC as compared to the undoped one. Specifically, the Voc keeps on decreasing slightly with the Nb doping level from 0.655 V (0 mol%) to 0.585 V (2 mol%), probably due to downward shift of the CB of TiO2 after Nb doping, which is observed in Fig. 1(l). The FF shows little fluctuation with the Nb doping levels (around 60%). It should be stressed here that the variations of Jsc make major contributions for

Fig. S3), Dn can be determined by the equation [29]: Dn ¼ d2 =4τd . As shown in Fig. 5(e), Dn increases sharply with light Nb doping (<1 mol %), confirming the role of Nb incorporation in promoting charge transport in the long TiO2 nanorods, while the value exhibits apparent

Fig. 5. (a) device structure of the DSSCs based on (Nb doped) TiO2 NRAs; (b) J-V curves of the DSSCs as a function of the Nb doping levels; correlation plots between (c) τd, (d) τr, (e) Dn, (f) Ln, (g) ηcc and light intensity as a function of the Nb doping levels, derived from IMPS/IMVS spectra; (h) EIS Nyquist plots of the DSSCs, inset is the equivalent circuit. 8

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Journal of Power Sources 450 (2020) 227715

decrease at high Nb doping of 2 mol%. Ln and ηcc are two important indexes evaluating the electron collection property considering both of the charge transport and recombination. Ln reflects the distance of electron diffusion in the TiO2 NRAs before back recombination, and ηcc represents a comprehensive efficiency for electron collection. Ln and ηcc pffiffiffiffiffiffiffiffiffi can be calculated by the following expressions [23,29]: Ln ¼ Dn τr , ηcc ¼ 1 τd =τr . As shown in Fig. 5(f), Ln first increases from ~14 μm (0 mol%) to ~ 20 μm (0.2 mol%), and then gets remarkable values of ~50 μm (1 mol%), and finally decreases to ~ 28 μm (2 mol%). Then variation of ηcc as shown in Fig. 5(g) is similar to that of Ln. Especially, the ηcc is substantially improved from 92% for the undoped DSSC to over 98% for the 1 mol% Nb doped device. It should be mentioned that the changing tendencies of Ln and ηcc are similar to those of the Jsc and PCEs of DSSCs with different Nb doping levels, suggesting that the modified electronic properties induced by Nb doping might be the key factor to affect device performance. With light Nb doping (<1 mol%), on one hand the CB of TiO2 is positively shifted, leading to an enlarged driving force for elec­ tron injection, namely the difference between the lowest unoccupied molecular orbital (LUMO) of the absorbed dye and the CB of TiO2. On the other hand, light doping could remarkably increase the electron density (n) as shown in Fig. 1(k), resulting in enhanced electrical con­ ductivity (σ) according to the equation [46]: σ ¼ neμ, where μ is the electron mobility. As a result, the improved electron injection efficiency and electrical conductivity of TiO2 nanorods synergistically lead to substantially-enhanced electron transport and collection, and thus increased Jsc and PCE of DSSCs. However, Nb doping level as high as 2 mol% might induce a high defect concentration or Nb–Nb pairings [53] in the TiO2 nanorods, which would decrease the electron mobility and thus the conductivity, enabling a degenerated device Jsc and PCE. EIS are also carried out to study the charge transport and recombi­ nation in DSSCs. Fig. 5(h) shows the EIS Nyquist plots of DSSCs using undoped and 1 mol% Nb doped TiO2 NRAs, and the inset is the used equivalent circuit, where Rseries represents the series resistance of the whole device, R1 and CPE1 stand for the charge transfer resistance and the interfacial capacitance at the electrolyte/Pt counter electrode interface, and R2 and CPE2 donate the recombination resistance and the chemical capacitance at the TiO2/dye/electrolyte interface respectively. As shown in Fig. 5(h), there exist two semicircles in the high-frequency (>1 kHz) and medium-frequency (1–100 Hz) areas, resulting from charge recombination at the electrolyte/Pt and TiO2/dye/electrolyte interfaces, respectively. The microscopic electrical parameters fitted are shown in Table 2. Rseries is obviously decreased from 14.8 Ohm to 10.4 Ohm after Nb doping, which is co-contributed by the improved electron injection efficiency, enhanced electrical conductivity and reduced recombination. Besides, after Nb doping R2 is significantly increased from 30 Ohm to 52.4 Ohm, indicating that the recombination across the photoanode/electrolyte interface is hindered. The electron lifetime constant (τ) can be determined by the equation: τ ¼ R2 � CPE2 . As shown in Table 2, τ is remarkably prolonged from 56.9 ms to 373 ms, an outstanding 6.5-fold enhancement upon Nb doping, which is ascribed to the improved electron transport and collection.

Table 2 Fitted electrical parameters derived from EIS. Device

Rseries/Ohm

R1/Ohm

CPE1/μF

R2/Ohm

CPE2/μF

τ/ms

undoped Nb doped

14.8 10.4

2.5 1.1

1.9 54.4

30.0 52.4

1894 7120

56.9 373

accompanying with change of crystalline orientations by using TiCl4 as the precursor. Compositional deviations and occurrence of low oxidative states of Ti3þ are also observed in the situations of heavy doping. Three types of emerging solar cells are fabricated to probe the Nb doping ef­ fects including PSCs, DSSCs and Cu2O based heterojunction devices. In all cases, Nb doping substantially improves the device performances. Specifically, in PSCs, the PCE enhancement of ~16% upon 0.1 mol % Nb incorporation is contributed by the increase of Voc and FF, resulted from reduced recombination at the perovskite/TiO2 interfaces, which is induced by better interfacial band alignments, reduced defects on TiO2 surface and improved charge transport property of nanorods. In DSSCs, 1 mol% Nb doping substantially prolongs electron diffusion lengths through TiO2 NRAs from ~14 μm to ~50 μm, leading to a sharp Jsc increase from 9.83 mAcm 2 to 14.43 mAcm 2 and thus a remarkable PCE enhancement of ~33%, which is due to a synergistic effect of increased electron injection, improved electron transport and reduced recombination. This study demonstrates that Nb doped hydrothermal TiO2 NRAs could be versatile electron transporting materials in various solar cells, and might find more applications in a diversity of areas such as PEC water splitting, sensors and energy storage devices. Declaration of competing interest All authors have no competing interests. Acknowledgements This work was financially supported by NATIONAL R&D PROGRAM of CHINA under Grant No. 2017YFA0207400, National Natural Science Foundation of China under Grant No. 11604250, 61634005, 21701129, 61704128, Natural Science Foundation of Shaanxi, China under Grant No. 2018JQ2014, 2018JQ6039, and the Fundamental Research Funds for the Central Universities under Grant No. JB181404, JB181408. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2020.227715. References [1] N.S. Lewis, D.G. Nocera, Proc. Natl. Acad. Sci. U.S.A. 103 (2006) 15729–15735. [2] A. Shah, P. Torres, R. Tscharner, N. Wyrsch, H. Keppner, Science 285 (1999) 692–698. [3] K. Yoshikawa, H. Kawasaki, W. Yoshida, T. Irie, K. Konishi, K. Nakano, T. Uto, D. Adachi, M. Kanematsu, H. Uzu, K. Yamamoto, Nat. Energy 2 (2017). [4] N.J. Jeon, J.H. Noh, Y.C. Kim, W.S. Yang, S. Ryu, S.I. Seok, Nat. Mater. 13 (2014) 897–903. [5] H. Zhou, Q. Chen, G. Li, S. Luo, T.-b. Song, H.-S. Duan, Z. Hong, J. You, Y. Liu, Y. Yang, Science 345 (2014) 542–546. [6] M. Liu, M.B. Johnston, H.J. Snaith, Nature 501 (2013) 395–þ. [7] D. Bi, B. Xu, P. Gao, L. Sun, M. Graetzel, A. Hagfeldt, Nano Energy 23 (2016) 138–144. [8] M. Kim, G.-H. Kim, T.K. Lee, I.W. Choi, H.W. Choi, Y. Jo, Y.J. Yoon, J.W. Kim, J. Lee, D. Huh, H. Lee, S.K. Kwak, J.Y. Kim, D.S. Kim, Joule. [9] U. Bach, D. Lupo, P. Comte, J.E. Moser, F. Weissortel, J. Salbeck, H. Spreitzer, M. Gratzel, Nature 395 (1998) 583–585. [10] S. Mathew, A. Yella, P. Gao, R. Humphry-Baker, B.F.E. Curchod, N. Ashari-Astani, I. Tavernelli, U. Rothlisberger, M.K. Nazeeruddin, M. Graetzel, Nat. Chem. 6 (2014) 242–247. [11] A. Yella, H.-W. Lee, H.N. Tsao, C. Yi, A.K. Chandiran, M.K. Nazeeruddin, E.W.G. Diau, C.-Y. Yeh, S.M. Zakeeruddin, M. Graetzel, Science 334 (2011) 629–634. [12] Z. Pan, I. Mora-Sero, Q. Shen, H. Zhang, Y. Li, K. Zhao, J. Wang, X. Zhong, J. Bisquert, J. Am. Chem. Soc. 136 (2014) 9203–9210.

4. Conclusion In summary, Nb doping is discovered to greatly affect the morpho­ logical, crystalline, compositional, optical and electronic properties of the TiO2 NRAs on the FTO substrates synthesized by a hydrothermal method. By using TBOT as the reaction precursor, light Nb doping causes obvious reduction of sizes of the nanorods and slight decrease of the band gap of TiO2, while the rod shapes, crystalline orientations and the chemical environment of the TiO2 nanorods remain stable. Besides, light Nb doping induces reduced defective oxides, increased electron con­ centration and shift of the energy band structure of TiO2 (i.e. ECB moving downwards by ~0.2 eV and EF getting closer to ECB). Furthermore, heavy Nb doping transforms thick NRAs into dense thin films, 9

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