Substantially Improved Na-Ion Storage Capability by Nanostructured Organic–Inorganic Polyaniline-TiO2 Composite Electrodes

Developing sodium (Na)-ion batteries is highly appealing because they offer the potential to be made from raw materials, which hold the promise to be less expensive, less toxic, and at the same time more abundant compared to state-of-the-art lithium (Li)-ion batteries. In this work, the Na-ion storage capability of nanostructured organic–inorganic polyaniline (PANI) titanium dioxide (TiO2) composite electrodes is studied. Self-organized, carbon-coated, and oxygen-deficient anatase TiO2–x-C nanotubes (NTs) are fabricated by a facile one-step anodic oxidation process followed by annealing at high temperatures in an argon–acetylene mixture. Subsequent electropolymerization of a thin film of PANI results in the fabrication of highly conductive and well-ordered, nanostructured organic–inorganic polyaniline-TiO2 composite electrodes. As a result, the PANI-coated TiO2–x-C NT composite electrodes exhibit higher Na storage capacities, significantly better capacity retention, advanced rate capability, and better Coulombic efficiencies compared to PANI-coated Ti metal and uncoated TiO2–x-C NTs for all current rates (C-rates) investigated.


■ INTRODUCTION
Over the last 30 years, the demand for lithium-ion batteries (LIBs) for powering a variety of applications, from handheld consumer electronics to power-demanding electric vehicles, has been constantly growing. 1 Most LIBs today employ transition-metal oxides, mainly LiCoO 2 , or LiFePO 4 as cathode-active materials, which, despite their apparent success, pose severe concerns for future large-scale energy storage. They are toxic, lack sustainability, and are related to a high carbon footprint upon production and subsequent recycling. 2 In combination with potential shortage in supply and subsequent increase in the price of Li, research on sodiumion batteries (SIBs) has rapidly gained momentum, as corroborated by a strong increase in the number of publications related to SIBs within the last 5−10 years. 3,4 Organic electrode materials are interesting candidates for next-generation environmentally benign battery materials, since they are less toxic, less limited, and can easily be recycled compared to commonly employed transition-metal oxides like LiCoO 2 or LiFePO 4 . In addition, organic electrode materials are expected to require less energy to be produced and consequently lower the battery-associated carbon footprint. 5−7 Organic p-type polymers that contain moieties, which can be reversibly reduced or oxidized, are particularly promising and are typically employed as cathode materials. 8−10 Most p-type polymers though possess limited electronic conductivity and are therefore prone to slow electrode kinetics and subsequent limited rate capabilities. Subsequently high cycling stabilities are challenging to achieve due to deterioration processes of the electrode material itself or of side reactions in combination with the electrolyte. 11 Polyaniline (PANI) represents a special case of conjugated polymers that can be doped by acid−base chemistry. In PANI, the transition from an insulator to a metal-like organic conductor occurs through a protonation-induced change in the π−electron system. A protonation via acid−base chemistry allows an internal redox reaction and thereby the transformation from "emeraldine base" with semiconductor properties to the metal-like "emeraldine salt" form of polyaniline. 12 MacDiarmid and his colleagues showed already in 1987 that it is possible to construct a rechargeable LIB employing the emeraldine base form of PANI as the cathode. Their rechargeable LIB showed already very promising characteristics such as a capacity around 148 mAh g −1 and an energy density of about 340 mWh g −1 . 13 More recently, PANI-coated nanostructures have also been investigated for their potential applications in supercapacitors. 14−16 The rate capabilities of organic polymers are intimately connected with the electron-and ion-transfer rates to the active redox units that can be reversibly reduced or oxidized. Therefore, issues regarding electrode kinetics may be circumvented by employing substrate nanomaterials in combination with carbon coating. 17−19 The carbon coating adds the required electronic conductivity to different nanostructures, while at the same time their reduced size significantly decreases the diffusion path for both ions and electrons.  Titanium dioxide (TiO 2 ) forms highly ordered, selforganized TiO 2 nanotubes (NTs) by simple anodic oxidation of the substrate Ti metal using a fluoride-containing electrolyte. 20 This facile, one-step method has been extensively investigated, and its mechanism is well characterized and understood. 21−23 The aligned pore structure of the TiO 2 NTs supports conduction of electrons and ions in one dimension, while the thin NT walls significantly shorten the solid-state ion diffusion path. Additionally, TiO 2 NTs show a high tolerance toward structural changes. It has been shown that TiO 2 NTs are capable of reversibly accommodating variations in volume that may occur upon sodiation and desodiation cycling when covered by a redox polymer capable of reversible Na-ion storage. 24 No conductive binder is necessary for connecting the TiO 2 nanotube array to the electrodes, rendering the system less complex compared to powder-based electrodes. The low complexity makes this approach particularly interesting and creates new prospects for basic research. The influence of thermal annealing of TiO 2 NTs in a reductive atmosphere has been previously reported by our group aiming at the improvement of LIB performance. 25−28 We were able to show that thermal annealing under an argon (Ar) atmosphere at 400°C led to a conversion of the as-grown amorphous TiO 2 NTs to oxygen-deficient TiO 2−x NTs with an anatase crystal structure. 29 The presence of oxygen vacancies (x < 2) is known to enhance the electronic conductivity of the TiO 2 NTs substantially, compared to stoichiometric TiO 2 , and as a consequence, the charge transfer increases due to lower ohmic drops at the Ti/TiO 2 interface. Furthermore, oxygen deficiency is known to support the phase transition that occurs upon Liion intercalation and deintercalation in TiO 2 . This enables a significantly better Li-ion battery performance compared to stoichiometric anatase NTs regarding intercalation capacity and rate capability. 25,30−32 The high-temperature annealing in an argon−acetylene gas mixture leads to the conversion of the TiO 2−x phase to a more conductive, carbon-coated TiO 2−x -C phase, which significantly enhances the electron and ion transport inside the nanotube array. 29 We herein investigate the new concept of combining these self-organized and anodically grown TiO 2−x -C NTs with electropolymerized PANI to form nanostructured organic− inorganic polyaniline-TiO 2 composite electrodes for use in SIBs.

■ RESULTS AND DISCUSSION
Well-aligned TiO 2 nanotube arrays are grown in situ, directly on a Ti metal substrate by anodization using the prescribed method. Top-view scanning electron microscopic (SEM) images of TiO 2−x -C NTs before PANI deposition are shown in Figure 1. The SEM top view displays a homogeneous NT array for TiO 2−x -C NTs with top tube diameters of about 165 nm and top wall thicknesses of about 15 nm (Figure 1c). In this form, the carbon-coated anatase TiO 2−x -C NT offers an ideal template, characterized by its high active surface area, important to foster good adhesion of the PANI-coated layers to the subjacent supporting substrate.
The electrodeposition of aniline on self-organized TiO 2−x -C NTs is schematically illustrated in Figure 2a. The deposition is initiated by a potentiodynamic electropolymerization from 1.2 to 0 V for 10 cycles with a scan rate of 50 mV s −1 in a 0.35 M H 2 SO 4 aqueous solution containing a 0.4 M aniline monomer. Sulfuric acid has been previously identified to be the most suitable medium for producing a uniform PANI-coated film on different metal oxides. 33 A clear peak current increase can be observed with an increasing number of cycles, indicating the development of a PANI-covered film onto the TiO 2−x -C NT surface (shown by red arrows in Figure 2b). Subsequent galvanostatic electropolymerization is applied at a constant current of 0.1 mA cm −2 for 45 min (Figure 2c), showing the typical initial steep increase in potential up to around 0.85 V until its stabilization potential around 0.8 V is established. 34 Figure 2d shows a comparison between the initial TiO 2−x -C NTs before (black line) and after (red line) electropolymerization. A substantial increase in the current response is observed, showing a more boxlike behavior superimposed by broad, faradic charge-transfer peaks. These peaks have only little separation in their anodic and cathodic peak potentials, suggesting only small polarization between the oxidized and reduced states. The broad, rectangular-like shape of the cyclic voltammogram (CV) response, after PANI deposition, may be described by a pseudocapacitive behavior, which is characteristic for an electrode surface covered with a thin film of organic conducting polymer films. 35 SEM images of the final PANI-coated TiO 2−x -C NTs reveal a clear coverage of the NT surface by a PANI film, leading to a decreased inner diameter and an increased wall thickness of about 75 nm ( Figure 1d). The surface morphology of the PANI-coated TiO 2−x -C NT composites is following the subjacent NT surface with some granular structure on top that is typical for electropolymerized PANI. 36 For comparison, PANI-coated flat Ti-substrates and pure mechanically polished flat Ti-substrates are shown in Figure S1. Electrodeposited PANI shows a characteristic granular structure. 37 The formation of a PANI surface film is supported by impedance measurements. The Nyquist plot for the pristine TiO 2−x -C NT electrode (recorded at open circuit potential OCP; Figure 2e, black squares) shows a large, not fully developed semicircle, with a charge-transfer resistance (R ct ) of 129.86 Ω cm 2 (57.46 Ω). Differently, the PANI-coated TiO 2−x -C NT electrode shows a significantly smaller and almost fully developed semicircle, with an R ct of 4.16 Ω cm 2 (1.84 Ω; Figure 2e, red circles). Accordingly, it can be inferred that a well-conductive PANI coating in the emeraldine salt form has been formed. The low-frequency branch for pristine TiO 2−x -C NTs is characterized through a slightly sloping line indicating the development of a second time constant, typically observed for nanostructured electrodes, where the diffusion is, in one dimension, bounded by a large porous electrode with capacitive walls. 38 Differently, for the PANI-coated TiO 2−x -C NTs, the low-frequency branch in the Nyquist plot develops in an almost vertical line, indicating a finite space diffusion behavior where a flooded porous layer is present, 39 in our case the freshly formed PANI film, which is terminated by a blocking outer interface of the PANI/electrolyte boundary. 40 This impedance response is expected for electrodes coated by a thin film of conducting polymer like PANI. Levi and Aurbach have thoroughly investigated the electrochemical behavior of thin-film polymer-coated electrodes for LIB applications. 39,41 By applying the developed finite space model from Levi and Aurbach, with its corresponding electrochemical equivalent circuit and a slightly adopted restricted diffusion element (shown in the Supporting Information Figure S2), the equivalent circuit fits well the measured impedance data (compare the Supporting Information Figure S3 and Table S1). The electronic equivalent circuit consists of two interfaces for the PANI film/TiO 2−x -C NT ACS Applied Energy Materials www.acsaem.org Article interface and the PANI film/solution interface. Since the impedance data is recorded at open circuit voltage (OCV), it is expected that electronic equilibrium occurs across the PANI film/TiO 2−x -C NT interface while simultaneously ionic equilibrium is established across the PANI film/solution interface. The two interfaces, modeled by R/C (R ct e of 4.16 Ω cm 2 , C dl e of 34.35 μF) and R/CPE (R ct i of 3.43 Ω cm 2 , C dl i of 0.75 μF) elements, are therefore placed in series to the modified Warburg element for adopted restricted diffusion. One additional R element (R el of 24.10 Ω cm 2 ) is used, describing the electrolyte resistance that accounts for the impedance response at high frequencies. The deviation-related weighted sum of squares (X 2 /Z) between the measured impedance response and the equivalent circuit fit for both the initial and final PANI-coated TiO 2−x -C NTs is low. The deviation-related weighted sum of squares (X 2 /Z) values are given with 1.13 × 10 −3 and 0.97 × 10 −3 , respectively, indicating that the equivalent circuit adequately describes the experimentally measured Nyquist plot. One has to point out also that for the initial TiO 2−x -C NTs without a PANI film present, the restricted diffusion element is not needed to fit the recorded impedance data. Additionally, the assumption of an

ACS Applied Energy Materials
www.acsaem.org Article electronic equilibrium occurring simultaneously across the interfaces may change under an applied potential. Especially in proximity to the redox peak potentials (around 0.5 V; Figure  2d), it is expected that space charge layers (evolving fast) may be countered by ionic concentration gradients that evolve slow in comparison.
A full summary of all parameters used for fitting together with their obtained values is given in the Supporting Information Table S1. Figure 3 shows SEM cross-section images of TiO 2−x -C NTs before PANI deposition (initial) and PANI-coated TiO 2−x -C NTs (final). Analyzing the SEM cross sections reveals 4.0 ± 0.3 μm average tube length, an average pore diameter of 165 ± 12 nm at the top, and a solid hemisphere with a diameter of 225 ± 30 nm at the bottom of each nanotube.
While the TiO 2−x -C NTs are characterized by splintered, sharp edges at the top rims (Figure 3a,c), the PANI-coated TiO 2−x -C NT composite shows distinguishable morphological features. Its surface becomes rough, and the nanotube wall thickness increases to 75 ± 10 nm, while the initial TiO 2−x -C NT geometry is not altered by the electrodeposition of PANI. The magnified SEM cross-section image (Figure 3d) shows that the PANI layer provides good coverage of the topmost tube surfaces. The film thickness determined by SEM imaging is additionally confirmed by the measured mass increase upon electropolymerization. A mass increase of 0.53 ± 0.15 mg was measured at a total active surface area of about 295 ± 65 cm 2 . The value for the active surface area is obtained by a mathematical model, using the NT geometry analyzed via SEM imaging and additionally verified by the capacitive current contribution in the CV measurement. Values of 230 and 360 cm 2 were obtained, respectively, which are in reasonable agreement. This infers that the entire NT surface is electrochemically active toward PANI deposition, suggesting the complete coverage of the NT surface by PANI. In summary, SEM top-view and cross-section images in combination with the values obtained for the active surface area and mass increase upon PANI electropolymerization suggest that the outer and inner surfaces of the TiO 2−x -C NTs are coated with a 20−30 nm thin PANI film. The PANI percentage in the composite electrode material is determined by about 31% (compare the experimental part).
X-ray photoelectron spectroscopy (XPS) is a truly important tool allowing a quantitative analysis of the electrode materials and their corresponding redox states. 42 Figure 4a depicts the survey XPS spectra of pristine and PANI-coated TiO 2−x -C NTs. The N 1s signal of the TiO 2−x -C NTs prior to PANI deposition is practically zero, while the Ti 2p, Ti 2s, and the O 1s signals, characteristic for TiO 2 , are clearly visible. For the PANI-coated TiO 2−x -C NTs (Figure 4a), the Ti 2s and Ti 2p signals are absent, while a strong peak for the N 1s signal is measured, corroborating the deposition of PANI on the surface of the TiO 2−x -C NTs. High-resolution spectra measured at the Ti 2p, C 1s as well as N 1s regions are shown in Figure 4b−e. The atomic concentrations of the different elements present obtained from the high-resolution spectra are summarized in the Supporting Information Table S2.
The high-resolution C 1s spectra for the pristine TiO 2−x -C NTs (Figure 4e) reveal the presence of carbon deposited at the surface through a TiO 2 /C ratio of about 77:22 atom %, originating mainly from NT annealing in Ar/C 2 H 2 during material synthesis with some possible carbon contributions adsorbed from ambient air. The main carbon component is present in the wwwC−C sp 3 state with a characteristic binding energy of 284.4 eV. For PANI-coated TiO 2−x -C NTs, the C/N ratio obtained from the high-resolution C 1s and N 1s signal (Figure 4c,d) is about 75:13 atom % (Table S1), which is fairly close to the theoretical 6:1 ratio for pure PANI. In the properly  Table S3 in the Supporting Information. Figure 5a shows the comparison of CV measurements for the PANI-coated Ti metal (dashed green), TiO 2−x -C NTs before PANI deposition (black), and PANI-coated TiO 2−x -C NTs (red) at a scan rate of 10 mV s −1 . The CV measurements of TiO 2−x -C NTs do not show any distinct reduction or oxidation peaks that can be correlated toward reversible intercalation of Na ions into TiO 2−x -C within the given potential range 2.0−4.0 V (Figure 5a). While TiO 2 NTs have been demonstrated to serve as good model electrodes for studying LIBs, 47−49 the exact Na-ion storage mechanism in TiO 2 NTs is still controversially discussed and, up to now, not fully understood. 24,50,51 Previous studies revealed that the measured current response upon voltage change is characterized by a simultaneous capacitive and insertion contribution. 24 The Na-ion insertion part takes place only well below 2 V (i.e., at ∼0.7/0.9 V at a scan rate of 10 mV s −1 compare the Supporting Information Figure S6). The Na-ion insertion at low potentials is known to represent the reversible storage of Na ions in the active material structure. This is followed by reduction of Ti 4+ (sodiation) and vice versa by oxidation of Ti 3+ (desodiation). 52 Therefore, within the given potential range 2.0−4.0 V, TiO 2−x -C NTs can be regarded as an inactive but important support material, offering an aligned, homogeneous pore structure; one-dimensional electronic and ionic conduction; and short solid-state Na-ion diffusion pathways accompanied by a high tolerance regarding structural alterations, and are consequently well qualified as support materials for nanostructured organic−inorganic polyaniline-TiO 2 composite electrodes. The current response from the PANI-coated Ti metal (Figure 5a, dashed green line) is substantially smaller compared to the PANI-coated TiO 2−x -C NTs, which results mainly from the lower active surface area. When the current response is multiplied by a factor of 10 ( Figure S7), a reductive peak is clearly visible but substantially shifted toward lower potentials (by about 410 mV−2.28 V compared to 2.69 V for PANI-coated TiO 2−x -C NTs). This may be explained by the unfavorable band positions of TiO 2 . The polished Ti metal is covered by a thin oxide layer. The lower edge of the conduction and the upper edge of the valence band are situated at about −4 and −7.3 eV vs the vacuum energy level (about −0.5 and +2.8 V vs NHE) for TiO 2 . 53,54 As it is expected that the band edges are pinned by surface states and will not be shifted upon external bias, electron injection into PANI is largely hindered, which is corroborated by the polarization overpotential of about 410 mV, measured in the CV current response. Figure 5b shows the Nyquist plots for TiO 2−x -C NTs before PANI deposition (black) and PANI-coated TiO 2−x -C NTs (red) at OCP in the battery half-cell using Na-metal foil as the counter electrode and 1 M NaFSI/EC:DMC as the electrolyte system. The Nyquist plots of both electrodes are characterized by a high-frequency semicircle and a low-frequency branch. For the TiO 2−x -C NTs, the low-frequency branch in the Nyquist plot is dominated by a larger imaginary part and immediately develops in an almost vertical line. This characteristic is in accordance with minute currents on the CV at these potentials. For the PANI/TiO 2−x -C NTs, the lowfrequency branch initially develops in a slightly sloping line and turns afterward to an almost vertical straight line. This impedance response is again characteristic for thin-film polymer-coated electrodes and can be fitted to the already previously introduced electronic equivalent circuit, depicted in Figure S2 and the comparison between the measured data and the best fit in Figure S3 (Nyquist and Bode plots). The values obtained are given by an R ct e of 2.19 Ω cm 2 (0.97 Ω), a C dl e of 11.3 μF, an R ct i of 133.58 Ω cm 2 (59.06 Ω), a C dl i of 3.35 μF, and an R el of 9.90 Ω cm 2 (4.38 Ω, Table S1). The deviationrelated weighted sum of squares (X 2 /Z) between the measured impedance response and the equivalent circuit fit is low with 2.0 × 10 −3 . Figure 5c,d shows the galvanostatic sodiation/desodiation cycling (GCPL) measurements of PANI-coated TiO 2−x -C NTs (red) and TiO 2−x -C NTs before PANI deposition (black) at different currents of 16, 80, 160, 800, and 1600 μA. When calculated to the amount of only the PANI material present (0.53 ± 0.15 mg) in the PANI-coated TiO 2−x -C NTs (red), these currents correspond to C-rates of about C/10, C/2, 1C, 5C, and 10C. The charge/discharge measurements of PANIcoated TiO 2−x -C NTs are described by the downward-sloping line featuring three distinct different regions: first, starting from 3.5 V to about 2.6 V, with a slightly downward-sloping trend; second, followed by a rapid decrease in potential from 2.6 V to about 2.2 V; and finally, from 2.2 V onward to 2.0 V again with a more gently downward-sloping trend. The GCPL measurements are congruent with the current-potential response recorded in the CV measurements (Figure 5a). When cycled between 3.5 and 2.0 V, with a current of 16 μA, the PANIcoated TiO 2−x -C NTs exhibit a specific capacity of 22.65 μAh cm −2 in the first and 21.05 μAh cm −2 in the second cycle. Even at 100 times higher constant current of 1600 μA, the PANIcoated TiO 2−x -C NTs still reveal a specific capacity of 6.27 ACS Applied Energy Materials www.acsaem.org Article μAh cm −2 . These characteristics are distinctly different compared to TiO 2−x -C NTs without a PANI coverage ( Figure  5d). The pure TiO 2−x -C NTs show a steep and monotone downward-sloping line with a specific capacity of 9.57 μAh cm −2 in the first and 7.23 μAh cm −2 in the second cycle. At a constant current of 1600 μA, the specific capacity drops below 1 μAh cm −2 . Flat electrodes of the PANI-coated Ti metal in comparison do not exhibit significant specific sodiation capacities with only around 0.52 μAh cm −2 in the first cycle and negligible specific capacities at higher current rates ( Figure  S8b). The reaction of Na ions with PANI following initial electropolymerization may be described according to the three-step sequence shown in Figure 6a−c. 13 After electropolymerization in a 0.35 M H 2 SO 4 aqueous solution, the PANI is in the emeraldine salt form (Figure 6a), where the acid protonates the imine nitrogen located within the polymer backbone and thereby induces charge carriers. After the electrode is transferred into the battery half-cell, the NaFSI electrolyte will absorb the protons, transferring PANI, to a large extent, from the emeraldine salt into the emeraldine base form (Figure 6b). This transformation is corroborated by an increase in the charge-transfer resistance in the corresponding Nyquist plot at OCV (Figure 5b). Subsequently, upon repeated sodiation and desodiation in the battery electrolyte, the PANI may transform from its emeraldine base to its leucoemeraldine base form ( Figure 6c) and back, respectively. According to the reaction shown in Figure 6, PANI in the emeraldine salt form has a theoretical specific capacity of about 148 mAh g −1 if the PANI weight only is considered and about 97 mAh g −1 when the weight of the HSO 4 − anion from the electropolymerization is included (or 74 mAh g −1 when the HSO 4 − anion is replaced by the heavier FSI − anion in the battery electrolyte). A detailed analysis of the specific capacities and corresponding Coulombic efficiencies of our PANI-coated Ti metal, TiO 2−x -C NTs, and PANI-coated TiO 2−x -C NTs, as a function of cycle number and applied current, is shown in Figure 7.
Both systems, the TiO 2−x -C NTs (Figure 7, black) and PANI-coated TiO 2−x -C NTs (Figure 7, red), undergo an initial capacity decay at a low sodiation current of 16 μA (C/10 rate based on PANI mass loading). The cycling trend for both electrodes is characterized by a more gradual capacity decay obtained for the PANI-coated TiO 2−x -C NTs compared to the TiO 2−x -C NTs. The specific capacity for PANI-coated TiO 2−x -C NTs drops from initially 22.65 μAh cm −2 (96 ± 28 mAh g −1 ) in the first cycle to about 15 μAh cm −2 after 10 galvanostatic sodiation/desodiation cycles at a current of 16 μA and tends to stabilize afterward, with 15.25 μAh cm −2 (65 ± 19 mAh g −1 ) in the 15th cycle. This initial capacity loss amounts to about 32% of the electrodes' first sodiation capacity. The specific gravimetric capacity of 65 ± 19 mAh g −1 in the 15th cycle still amounts to about 88% of the theoretical maximum capacity of PANI, including the FSI anion. In comparison to the PANI-coated TiO 2−x -C NTs, the specific capacity for TiO 2−x -C NTs drops from initially 9.57 μAh cm −2 in the first cycle to 3.12 μAh cm −2 after 15 galvanostatic sodiation/desodiation cycles at a current of 16 μA, corresponding to a capacity loss of about 67% of the electrodes' initial capacity. If we further take into account that the absolute capacity decrease for PANI-coated TiO 2−x -C NTs is 7.4 μAh cm −2 compared to 6.45 μAh cm −2 for the pure TiO 2−x -C NTs, then only 0.95 μAh cm −2 or 4.2% of the initial capacity decrease is related to the PANI material itself and the rest must be attributed to electrolyte decomposition and/or side reactions. This is consistent with galvanostatic cycling measurements of flat, PANI-coated Ti metal electrodes ( Figure  7, green) where the initial capacity decay is on the same order (about 0.4 μAh cm −2 ). The corresponding Coulombic efficiency of the galvanostatic sodiation desodiation cycling is shown in Figure 7b. PANI-coated TiO 2−x -C NTs exhibit higher Coulombic efficiencies compared to PANI-coated Ti metal and uncoated TiO 2−x -C NTs for all C-rates. In general, the Coulombic efficiencies are almost 100% at elevated Crates, while at very slow rates (C/10), Coulombic efficiencies below 100% are mainly related to potential electrolyte decomposition and/or side reactions. Nyquist impedance plots and corresponding ex situ ATR−FTIR measurements after the galvanostatic cycling measurements shown in Figure 7 also confirm the excellent PANI film integrity (Figures S5c and  S9) on the TiO 2−x -C NT support.  Although an extensive comparison study of the available literature regarding polymer cathode materials for Na-ion batteries, including specific material and battery-related properties, is beyond the scope of this research article, the performance of PANI-coated TiO 2−x -C NTs is comparable to other, recent reports on advanced polymer composite cathode materials. A summary of the most relevant literature, comparing the operation voltages vs specific gravimetric capacities after 100 charge/discharge cycles of different Naion organic and organic composite electrodes, is shown in Figure 8. 55−63 For example, a poly(N-vinylcarbazole) polymer electrode has been reported as a dual-intercalation cathode for Na-ion batteries, which demonstrated specific discharge capacities of about 110 mAh g −1 for 100 cycles. 55 An aluminum-coordinated poly(tetrahydroxybenzoquinone) electrode demonstrated a reversible capacity of 113 mAh g −1 and stable cycle performance over 100 cycles. 56 Graphene-wrapped poly 2,5-dihydroxy-1,4-benzoquinone-3,6-methylene nanocomposites characterized by three-dimensional nanoflowerlike structures recently showed stable specific capacities of about 121 mAh g −1 after 100 cycles. 57 Na 2 FePO 4 F and Na 2.4 V 2 (PO 4 ) 3 encircled by a nanolayer composed of poly(3,4 ethylenedioxythiophene) have been reported with stable cycling performance (over 500 cycles) and specific capacities of 123.1 and 112.4 mAh g −1 , respectively. 58,59 A good review article, summarizing the most recent progress in advanced organic electrode materials and comparing the performance of different organic polymers as cathode materials for Na-ion batteries, is that by Zhao et al. 64 To further investigate potential changes in the NT crystal structure during galvanostatic cycling, ex situ X-ray diffraction (XRD) patterns regarding the pristine and PANI-coated TiO 2−x -C NTs are measured before and after the GCPL measurements and are shown in Figure 7. The XRD patterns for both the TiO 2−x -C NTs and PANI-coated TiO 2−x -C NTs are summarized by intense peaks at 2θ of ∼25, 38, 48, 54, and 56°corresponding to (101), (004), (200), (105), and (211) of the anatase structure ( Figure 9). 47,65 XRD measurements show that after repeated electrochemical sodiation between 3.5 and 2.0 V, no Na-related new phase is detected in the XRD patterns. Therefore, TiO 2−x -C NTs maintain their initial anatase structure, which suggests that the sodiation process for PANI-coated TiO 2−x -C NTs does not change the crystallographic structure of the subjacent NTs significantly. This further corroborates that Na-ion storage does not happen in the TiO 2−x -C NT host structure within the given potential range, since previous studies by our group and others have clearly shown that Na-ion storage in anatase TiO 2 is followed by a continuous intensity decrease of the anatase diffraction peaks upon sodiation. 24,52 ■ CONCLUSIONS The good cycling performance of PANI-coated TiO 2−x -C NTs (Figure 7), together with its well-maintained anatase NT structure after long-term galvanostatic cycling (Figure 9), proves that the carbon-coated anatase TiO 2−x -C NT array provides an ideal nanostructured solid template for PANI deposition. The large available surface area of the NT array also appears to be favorable for good adhesion regarding the PANI coating to the supporting substrate, showing little material dissolution upon repeated sodiation/desodiation measurements. In summary, PANI-coated TiO 2−x -C NTs exhibit higher Na storage capacities, better capacity retention, superior rate capabilities, and higher Coulombic efficiencies compared to PANI-coated Ti metal and uncoated TiO 2−x -C NTs for all C-rates. As for other cathode and/or anode materials, PANI-coated TiO 2−x -C NTs require, in the asprepared state, presodiation (or prelithiation) before it can be used in a full-cell battery. 66−68 However, its half-cell battery performance will prove to be a good alternative to common cathode materials employed for LIBs if high electron-and iontransfer rates can be realized, for example, by employing substrate nanomaterials like TiO 2−x -C NTs where nanostructures are coated by a thin layer of conductive carbon. PANI on TiO 2−x -C NTs may therefore present a cost-effective, abundant, and environmentally benign cathode material for future rechargeable SIBs.

■ EXPERIMENTAL SECTION
Synthesis of the TiO 2−x -C NTs has been performed by a slightly modified procedure reported previously. 25,48 A summary of the synthesis procedure and a photograph of the amorphous TiO 2 NT samples before and anatase TiO 2−x -C NTs samples after carbothermal annealing ( Figure S11a,b) can be found in the Supporting Information. Electropolymerization of PANI onto TiO 2−x -C NTs and Ti metal discs has been carried out in an aqueous solution containing 0.4 M aniline (>99.5%, Sigma-Aldrich, used as received) in 0.35 M H 2 SO 4 (Merck, used as received), employing a three-electrode system with TiO 2−x -C NTs as the working electrode, a Pt plate as the counter electrode, and a mercury-mercurous sulfate (Hg/Hg 2 SO 4 ) reference electrode. Electropolymerization of PANI is performed in a two-step process, starting with 10 cycles of potentiodynamic polymerization between 1.2 V vs NHE and 0.0 V vs NHE at 50 mV s −1 and followed by a constant current polymerization at 0.1 mA cm −2 for 45 min. A photograph of a pristine polished Ti metal disk, Ti metal disk after 10 cycles of potentiodynamic polymerization, and after further constant current polymerization is shown in the Supporting Information, Figure  S11c.
For defining the TiO 2−x -C NT volume, their morphology has been evaluated by analyzing the SEM images (average tube length: 4.0 ± 0.3 μm; 165 ± 12 nm average pore diameter at the top, and a solid hemisphere at the bottom of each nanotube 225 ± 30 nm; Figure  S10). The total volume of the TiO 2−x -C NTs is calculated according to the NT morphology, given by a cone-shaped NT configuration. Each NT is closed by a solid hemisphere at the bottom. With this approach, the TiO 2−x -C volume has been calculated to approx. 0.445 mm 3 . The resultant active mass of the TiO 2−x -C NTs on the electrode has been calculated by multiplying the TiO 2−x -C volume by the density of anatase (3.84 g cm −3 ) to 1708.7 μg. Therefore, the PANI percentage in the composite is about 31%.
Battery half-cell measurements were carried out using a commercially available battery cell (ECC-Ref Cell from EL-Cell) in a three-electrode configuration. Details are given in the Supporting Information. The mass specific capacity in Figure 7 is determined by dividing the measured capacity in mAh by the active electrode material mass, and hence the mass of PANI, being 0.53 ± 0.15 mg (the areal capacity has been determined accordingly, by dividing the measured capacity in mAh by the electrodes' geometrically exposed area of 2.26 cm 2 ).
SEM imaging was performed with a JEOL JSM-7601F field emission electron microscope. Secondary electrons were detected by an in-lens detector to characterize the surface and cross-sectional morphology prior to and after PANI deposition. An acceleration voltage of 10 kV was chosen. Cross sections were simply prepared by scratching the sample and imaging the fraction of the NTs.
XPS (Thermo MultiLab 2000 spectrometer equipped with a hemispherical energy analyzer) has been utilized to examine the chemical composition of the PANI-coated TiO 2−x -C NTs. Additionally, XPS was employed to determine the oxidation states of the individual material components prior to and also after PANI electrodeposition. The XPS is measured using a take-off angle of zero with respect to the surface normal. Furthermore, a monochromatized Al Kα X-ray source (1486.6 eV) was used. The step size for the high-resolution spectra of the Ti 2p, C 1s, and N 1s regions was 0.02 eV with a dwell time of 0.2 s. The quantitative analysis of the XPS measurements was performed with CasaXPS software (version 2.3.16). A charging correction has been employed for all spectra by shifting the spectra with respect to the location of the Ti(IV) 2p 3/2 peak of TiO 2 at 458.6 eV.
Infrared spectra (diamond ATR, PIKE GaldiATR, Bruker Vertex 70) were measured in the range of 4000−400 cm −1 with a resolution of 2 cm −1 (32 scans per spectrum). Raman spectra (Bruker BRAVO) were recorded with a resolution of 2 cm −1 from 160 to 3200 cm −1 . The spectrometer suppresses the fluorescence by the patented SSE technology (patent number: US8570507B1) and has two temperature-shifted excitation lasers (DuoLaser) with wavelengths of probably 785 and 1064 nm (obscured by the manufacturer). IR and Raman spectra were processed in OPUS spectroscopic software.
X-ray powder diffraction (XRD) analysis was performed on a Siemens D5000 X-ray diffractometer, with Cu Kα emission. Diffractograms were measured between 15 and 75°(2θ) and a step size of 0.02°(2θ). The acquisition time for each step was 1 s.

* sı Supporting Information
The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acsaem.9b02541. SEM images of mechanically polished Ti metal substrates and final PANI-coated mechanically polished Ti metal substrates ( Figure S1); equivalent electric circuit used for fitting all of the EIS data ( Figure S2); Nyquist and Bode plots for initial TiO 2−x -C NTs before PANI deposition and PANI-coated TiO 2−x -C NTs ( Figure S3); summary of all EIS fitting parameters (Table S1); atomic concentrations of the different elements obtained from XPS analysis (Table S2); Raman spectra of PANI-coated TiO 2−x -C NTs, TiO 2−x -C NTs before PANI deposition, and PANIcoated, mechanically polished Ti metal substrates ( Figure S4); ATR−FTIR spectra of TiO 2−x -C NTs before PANI deposition, PANI-coated TiO 2−x -C NTs before battery cycling, and PANI-coated TiO 2−x -C NTs after 150 galvanostatic charge/discharge battery cycles ( Figure S5); summary of the evaluation of the most prominent, characteristic vibrational modes for PANIcoated TiO 2−x -C NTs obtained from Raman and ATR− FTIR measurements (Table S3); CVs of TiO 2−x -C NTs between 3.0 and 0.1 V in a Na-ion battery half-cell ( Figure S6); CV measurements of PANI-coated Ti metal compared to PANI-coated TiO 2−x -C NTs ( Figure  S7); galvanostatic charge/discharge performances (Figure S9); Nyquist plots of PANI-coated TiO 2−x -C NTs, TiO 2−x -C NTs, and PANI-coated, mechanically polished Ti metal substrates before and after galvanostatic charge/discharge cycling ( Figure S9); cross-section SEM images used for determining the average tube length and tube bottom diameter ( Figure S10); photographs of different working electrodes used in this study ( Figure S11); summary on the synthesis of TiO 2−x -C NTs (PDF) Julia Kunze-Liebhaüser − Institute of Physical Chemistry, University of Innsbruck, A-6020 Innsbruck, Austria; orcid.org/0000-0002-8225-3110 Complete contact information is available at: https://pubs.acs.org/10.1021/acsaem.9b02541

Author Contributions
The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.