- Email: [email protected]
Hierarchically porous titania xerogel monoliths: synthesis, characterization and electrochemical properties
Materials Research Bulletin 73 (2016) 48–55
Contents lists available at ScienceDirect
Materials Research Bulletin journal homepage: www.elsevier.com/locate/matresbu
Hierarchically porous titania xerogel monoliths: synthesis, characterization and electrochemical properties Wenjun Zhu, Hui Yang, Yuan Xie, Sai Sun, Xingzhong Guo* School of Materials Science and Engineering, Zhejiang University, 38 Zheda Road, Xihu District, Hangzhou 310027, China
A R T I C L E I N F O
A B S T R A C T
Article history: Received 1 March 2015 Received in revised form 17 August 2015 Accepted 21 August 2015 Available online xxx
Hierarchically porous titania (TiO2) xerogel monoliths were successfully synthesized by a facile sol–gel process accompanied by phase separation in the presence of formamide (FA) and triblock copolymer F127, followed by ambient pressure drying to remove the solvents. Gelation of the system has been mediated by FA, while F127 enhances a phase-separation. The resultant dried xerogel presents a hierarchically porous structure constructed by large macropores and mesoporous skeletons, which exhibits a high BET surface area of 444 m2 g 1. Heat-treatment at 800 C results in the formation of the crystalline anatase phase with high thermal stability, without spoiling the macrostructure. The asprepared porous TiO2 xerogel displays superior electrochemical properties, and the specific capacity of TiO2 xerogel could remain over 114 mAh g 1 (69 mAh cm 3) at 1 C after 100 cycles, representing a good cycling stability. ã 2015 Elsevier Ltd. All rights reserved.
Keywords: Oxides Sol–gel chemistry Electrochemical properties Microstructure Energy storage
1. Introduction Lithium-ion batteries (LIBs) have attracted considerable interests as a priority power source in portable electronic devices and electric/hybrid vehicles due to the advantages of high energy density, long life cycle, low toxicity and self-discharge rate [1–6]. However, the commercial carbon anodes suffer several limitations for practical vehicle applications. Firstly, the dendritic lithium deposits on the electrode surface because of the low lithium insertion potential, which results in serious safety concerns. Secondly, the solid electrolyte interface (SEI) layer formed at low potential (0 V) is an electronic insulator, which restrains the high rate performance of electrode to a large extent. Thirdly, the conventional carbon materials display a poor performance at fast discharge-charge cycling resulting from low lithium diffusion coefficients [7,8]. Thus, it is highly desirable to explore safe and high performance anode materials of LIBs for practical applications in electric vehicles and renewable energy storage. Nowadays, titanium dioxide (TiO2) has received increasing attention as a promising anode material because of its chemical stability, low cost, and environmental benignity [9–14]. Moreover, compared with the traditional carbon anode, TiO2 anode delivers a relatively higher lithium insertion/extraction voltage (higher than 1.5 V vs Li+/Li), which can ensure adequate safety by efficiently
* Corresponding author. Fax: +86 571 87953054. E-mail address: [email protected] (X. Guo). http://dx.doi.org/10.1016/j.materresbull.2015.08.025 0025-5408/ ã 2015 Elsevier Ltd. All rights reserved.
avoid the formation of SEI layers and lithium dendrites on the electrode [15–17]. Nevertheless, the practical application of TiO2 is still largely restrained by the limited rate capability resulting from its low intrinsic electrical conductivity (10 12 S cm 1) [18,19]. Recently, considerable efforts have been made to enhance the rate performance of TiO2 electrodes. Porous materials have played important roles in improving the electrochemical performances by shortening the diffusion paths of both electrons and lithium ions and ensuring better contact between the electrode material and the electrolyte through large specific surface area. Furthermore, a large number of active sites can be provided for inserting Li ions and facilitating the transfer of Li ions due to open pore channels, especially the mesoporous, leading to high rate capability [7,9,11,19–23]. Shao et al. [7] reported a 3D porous architecture composed of TiO2 nanotubes connected with a carbon nanofiber matrix for lithium-ion battery, displaying excellent rate performance (112 mA h g 1 at 30 C rate). Zapien et al. [21] fabricated porous and dense TiO2 nanospheres for lithium-ion battery applications, and the porous nanospheres present superior electrochemical performance. Xerogels (aerogels) have garnered particular interests in the applications of LIBs due to the unique porous structure with an interconnected structure and particular properties such as extremely low density and thermal insulating [24–26]. Liu et al. [27] synthesized Fe3O4/C composites from Fe-based xerogels for anode materials of LIBs, presenting enhanced electrochemical properties. Hao et al. [28] have reported Co3O4/Carbon aerogel hybrids as LIBs anode materials, and a high specific discharge
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
capacity of 779 mAh g 1 could be obtained after 50 cycles. However, there have been few reports on the electrochemical performance of TiO2 xerogel. In this work, we demonstrate a facile strategy to prepare hierarchically porous TiO2 xerogel monolith through a sol–gel process accompanied by phase separation, followed by ambient pressure drying to remove the solvents. The addition of formamide (FA) mediates the gelation of the system, while the triblock copolymer F127 induces the phase separation. The hierarchically porous structure constructed by large macropores and mesoporous skeletons can be obtained, and the electrochemical properties of the resultant xerogel monoliths were firstly evaluated in detail. 2. Experimental 2.1. Synthesis of xerogel monolith Tetrabutyl titanate (Ti(OC4H9)4, TBOT, Sinopharm Chemical Reagent Co., Ltd (China)) was used as titanium source. Mixtures of distilled water (H2O) and ethanol (EtOH, Sinopharm Chemical Reagent Co., Ltd (China), 99.5%) were used as the solvent. Ethylic acid (HAc, Sinopharm Chemical Reagent Co., Ltd (China)) was used as catalyst as well as chelating agent. Formamide (FA, Shanghai Lingfeng Chemical Reagent Co., Ltd) was used to initiate gelation as well as a drying control chemical additive (DCCA). Tetraethoxysilane (TEOS, Sinopharm Chemical Reagent Co., Ltd (China)), ethanol (EtOH) and n-hexane were used to solvent exchange. Triblock copolymer F127 (Shanghai Lingfeng Chemical Reagent Co., Ltd) was used as a phase separation inducer. All the chemical reagents were used as received. The preparation process of TiO2 xerogel is shown in Fig. 1, and the typical synthesis is described as follows. 5 mL TBOT was dissolved in the 10 mL EtOH, and the mixture was named as solution A. 1.5 mL HAc, 0.092–0.552 g F127, 1.5 mL deionized water and 10 mL EtOH were mixed, and the mixture was named as solution B. The solution B was dropped into the solution A by 40– 50 drops per minute under a strong stirring condition. After continuously stirring for 20 min, 0.3 mL FA was then added to the mixed solutions under ambient conditions (25 C). The container was sealed and kept at 60 C for gelation. After gelation, the obtained wet gel was aged at 60 C for 48 h. The aged gels were solvent exchanged in EtOH at 60 C for 24 h three times, in TEOS/ EtOH (TEOS: EtOH = 1:4, volume ratio) solution at 60 C for 18 h twice, in n-hexane at 60 C for 48 h. Some of the dried gels were heat-treated at various temperatures between 500 and 900 C for 2 h in air. In this study, the molar ratios of TBOT: EtOH: HAc: water:
49
FA: F127 are designed to be 1: 24: 1.8: 6: 0.6: x (x = 0.0005, 0.001, 0.0015, 0.002, 0.0025 and 0.003), which were named as F05, F10, F15, F20, F25 and F30. The sample without adding F127 was named as F00. 2.2. Characterization of xerogels Apparent density was determined by measuring the volume (V) and mass (m) of xerogel samples. The xerogel samples with different amount of F127 have been cut into regular geometry to measure the volume and mass. Porosity was calculated by using apparent density (ra) and true density (rt, 3.9 g cm 3) as follows: porosity = (rt ra)/rt 100%. Differential thermal analysis (DTA) was performed using a WRT-3P analyzer. The crystalline phases of xerogels were analyzed by the X-ray diffraction (XRD) on a Rigaku. D/Max-RA X-ray diffractometer using nickel filtered Cu-Ka radiation in the range of 2u is 10–80 with a scanning speed of 2 /min. The morphology and microstructure were examined by scanning electron microscope (SEM, FEI SIRION, 25 KV) and transmission electron microscopy (TEM; FEI, Tecnai G2 F20). The nitrogen adsorption–desorption of xerogels were characterized by a nitrogen adsorption–desorption apparatus (Quantachrome Autosorb-1-C). Before each nitrogen adsorption–desorption measurement, the samples were degassed at 200 C under vacuum for more than 6 h. Specific surface areas were calculated using the Brunauer–Emmett–Teller (BET), meanwhile, the pore size distributions were calculated using the Barrett–Joyner–Halenda (BJH) method based on the desorption branch. 2.3. Electrochemical measurements Electrochemical experiments were carried out using standard CR 2025 type coin cells. The working electrode was prepared by mixing 70 wt % the active material (TiO2 xerogel powder), 20 wt % carbon black (Super P) and 10 wt % polyvinylidene fluoride (PVDF) binder. The slurry was pasted on a Cu foil and dried in a vacuum oven at 120 C for 12 h to serve as the working electrode. The weight of the active material in the electrode sheet was about 3 mg cm 2. Cells were assembled in in a glove box under highpurity argon atmosphere with a pure lithium foil as both counter electrode and reference electrode. The electrolyte solution was 1 M LiPF6 dissolved in a mixture of ethylene carbonate (EC)/dimethyl carbonate (DMC) at a volume ratio of 1:1. Cellgard 2300 polypropylene microporous film was used as the separator. Galvanostatical charge–discharge experiments were carried out in a voltage range of 1.0–3.0 V on a battery test system (Shenzhen Neware Battery,
Fig. 1. Schematic illustration of the fabrication process of hierarchically porous TiO2 xerogel monoliths.
50
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
China). Cyclic voltammetry (CV) measurements were conducted using a CHI650B electrochemical workstation (Shanghai Chenhua, China) at a scan rate of 0.1 mV s 1 in the potential range of 1.0–3.0 V and the impedance spectra were tested on the same workstation with the frequency ranging from 0.1 Hz to 100 KHz. 3. Results and discussion 3.1. Synthesis and characterization of the TiO2 xerogel monolith The synthesis and characterization of the TiO2 xerogel monolith were investigated in detail. During the reaction process, an alcogelgel with a spatial network structure is formed due to the hydrolysis and polycondensation of tetrabutyl titanate, consisting of titanium-oxygen based (RTi O TiR) polymers. The gelation time of TiO2 sols is related to the content of F127. As shown in Fig. 2a, with the increase of F127 content, the gelation time of TiO2 sol increases firstly (see F05 sample), and then rapidly decreases with the further increase of F127. F00 sample shows a gelation time of 18 min, while the gelation time of F30 sample is only about 6 min. It indicates that the gelation of TiO2 sol could be accelerated by the addition of F127. And white precipitations are observed in the experiment with the molar ratio of 0.0035 (F127:TBOT). Triblock copolymer F127 possesses the characteristic of relatively hydrophilic polyethylene oxide (PEO) block and relatively hydrophobic polypropylene oxide (PPO) block [29,30]. Additionally, F127 has a long molecular chain with a big molecular weight (12,600 g mol 1). It indicates that F127 has more significant impact on the gelation performance of TiO2 sol and the construction of gel network. As is well known, F127 presents two modification functions: wrapping and crosslinking [30]. The modification of wrapping restricts the single particle to contact or aggregate, while the modification of crosslinking accelerates the formation of much huger clusters. That is to say, when the added amount of the F127 is small, it is difficult for TiO2 sol particles to polycondensation due to the steric hindrance derived from wrapping, resulting in a long gelation time (as see F05 sample in Fig. 2a). The crosslinking function of F127 becomes dominant gradually with the increase of the added amount, and the gelation time of TiO2 sol reduces due to the strong polycondensation and the increase in amount and diameter of clusters. Fig. 2b displays the apparent densities and porosities of TiO2 xerogel with different F127 contents. It can be seen that the apparent density of TiO2 xerogel obviously decreases with the increase of F127 content, resulting in rapid increase of the
corresponding porosity. The minimum apparent density of 0.261 g 3 and the corresponding high porosity of 93.3% could be obtained at the molar ratio of 0.003 (F127:TBOT, sample F30). The result confirms that the addition of F127 could effectively promote the generation of porous structure, which leads to the decrease of the apparent density and the increase of the corresponding porosity for prepared TiO2 xerogel. Scanning electron microscopy (SEM) was undertaken to study the morphology of TiO2 xerogels. Fig. 3 shows the SEM images of fracture surface of TiO2 xerogels with different F127 contents. It is observed that only particulate xerogel with a lot of nanopores is obtained without F127 (F00). Compared with the F00 sample, the F05 xerogel sample displays a uniform morphology with smaller pore size and uniform pore size distribution, which is attributed to the steric hindrance of F127 as above mentioned. With the increase of F127 content, the macropores gradually form, and the particles aggregate to generate skeletons. Finally, the macroporous morphology with macropores and loose skeletons can be obtained with the mole ratio of F127:TBOT at 0.0025 (F25), and the xerogel exists in the form of monolith, as shown in Fig. 3 h. With the further increase of the F127 content (F30), the macropores become very large with relatively dense skeletons, and the xerogel presents granular shape again (no shown). The formation of unique porous structure can be attributed to the crosslinking function of F127 as above mentioned. During the process, F127 offers more entanglement with the skeleton particles, resulting in the networks increasingly loose, and the pores size rapidly increases to form macropores. In addition, the hydrophilic EO units of F127 are connected with surface hydroxyls of the gel networks via hydrogen bonding during and after the gelation. F127 will be away from the solvent phase and exposed toward the gel network through hydrophobic interaction of PO units [29,30]. The polycondensation of TiO2 sol particles after modification will increase the incompatibility of the mixed system, which results in the phase separation and the formation of macroporous structure [30]. Nitrogen isothermal absorption–desorption measurements were performed to determine the BET surface area and the porosity of the typical TiO2 xerogels. As revealed in Fig. 4. The isotherms of F00 and F25 samples present type IV with H3 hysteresis loops at relative pressure of 0.7–1.0, signifying the existence of slit-shaped pores [26,31,32]. While the isotherm of F30 sample is changed into type II, indicating that the formation of macroporous structure. It can be detected from the BJH pore size distributions (Fig. 4b), the major pores are distributed in large
cm
Fig. 2. (a) The gelation time of TiO2 sols with different F127 contents. (b) Apparent densities and porosities of TiO2 xerogels with different F127 contents.
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
51
Fig. 3. Morphology of fracture surface of TiO2 xerogels with different F127:TBOT ratios (a-without F127(F00), b-F127:TBOT = 0.0005(F05), c-F127:TBOT = 0.001(F10), d-F127: TBOT = 0.0015(F15), e-F127:TBOT = 0.002(F20), f-F127:TBOT=0.0025(F25), g-F127:TBOT = 0.003(F30), h-xerogel monolith of F25 sample).
(a)
(b)
Fig. 4. N2 adsorption–desorption isotherms (a) and pore size distribution of the typical TiO2 xerogels (b).
meso- to macropore region. And the pore size distribution of the F25 sample is relatively sharper than those of F00 and F30 samples. The variation of the size and shape of pores may be associated with the phase separation induced by F127. The obtained specific parameters of the BET surface area and median pore size are shown in Table 1, the dried gel without surfactant possesses a surface area
of 415 m2 g 1 with most pores at around 8 nm and a whole pore volume of 0.86 cm3 g 1. The addition of surfactant (F25 sample) increases surface area to 444 m2 g 1, and deceases the pore size to 7.0 nm and the pore volume to 0.82 cm3 g 1, respectively. With the further increases in surfactant (F30 sample), the surface area decreases to 397 m2 g 1, and the pore size and volume rapidly
Table 1 Pore characteristics of typical titania xerogels with different F127 contents. Sample
Specific surface area/m2 g
F00 F25 F30
415 444 397
1
Total pore volume/cm3 g 0.86 0.82 0.25
1
Average pore size/nm 8 7 3
52
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
decrease to 3 nm and 0.25 cm3 g 1 respectively, which is attributed to the formation of relatively dense skeletons. Fig. 5a shows the DTA curve of monolithic TiO2 xerogel (F25 sample). Two endothermic peaks near 90 and 225 C are mainly due to the volatilization of adsorbed water and the fracturing of -OH group, and a small exothermic peak at 305 C is attributed to the pyrolysis of organic species [33]. The very small exothermic peak locates at 665 C is possibly ascribed to the crystallization of anatase phase. Stengl et al. [34] reported that an exotherm appears at 420 C for anatase crystallization in titania gel gained by low temperature supercritical drying. In contrast, the high crystallization temperature in Fig. 5b indicates that the crystallization of anatase from titania gel is restricted to some extent, which could be attributed to the formation of macropore structure derived from phase separation [22]. Moreover, the small exothermic peak (305 C) corresponding to the pyrolysis of organic species indicates that the formation of interconnected macropores is beneficial to the removing of organic species during the solvent exchanging. The heat-treatment of the typical monolithic TiO2 xerogel monolith (F25 sample) was carried out at 500–900 C with a heating rate of 1–2 C/min. Fig. 5b shows XRD patterns of typical TiO2 xerogels at different heat-treatment temperatures. No discernible peaks can be detected after the heat-treated at 500 C. Very small peaks appear at 700 C, which is attributed to anatase phase, implying that amorphous structure begins to transform into anatase phase at this temperature. It is also seen that the diffraction peaks become strong with the increase of temperature. The diffraction peaks ascribed to anatase phase appear completely at 800 C, and these diffraction peaks show a characteristic of very broad width due to the formation of nanoscale anatase crystallites. When the heat-treatment temperature increases to 900 C, the anatase phase with broad diffraction peaks keep steadily, and there are no other phases can be detected, implying that the anatase phase has high thermal stability of phase transformation. In general, the transition to anatase of TiO2 xerogel happens at 420 C, and the transition of anatase-rutile occurs at 600–700 C [34,35]. It also confirms that the crystallization of anatase phase and the transition of anatase-rutile are restricted, which could be attributed to the unique porous structure. The result is consistent with the results of DTA (see Fig. 5a). Furthermore, the density of the F25 samples after heat-treatment is larger than that of dried xerogel (Fig. 2b) due to the volume shrinkage. After calcination at 500, 700, 800 and 900 C, the calculated densities are 0.47, 0.49, 0.61 and 0.60 g cm 3, and the
corresponding porosities are 87.9, 87.4, 84.3 and 84.6%, respectively. The increase in density with heat-treatment temperature can be ascribed to the decrease of porosity derived from the growth and agglomeration of grains. After heat treatment, the crystalline TiO2 xerogel can keep the monolithic shape basically and have a volume shrinkage of about 10 % compared with the dried xerogel. Fig. 6 shows the SEM image and N2 adsorption–desorption isotherm of TiO2 xerogel heattreated at 800 C. Although some agglomerations exist, the mophology is basically maintained after heat-treatment, implying that the formation of crystalline phase does not spoil the morphology of xerogel (Fig. 6a). In contrast to the macrostructure, the micro-mesoporous structure is significantly altered by the heat-treatment. Fig. 6b presents the isotherm of the xerogel heattreated at 800 C with a typical type II according to the IUPAC classification, representing a non-porous or macroporous structure. Compared with the dried gel (Fig. 4a), the specific surface area of the heat-treated gels (Fig. 6b) decreases rapidly (from 444 m2 g 1 to 61 m2 g 1). Moreover, the number of the pores, espically micro–mesopores decreases and the xerogel becomes much denser. After the heat-treatment at 800 C, the pores are basically distributed in large meso- to macropore region with micropores disappeared. The big variation of surface area and pore size after the heat-treatments is mainly attributed to the formation of anatase phase and the sintering effect of titania particles. The detail microstructure of TiO2 aerogel (F25 calcined at 800 C) was further investigated by TEM and HRTEM analysis. As shown in Fig. 7, the aerogels were constructed by nanoparticles with dozens of nanometers in size, and the nanoparticles interconnect into a network porous structure (Fig. 7a). It can be observed from the HRTEM image (Fig. 7b) that the lattice interplanar spacings of approximately 0.35 and 0.237 nm correspond to the (1 0 1) and (0 0 4) plane of tetragonal anatase TiO2, respectively, in good agreement with the results of XRD (Fig. 5). 3.2. Electrochemical performance It is reported that the morphology and microstructure of materials play an important role in the final electrochemical performance. In this work, the F25 sample shows a uniform morphology with macropores and loose skeletons, in comparison with other samples. Based on the previous works [9,19,20,36,37], it indicates that the F25 sample possesses a surperior electrochemical performance. Therefore, we investigate the electrochemical performance of the F25 xerogel sample after calcination by a series
Fig. 5. (a) DTA curve of TiO2 xerogel monolith (F25 sample). (b) XRD patterns of typical TiO2 xerogel (F25 sample) at different heat-treatment temperatures.
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
53
Fig. 6. Morphology (a) and N2 adsorption–desorption isotherm (b) of F25 xerogel monolith heat-treated at 800 C (inset is pore size distribution).
Fig. 7. Low (a) and high (b) magnification TEM image of the resulting TiO2 aerogel (F25 calcined at 800 C).
of electrochemical measurements, and the different calcination temperatures have been considered. As is well known, volumic capacity is an important parameter to evaluate the electrochemical performance of porous materials for the practical application [38,39]. In this work, the capacities of the samples are addressed not only in weight dimensionality (mAh g 1) but also by volume (mAh cm 3). The initial three charge–discharge profiles at 0.5 C in the voltage range of 1.0–3.0 V are presented in Fig. 8a. There are no obvious voltage plateau observed during the lithiation/delithiation processes, which can be ascribed to the solid solution behavior or the relative low crystallinity of the nanocomposite electrode, in accordance with previous reports on nanosized TiO2-based Li-ion batteries [40– 42]. In the first discharge step, a large specific capacity of 280 mAh g 1 (170 mAh cm 3) is obtained, indicating that 0.8 mol of Li ions are incorporated into the TiO2 xerogel electrode (Li0.8TiO2). The high discharge specific capacity in the first cycle can be possibly derived from an undesirable side reaction with the electrolyte or an irreversible change in the structure of the TiO2 xerogel upon deeper Li-ion insertion [40,43]. It is generally believed that the formation of SEI film happens when discharged to 0 V. In consideration of the high voltage plateau of TiO2 (1.5 V), the galvanostatical charge–discharge experiments were carried out in the voltage range of 1.0-3.0 V. Thus, the formation of the SEI layer during the electrochemical process could be effectively avoided. The large irreversible capacity has been observed in the first cycle could be attributed to the irreversible electrochemical decomposition of electrolyte or impurity phase over the TiO2 xerogel surface due to the porous structure [44,45]. Although a large irreversible capacity of about 130 mAh g 1 appears in the first cycle, the specific capacities of the TiO2
xerogel electrode are stabilized on following cycles, and the discharge and charge capacities in the third cycle are 140 and 135 mAh g 1, corresponding to 85 and 82 mAh cm 3, respectively. Moreover, sloped behaviors after the first discharge reaction are observed, which can be ascribed to the irreversible formation of a “nanocomposite” of crystalline grains and amorphous regions [46,47]. Fig. 8b shows the CVs of the TiO2 xerogel electrode at a scan rate of 0.1 mV s 1. The result is consistent with charge–discharge curves in Fig. 8a, there are no obvious intense peak detected during the scan process. Two broad current peaks are observed at around 1.7 V (cathodic sweep) and 2.0 V (anodic sweep), corresponding to the biphasic transition between tetragonal anatase and orthorhombic lithium titanate. The charge–discharge capacities versus cycle number for the F25 samples calcined at different temperture are shown in Fig. 8c. It can be seen that the reversible capacity of the F25 sample increases with the heat-treatment temperature in the range of 500–800 C, which is attributed to the increase of crystallinity. When the temperture increases to 900 C, the sample has a bad electrochemical performance, which might results from the collapse of porous structure. In contrast, the TiO2 xerogel electrode calcined at 800 C presents the best cycle performance, which delivers an initial discharge capacity of 180 mAh g 1 (109 mAh cm 3) and charge capacity of 135 mAh g 1 (82 mAh cm 3) with a coulombic efficiency of 75%. The coulombic efficiency increases upon cycling, which reaches to around 100% after the 10th cycle and keep unchanged during the whole cycles. Meanwhile, accompanied with the cycle number increasing, the reversible capacity decreases in the initial 10 cycles and then gradually levels off. It can still remain above 114 mAh g 1 (69 mAh cm 3) after 100 cycles, showing a low average capacity
54
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
Fig. 8. (a) Initial three charge–discharge curves (at 0.5 C, 1 C = 168 mAh g 1) and (b) CVs at a scan rate of 0.1 mV s 1of TiO2 xerogel (F25 calcined at 800 C). (c) Cycling performance (at 1 C) of TiO2 xerogels (F25) calcined at different temperatures. (d) Nyquist plots and corresponding simulation results of TiO2 xerogel (F25 calcined at 800 C), the inset is the equivalent circuit for plot fitting.
fading of 0.19% per cycle from the second cycle to 100th cycle. Furthermore, the reversible capacities of the samples calcined at 500, 700 and 900 C are 69, 76 and 67 mAh g 1, respectively, and the corresponding volumic capacities are 32, 37 and 40 mAh cm 3, respectively. Although the smallest porosity it has, the TiO2 xerogel calcined at 800 C presents the highest capacity both in mAh g 1 and mAh cm 3. The high charge–discharge capacities and superior cycle stability are due to the synergistic effect of crystallinity and hierarchically porous structure, guaranteeing the electrochemical
reaction and ensuring better contact between the electrode material and the electrolyte. Electrochemical impedance spectroscopy (EIS) was performed for the TiO2 xerogel electrode before and after the galvanostatic cycling and the corresponding Nyquist plots are shown in Fig. 8d. It can be seen that all the Nyquist plots exhibit the characteristic of a semicircle in the high frequency region and a sloping straight line in the low frequency region. The semicircle in high frequency is ascribed to the charge-transfer impedance at the electrolyte/
Fig. 9. Schematic diagram of the insertion/extraction process of the hierarchically porous TiO2 xerogel monoliths.
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
electrode interface, and the sloping straight line of low frequency region corresponds to the lithium ion diffusion process in the electrodes [31,48]. The inset shows the equivalent circuit for the TiO2 xerogel to simulate the experimental EIS plots, and the corresponding fitting curves are drawn in Fig. 8d, respectively. Rs is attributed to the solution resistance (including total ohmic resistance of the separator, electrolyte, and contacts between the current collector and the electrode materials), CPE (i) and Rsl (i) (i = 1, 2) represent the capacity of the layer and the migration of hydrogen ions, respectively. Rct and Cdl denote the charge-transfer resistance and a double-layer capacitance, respectively. ZW is the Warburg impedance, representing the diffusion of electrolyte ions into the electrode materials, which is associated with the inclined line at low frequencies. As indicated in Fig. 8d, a small semicircle was obtained before the cycling, indicating low resistances of Li+ diffusion in charge-transfer reaction. After cycling, the semicircle is enlarged, which suggests the resistance of Li+ diffusion increases upon cycles (Fig. 9). This is probably due to the large volume change during charge–discharge cycling that gradually deteriorates the contact between TiO2 xerogel and the collector, which resulting in the 20 % fade of capacity during the specific capacities versus cycle test in Fig. 8c. 4. Conclusions In summary, monolithic titania (TiO2) xerogels were synthesized by combining a sol–gel process with phase separation, followed by ambient pressure drying to remove the solvents. The as-prepared titania xerogel monolith exhibits an unique hierarchically porous structure constructed by large macropores and mesoporous skeletons. The monolithic shape and morphology of xerogels are not spoiled after heat-treatment, the BET surface area rapidly decreases from 444 to 61 m2 g 1 because of the formation and sintering of anatase nanoparticles. The crystallization of anatase phase and the transition of anatase-rutile are restricted due to the unique porous structure. The TiO2 xerogel calcined at 800 C presents the best electrochemical performance, the reversible capacity could remain over 114 mAh g 1 (69 mAh cm 3) at 1 C after 100 cycles with a good cycling stability. The route provides a facile strategy for the preparation of hierarchically macroporous xerogel monoliths, and we believe the hierarchically porous titania (TiO2) xerogel monolith will also extend its application fields such as catalysis, separation, absorption and energy conversion. Acknowledgements This work is supported by the National Natural Science Foundation of China (51372225), and Zhejiang Provincial Natural Science Foundation of China (LY13B010001). References [1] J.M. Tarascon, M. Armand, Nature 414 (2001) 359. [2] Y. Xia, Z. Xiao, X. Dou, H. Huang, X.H. Lu, R.J. Yan, Y.P. Gan, W.J. Zhu, J.P. Tu, W.K. Zhang, X.Y. Tao, ACS Nano 7 (2013) 7083. [3] L.X. Yuan, Z.H. Wang, W.X. Zhang, X.L. Hu, J.T. Chen, Y.H. Huang, J.B. Goodenough, Energy Environ. Sci. 4 (2011) 269. [4] H. Huang, W. Zhu, X. Tao, Y. Xia, Z. Yu, J. Fang, Y. Gan, W. Zhang, ACS Appl. Mater. Interface 4 (2012) 5974.
55
[5] W. Zhu, H. Huang, W. Zhang, X. Tao, Y. Gan, Y. Xia, H. Yang, X. Guo, Electrochim. Acta 152 (2014) 286. [6] N. Brun, S.R.S. Prabaharan, C. Surcin, M. Morcrette, H. Deleuze, M. Birot, O. Babot, M.F. Achard, R. Backov, J. Phys. Chem. C 116 (2012) 1408. [7] B.T. Zhao, S.M. Jiang, C. Su, R. Cai, R. Ran, M.O. Tade, Z.P. Shao, J. Mater. Chem. A 1 (2013) 12310. [8] X. Xin, X.F. Zhou, J.H. Wu, X.Y. Yao, Z.P. Liu, ACS Nano 6 (2012) 11035. [9] Y.G. Guo, Y.S. Hu, W. Sigle, J. Maier, Adv. Mater. 19 (2007) 2087. [10] D. Dambournet, I. Belharouak, K. Amine, Chem. Mater. 22 (2010) 1173. [11] N. Li, G. Liu, C. Zhen, F. Li, L.L. Zhang, H.M. Cheng, Adv. Funct. Mater. 21 (2011) 1717. [12] P. Kubiaka, J. Geserick, N. Husing, A. Wohfahrt-Mehrens, J. Power Sources 175 (2008) 510. [13] F. Carn, A. Colin, M.F. Achard, H. Deleuze, C. Sanchez, R. Backov, Adv. Mater. 17 (2005) 62. [14] F. Carn, M.F. Achard, O. Babot, H. Deleuze, S. Reculusa, R. Backov, J. Mater. Chem. 15 (2005) 3887. [15] J.W. Zhang, X.X. Yan, J.W. Zhang, W. Cai, Z.S. Wu, Z.J. Zhang, J. Power Sources 198 (2012) 223. [16] S.T. Myung, N. Takahashi, S. Komaba, C.S. Yoon, Y.K. Sun, K. Amine, H. Yashiro, Adv. Funct. Mater. 21 (2011) 3231. [17] Y.-H. Ding, P. Zhang, H.-M. Ren, Q. Zhuo, Z.-M. Yang, Y. Jiang, Mater Res Bull 46 (2011) 2403. [18] Z.X. Yang, G.D. Du, Q. Meng, Z.P. Guo, X.B. Yu, Z.X. Chen, T.L. Guo, R. Zeng, J. Mater. Chem. 22 (2012) 5848. [19] I. Moriguchi, R. Hidaka, H. Yamada, T. Kudo, H. Murakami, N. Nakashima, Adv Mater 18 (2006) 69. [20] V. Gentili, S. Brutti, L.J. Hardwick, A.R. Armstrong, S. Panero, P.G. Bruce, Chem Mater 24 (2012) 4468. [21] H.E. Wang, H. Cheng, C.P. Liu, X. Chen, Q.L. Jiang, Z.G. Lu, Y.Y. Li, C.Y. Chung, W.Y. Zhang, J.A. Zapien, L. Martinu, I. Bello, J. Power Sources 196 (2011) 6394. [22] G. Hasegawa, T. Sato, K. Kanamori, K. Nakanishi, T. Abe, New J Chem 38 (2014) 1380. [23] N. Brun, S.R.S. Prabaharan, M. Morcrette, C. Sanchez, G. Pecastaings, A. Derre, A. Soum, H. Deleuze, M. Birot, R. Backov, Adv. Funct. Mater. 19 (2009) 3136. [24] M. Zeller, V. Lorrmann, G. Reichenauer, M. Wiener, J. Pflaum, Adv Energy Mater 2 (2012) 598. [25] R.P. Maloney, H.J. Kim, J.S. Sakamoto, ACS Appl. Mater. Interface 4 (2012) 2318. [26] H. Yang, W.J. Zhu, S. Sun, X.Z. Guo, RSC Adv. 4 (2014) 32934. [27] X.Y. Chen, B.H. Liu, Z.P. Li, Solid State Ionics 261 (2014) 45. [28] F.B. Hao, Z.W. Zhang, L.W. Yin, ACS Appl. Mater. Interface 5 (2013) 8337. [29] R. Barreiro-Iglesias, C. Alvarez-Lorenzo, A. Concheiro, J. Control. Release 93 (2003) 319. [30] M. Kurahashi, K. Kanamori, K. Takeda, H. Kaji, K. Nakanishi, RSC Adv. 2 (2012) 7166. [31] W.J. Zhu, H. Huang, Y.P. Gan, X.Y. Tao, Y. Xia, W.K. Zhang, Electrochim. Acta 138 (2014) 376. [32] W.Y. Li, X.Z. Guo, Y. Zhu, Y. Hui, K. Kanamori, K. Nakanishi, J Sol–Gel Sci. Technol. 67 (2013) 639. [33] Z.Q. Zhao, X.L. Jiao, D.R. Chen, J. Mater. Chem. 19 (2009) 3078. [34] V. Stengl, S. Bakardjieva, J. Subrt, L. Szatmary, Microporous Mesoporous Mater. 91 (2006) 1. [35] M. Popa, D. Macovei, E. Indrea, I. Mercioniu, I.C. Popescu, V. Danciu, Microporous Mesoporous Mater. 132 (2010) 80. [36] W. Zhu, H. Yang, K. Nakanishi, K. Kanamori, X. Guo, RSC Adv. 5 (2015) 24803. [37] X.B. Wang, Y.Y. Wang, L. Yang, K. Wang, X.D. Lou, B.B. Cai, J. Power Sources 262 (2014) 72. [38] M. Depardieu, R. Janot, C. Sanchez, A. Bentaleb, R. Demir-Cakan, C. Gervais, M. Birot, M. Morcrette, R. Backov, J. Mater. Chem. A 2 (2014) 18047. [39] M. Depardieu, R. Janot, C. Sanchez, A. Bentaleb, C. Gervais, M. Birot, R. DemirCakan, R. Backov, M. Morcrette, RSC Adv. 4 (2014) 23971. [40] D.H. Kim, K.M. Min, K.S. Park, I.J. Park, I.S. Cho, W.M. Seong, D.W. Kim, K.S. Hong, Dalton Trans. 42 (2013) 4278. [41] K.S. Park, K.M. Min, Y.H. Jin, S.D. Seo, G.H. Lee, H.W. Shim, D.W. Kim, J. Mater. Chem. 22 (2012) 15981. [42] Z.S. Hong, M.D. Wei, T.B. Lan, L.L. Jiang, G.Z. Cao, Energy Environ. Sci. 5 (2012) 5408. [43] K.S. Park, J.G. Kang, Y.J. Choi, S. Lee, D.W. Kim, J.G. Park, Energy Environ. Sci. 4 (2011) 1796. [44] M. Rahman, J.Z. Wang, M.F. Hassan, D. Wexler, H.K. Liu, Adv. Energy Mater. 1 (2011) 212. [45] F. Wu, X. Li, Z. Wang, H. Guo, Nanoscale 5 (2013) 6936. [46] J. Jamnik, J. Maier, Phys. Chem. Chem. Phys. 5 (2003) 5215. [47] P. Balaya, H. Li, L. Kienle, J. Maier, Adv. Funct. Mater. 13 (2003) 621. [48] G.F. Cai, J.P. Tu, D. Zhou, J.H. Zhang, Q.Q. Xiong, X.Y. Zhao, X.L. Wang, C.D. Guts, J. Phys. Chem. C 117 (2013) 15967.
Contents lists available at ScienceDirect
Materials Research Bulletin journal homepage: www.elsevier.com/locate/matresbu
Hierarchically porous titania xerogel monoliths: synthesis, characterization and electrochemical properties Wenjun Zhu, Hui Yang, Yuan Xie, Sai Sun, Xingzhong Guo* School of Materials Science and Engineering, Zhejiang University, 38 Zheda Road, Xihu District, Hangzhou 310027, China
A R T I C L E I N F O
A B S T R A C T
Article history: Received 1 March 2015 Received in revised form 17 August 2015 Accepted 21 August 2015 Available online xxx
Hierarchically porous titania (TiO2) xerogel monoliths were successfully synthesized by a facile sol–gel process accompanied by phase separation in the presence of formamide (FA) and triblock copolymer F127, followed by ambient pressure drying to remove the solvents. Gelation of the system has been mediated by FA, while F127 enhances a phase-separation. The resultant dried xerogel presents a hierarchically porous structure constructed by large macropores and mesoporous skeletons, which exhibits a high BET surface area of 444 m2 g 1. Heat-treatment at 800 C results in the formation of the crystalline anatase phase with high thermal stability, without spoiling the macrostructure. The asprepared porous TiO2 xerogel displays superior electrochemical properties, and the specific capacity of TiO2 xerogel could remain over 114 mAh g 1 (69 mAh cm 3) at 1 C after 100 cycles, representing a good cycling stability. ã 2015 Elsevier Ltd. All rights reserved.
Keywords: Oxides Sol–gel chemistry Electrochemical properties Microstructure Energy storage
1. Introduction Lithium-ion batteries (LIBs) have attracted considerable interests as a priority power source in portable electronic devices and electric/hybrid vehicles due to the advantages of high energy density, long life cycle, low toxicity and self-discharge rate [1–6]. However, the commercial carbon anodes suffer several limitations for practical vehicle applications. Firstly, the dendritic lithium deposits on the electrode surface because of the low lithium insertion potential, which results in serious safety concerns. Secondly, the solid electrolyte interface (SEI) layer formed at low potential (0 V) is an electronic insulator, which restrains the high rate performance of electrode to a large extent. Thirdly, the conventional carbon materials display a poor performance at fast discharge-charge cycling resulting from low lithium diffusion coefficients [7,8]. Thus, it is highly desirable to explore safe and high performance anode materials of LIBs for practical applications in electric vehicles and renewable energy storage. Nowadays, titanium dioxide (TiO2) has received increasing attention as a promising anode material because of its chemical stability, low cost, and environmental benignity [9–14]. Moreover, compared with the traditional carbon anode, TiO2 anode delivers a relatively higher lithium insertion/extraction voltage (higher than 1.5 V vs Li+/Li), which can ensure adequate safety by efficiently
* Corresponding author. Fax: +86 571 87953054. E-mail address: [email protected] (X. Guo). http://dx.doi.org/10.1016/j.materresbull.2015.08.025 0025-5408/ ã 2015 Elsevier Ltd. All rights reserved.
avoid the formation of SEI layers and lithium dendrites on the electrode [15–17]. Nevertheless, the practical application of TiO2 is still largely restrained by the limited rate capability resulting from its low intrinsic electrical conductivity (10 12 S cm 1) [18,19]. Recently, considerable efforts have been made to enhance the rate performance of TiO2 electrodes. Porous materials have played important roles in improving the electrochemical performances by shortening the diffusion paths of both electrons and lithium ions and ensuring better contact between the electrode material and the electrolyte through large specific surface area. Furthermore, a large number of active sites can be provided for inserting Li ions and facilitating the transfer of Li ions due to open pore channels, especially the mesoporous, leading to high rate capability [7,9,11,19–23]. Shao et al. [7] reported a 3D porous architecture composed of TiO2 nanotubes connected with a carbon nanofiber matrix for lithium-ion battery, displaying excellent rate performance (112 mA h g 1 at 30 C rate). Zapien et al. [21] fabricated porous and dense TiO2 nanospheres for lithium-ion battery applications, and the porous nanospheres present superior electrochemical performance. Xerogels (aerogels) have garnered particular interests in the applications of LIBs due to the unique porous structure with an interconnected structure and particular properties such as extremely low density and thermal insulating [24–26]. Liu et al. [27] synthesized Fe3O4/C composites from Fe-based xerogels for anode materials of LIBs, presenting enhanced electrochemical properties. Hao et al. [28] have reported Co3O4/Carbon aerogel hybrids as LIBs anode materials, and a high specific discharge
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
capacity of 779 mAh g 1 could be obtained after 50 cycles. However, there have been few reports on the electrochemical performance of TiO2 xerogel. In this work, we demonstrate a facile strategy to prepare hierarchically porous TiO2 xerogel monolith through a sol–gel process accompanied by phase separation, followed by ambient pressure drying to remove the solvents. The addition of formamide (FA) mediates the gelation of the system, while the triblock copolymer F127 induces the phase separation. The hierarchically porous structure constructed by large macropores and mesoporous skeletons can be obtained, and the electrochemical properties of the resultant xerogel monoliths were firstly evaluated in detail. 2. Experimental 2.1. Synthesis of xerogel monolith Tetrabutyl titanate (Ti(OC4H9)4, TBOT, Sinopharm Chemical Reagent Co., Ltd (China)) was used as titanium source. Mixtures of distilled water (H2O) and ethanol (EtOH, Sinopharm Chemical Reagent Co., Ltd (China), 99.5%) were used as the solvent. Ethylic acid (HAc, Sinopharm Chemical Reagent Co., Ltd (China)) was used as catalyst as well as chelating agent. Formamide (FA, Shanghai Lingfeng Chemical Reagent Co., Ltd) was used to initiate gelation as well as a drying control chemical additive (DCCA). Tetraethoxysilane (TEOS, Sinopharm Chemical Reagent Co., Ltd (China)), ethanol (EtOH) and n-hexane were used to solvent exchange. Triblock copolymer F127 (Shanghai Lingfeng Chemical Reagent Co., Ltd) was used as a phase separation inducer. All the chemical reagents were used as received. The preparation process of TiO2 xerogel is shown in Fig. 1, and the typical synthesis is described as follows. 5 mL TBOT was dissolved in the 10 mL EtOH, and the mixture was named as solution A. 1.5 mL HAc, 0.092–0.552 g F127, 1.5 mL deionized water and 10 mL EtOH were mixed, and the mixture was named as solution B. The solution B was dropped into the solution A by 40– 50 drops per minute under a strong stirring condition. After continuously stirring for 20 min, 0.3 mL FA was then added to the mixed solutions under ambient conditions (25 C). The container was sealed and kept at 60 C for gelation. After gelation, the obtained wet gel was aged at 60 C for 48 h. The aged gels were solvent exchanged in EtOH at 60 C for 24 h three times, in TEOS/ EtOH (TEOS: EtOH = 1:4, volume ratio) solution at 60 C for 18 h twice, in n-hexane at 60 C for 48 h. Some of the dried gels were heat-treated at various temperatures between 500 and 900 C for 2 h in air. In this study, the molar ratios of TBOT: EtOH: HAc: water:
49
FA: F127 are designed to be 1: 24: 1.8: 6: 0.6: x (x = 0.0005, 0.001, 0.0015, 0.002, 0.0025 and 0.003), which were named as F05, F10, F15, F20, F25 and F30. The sample without adding F127 was named as F00. 2.2. Characterization of xerogels Apparent density was determined by measuring the volume (V) and mass (m) of xerogel samples. The xerogel samples with different amount of F127 have been cut into regular geometry to measure the volume and mass. Porosity was calculated by using apparent density (ra) and true density (rt, 3.9 g cm 3) as follows: porosity = (rt ra)/rt 100%. Differential thermal analysis (DTA) was performed using a WRT-3P analyzer. The crystalline phases of xerogels were analyzed by the X-ray diffraction (XRD) on a Rigaku. D/Max-RA X-ray diffractometer using nickel filtered Cu-Ka radiation in the range of 2u is 10–80 with a scanning speed of 2 /min. The morphology and microstructure were examined by scanning electron microscope (SEM, FEI SIRION, 25 KV) and transmission electron microscopy (TEM; FEI, Tecnai G2 F20). The nitrogen adsorption–desorption of xerogels were characterized by a nitrogen adsorption–desorption apparatus (Quantachrome Autosorb-1-C). Before each nitrogen adsorption–desorption measurement, the samples were degassed at 200 C under vacuum for more than 6 h. Specific surface areas were calculated using the Brunauer–Emmett–Teller (BET), meanwhile, the pore size distributions were calculated using the Barrett–Joyner–Halenda (BJH) method based on the desorption branch. 2.3. Electrochemical measurements Electrochemical experiments were carried out using standard CR 2025 type coin cells. The working electrode was prepared by mixing 70 wt % the active material (TiO2 xerogel powder), 20 wt % carbon black (Super P) and 10 wt % polyvinylidene fluoride (PVDF) binder. The slurry was pasted on a Cu foil and dried in a vacuum oven at 120 C for 12 h to serve as the working electrode. The weight of the active material in the electrode sheet was about 3 mg cm 2. Cells were assembled in in a glove box under highpurity argon atmosphere with a pure lithium foil as both counter electrode and reference electrode. The electrolyte solution was 1 M LiPF6 dissolved in a mixture of ethylene carbonate (EC)/dimethyl carbonate (DMC) at a volume ratio of 1:1. Cellgard 2300 polypropylene microporous film was used as the separator. Galvanostatical charge–discharge experiments were carried out in a voltage range of 1.0–3.0 V on a battery test system (Shenzhen Neware Battery,
Fig. 1. Schematic illustration of the fabrication process of hierarchically porous TiO2 xerogel monoliths.
50
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
China). Cyclic voltammetry (CV) measurements were conducted using a CHI650B electrochemical workstation (Shanghai Chenhua, China) at a scan rate of 0.1 mV s 1 in the potential range of 1.0–3.0 V and the impedance spectra were tested on the same workstation with the frequency ranging from 0.1 Hz to 100 KHz. 3. Results and discussion 3.1. Synthesis and characterization of the TiO2 xerogel monolith The synthesis and characterization of the TiO2 xerogel monolith were investigated in detail. During the reaction process, an alcogelgel with a spatial network structure is formed due to the hydrolysis and polycondensation of tetrabutyl titanate, consisting of titanium-oxygen based (RTi O TiR) polymers. The gelation time of TiO2 sols is related to the content of F127. As shown in Fig. 2a, with the increase of F127 content, the gelation time of TiO2 sol increases firstly (see F05 sample), and then rapidly decreases with the further increase of F127. F00 sample shows a gelation time of 18 min, while the gelation time of F30 sample is only about 6 min. It indicates that the gelation of TiO2 sol could be accelerated by the addition of F127. And white precipitations are observed in the experiment with the molar ratio of 0.0035 (F127:TBOT). Triblock copolymer F127 possesses the characteristic of relatively hydrophilic polyethylene oxide (PEO) block and relatively hydrophobic polypropylene oxide (PPO) block [29,30]. Additionally, F127 has a long molecular chain with a big molecular weight (12,600 g mol 1). It indicates that F127 has more significant impact on the gelation performance of TiO2 sol and the construction of gel network. As is well known, F127 presents two modification functions: wrapping and crosslinking [30]. The modification of wrapping restricts the single particle to contact or aggregate, while the modification of crosslinking accelerates the formation of much huger clusters. That is to say, when the added amount of the F127 is small, it is difficult for TiO2 sol particles to polycondensation due to the steric hindrance derived from wrapping, resulting in a long gelation time (as see F05 sample in Fig. 2a). The crosslinking function of F127 becomes dominant gradually with the increase of the added amount, and the gelation time of TiO2 sol reduces due to the strong polycondensation and the increase in amount and diameter of clusters. Fig. 2b displays the apparent densities and porosities of TiO2 xerogel with different F127 contents. It can be seen that the apparent density of TiO2 xerogel obviously decreases with the increase of F127 content, resulting in rapid increase of the
corresponding porosity. The minimum apparent density of 0.261 g 3 and the corresponding high porosity of 93.3% could be obtained at the molar ratio of 0.003 (F127:TBOT, sample F30). The result confirms that the addition of F127 could effectively promote the generation of porous structure, which leads to the decrease of the apparent density and the increase of the corresponding porosity for prepared TiO2 xerogel. Scanning electron microscopy (SEM) was undertaken to study the morphology of TiO2 xerogels. Fig. 3 shows the SEM images of fracture surface of TiO2 xerogels with different F127 contents. It is observed that only particulate xerogel with a lot of nanopores is obtained without F127 (F00). Compared with the F00 sample, the F05 xerogel sample displays a uniform morphology with smaller pore size and uniform pore size distribution, which is attributed to the steric hindrance of F127 as above mentioned. With the increase of F127 content, the macropores gradually form, and the particles aggregate to generate skeletons. Finally, the macroporous morphology with macropores and loose skeletons can be obtained with the mole ratio of F127:TBOT at 0.0025 (F25), and the xerogel exists in the form of monolith, as shown in Fig. 3 h. With the further increase of the F127 content (F30), the macropores become very large with relatively dense skeletons, and the xerogel presents granular shape again (no shown). The formation of unique porous structure can be attributed to the crosslinking function of F127 as above mentioned. During the process, F127 offers more entanglement with the skeleton particles, resulting in the networks increasingly loose, and the pores size rapidly increases to form macropores. In addition, the hydrophilic EO units of F127 are connected with surface hydroxyls of the gel networks via hydrogen bonding during and after the gelation. F127 will be away from the solvent phase and exposed toward the gel network through hydrophobic interaction of PO units [29,30]. The polycondensation of TiO2 sol particles after modification will increase the incompatibility of the mixed system, which results in the phase separation and the formation of macroporous structure [30]. Nitrogen isothermal absorption–desorption measurements were performed to determine the BET surface area and the porosity of the typical TiO2 xerogels. As revealed in Fig. 4. The isotherms of F00 and F25 samples present type IV with H3 hysteresis loops at relative pressure of 0.7–1.0, signifying the existence of slit-shaped pores [26,31,32]. While the isotherm of F30 sample is changed into type II, indicating that the formation of macroporous structure. It can be detected from the BJH pore size distributions (Fig. 4b), the major pores are distributed in large
cm
Fig. 2. (a) The gelation time of TiO2 sols with different F127 contents. (b) Apparent densities and porosities of TiO2 xerogels with different F127 contents.
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
51
Fig. 3. Morphology of fracture surface of TiO2 xerogels with different F127:TBOT ratios (a-without F127(F00), b-F127:TBOT = 0.0005(F05), c-F127:TBOT = 0.001(F10), d-F127: TBOT = 0.0015(F15), e-F127:TBOT = 0.002(F20), f-F127:TBOT=0.0025(F25), g-F127:TBOT = 0.003(F30), h-xerogel monolith of F25 sample).
(a)
(b)
Fig. 4. N2 adsorption–desorption isotherms (a) and pore size distribution of the typical TiO2 xerogels (b).
meso- to macropore region. And the pore size distribution of the F25 sample is relatively sharper than those of F00 and F30 samples. The variation of the size and shape of pores may be associated with the phase separation induced by F127. The obtained specific parameters of the BET surface area and median pore size are shown in Table 1, the dried gel without surfactant possesses a surface area
of 415 m2 g 1 with most pores at around 8 nm and a whole pore volume of 0.86 cm3 g 1. The addition of surfactant (F25 sample) increases surface area to 444 m2 g 1, and deceases the pore size to 7.0 nm and the pore volume to 0.82 cm3 g 1, respectively. With the further increases in surfactant (F30 sample), the surface area decreases to 397 m2 g 1, and the pore size and volume rapidly
Table 1 Pore characteristics of typical titania xerogels with different F127 contents. Sample
Specific surface area/m2 g
F00 F25 F30
415 444 397
1
Total pore volume/cm3 g 0.86 0.82 0.25
1
Average pore size/nm 8 7 3
52
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
decrease to 3 nm and 0.25 cm3 g 1 respectively, which is attributed to the formation of relatively dense skeletons. Fig. 5a shows the DTA curve of monolithic TiO2 xerogel (F25 sample). Two endothermic peaks near 90 and 225 C are mainly due to the volatilization of adsorbed water and the fracturing of -OH group, and a small exothermic peak at 305 C is attributed to the pyrolysis of organic species [33]. The very small exothermic peak locates at 665 C is possibly ascribed to the crystallization of anatase phase. Stengl et al. [34] reported that an exotherm appears at 420 C for anatase crystallization in titania gel gained by low temperature supercritical drying. In contrast, the high crystallization temperature in Fig. 5b indicates that the crystallization of anatase from titania gel is restricted to some extent, which could be attributed to the formation of macropore structure derived from phase separation [22]. Moreover, the small exothermic peak (305 C) corresponding to the pyrolysis of organic species indicates that the formation of interconnected macropores is beneficial to the removing of organic species during the solvent exchanging. The heat-treatment of the typical monolithic TiO2 xerogel monolith (F25 sample) was carried out at 500–900 C with a heating rate of 1–2 C/min. Fig. 5b shows XRD patterns of typical TiO2 xerogels at different heat-treatment temperatures. No discernible peaks can be detected after the heat-treated at 500 C. Very small peaks appear at 700 C, which is attributed to anatase phase, implying that amorphous structure begins to transform into anatase phase at this temperature. It is also seen that the diffraction peaks become strong with the increase of temperature. The diffraction peaks ascribed to anatase phase appear completely at 800 C, and these diffraction peaks show a characteristic of very broad width due to the formation of nanoscale anatase crystallites. When the heat-treatment temperature increases to 900 C, the anatase phase with broad diffraction peaks keep steadily, and there are no other phases can be detected, implying that the anatase phase has high thermal stability of phase transformation. In general, the transition to anatase of TiO2 xerogel happens at 420 C, and the transition of anatase-rutile occurs at 600–700 C [34,35]. It also confirms that the crystallization of anatase phase and the transition of anatase-rutile are restricted, which could be attributed to the unique porous structure. The result is consistent with the results of DTA (see Fig. 5a). Furthermore, the density of the F25 samples after heat-treatment is larger than that of dried xerogel (Fig. 2b) due to the volume shrinkage. After calcination at 500, 700, 800 and 900 C, the calculated densities are 0.47, 0.49, 0.61 and 0.60 g cm 3, and the
corresponding porosities are 87.9, 87.4, 84.3 and 84.6%, respectively. The increase in density with heat-treatment temperature can be ascribed to the decrease of porosity derived from the growth and agglomeration of grains. After heat treatment, the crystalline TiO2 xerogel can keep the monolithic shape basically and have a volume shrinkage of about 10 % compared with the dried xerogel. Fig. 6 shows the SEM image and N2 adsorption–desorption isotherm of TiO2 xerogel heattreated at 800 C. Although some agglomerations exist, the mophology is basically maintained after heat-treatment, implying that the formation of crystalline phase does not spoil the morphology of xerogel (Fig. 6a). In contrast to the macrostructure, the micro-mesoporous structure is significantly altered by the heat-treatment. Fig. 6b presents the isotherm of the xerogel heattreated at 800 C with a typical type II according to the IUPAC classification, representing a non-porous or macroporous structure. Compared with the dried gel (Fig. 4a), the specific surface area of the heat-treated gels (Fig. 6b) decreases rapidly (from 444 m2 g 1 to 61 m2 g 1). Moreover, the number of the pores, espically micro–mesopores decreases and the xerogel becomes much denser. After the heat-treatment at 800 C, the pores are basically distributed in large meso- to macropore region with micropores disappeared. The big variation of surface area and pore size after the heat-treatments is mainly attributed to the formation of anatase phase and the sintering effect of titania particles. The detail microstructure of TiO2 aerogel (F25 calcined at 800 C) was further investigated by TEM and HRTEM analysis. As shown in Fig. 7, the aerogels were constructed by nanoparticles with dozens of nanometers in size, and the nanoparticles interconnect into a network porous structure (Fig. 7a). It can be observed from the HRTEM image (Fig. 7b) that the lattice interplanar spacings of approximately 0.35 and 0.237 nm correspond to the (1 0 1) and (0 0 4) plane of tetragonal anatase TiO2, respectively, in good agreement with the results of XRD (Fig. 5). 3.2. Electrochemical performance It is reported that the morphology and microstructure of materials play an important role in the final electrochemical performance. In this work, the F25 sample shows a uniform morphology with macropores and loose skeletons, in comparison with other samples. Based on the previous works [9,19,20,36,37], it indicates that the F25 sample possesses a surperior electrochemical performance. Therefore, we investigate the electrochemical performance of the F25 xerogel sample after calcination by a series
Fig. 5. (a) DTA curve of TiO2 xerogel monolith (F25 sample). (b) XRD patterns of typical TiO2 xerogel (F25 sample) at different heat-treatment temperatures.
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
53
Fig. 6. Morphology (a) and N2 adsorption–desorption isotherm (b) of F25 xerogel monolith heat-treated at 800 C (inset is pore size distribution).
Fig. 7. Low (a) and high (b) magnification TEM image of the resulting TiO2 aerogel (F25 calcined at 800 C).
of electrochemical measurements, and the different calcination temperatures have been considered. As is well known, volumic capacity is an important parameter to evaluate the electrochemical performance of porous materials for the practical application [38,39]. In this work, the capacities of the samples are addressed not only in weight dimensionality (mAh g 1) but also by volume (mAh cm 3). The initial three charge–discharge profiles at 0.5 C in the voltage range of 1.0–3.0 V are presented in Fig. 8a. There are no obvious voltage plateau observed during the lithiation/delithiation processes, which can be ascribed to the solid solution behavior or the relative low crystallinity of the nanocomposite electrode, in accordance with previous reports on nanosized TiO2-based Li-ion batteries [40– 42]. In the first discharge step, a large specific capacity of 280 mAh g 1 (170 mAh cm 3) is obtained, indicating that 0.8 mol of Li ions are incorporated into the TiO2 xerogel electrode (Li0.8TiO2). The high discharge specific capacity in the first cycle can be possibly derived from an undesirable side reaction with the electrolyte or an irreversible change in the structure of the TiO2 xerogel upon deeper Li-ion insertion [40,43]. It is generally believed that the formation of SEI film happens when discharged to 0 V. In consideration of the high voltage plateau of TiO2 (1.5 V), the galvanostatical charge–discharge experiments were carried out in the voltage range of 1.0-3.0 V. Thus, the formation of the SEI layer during the electrochemical process could be effectively avoided. The large irreversible capacity has been observed in the first cycle could be attributed to the irreversible electrochemical decomposition of electrolyte or impurity phase over the TiO2 xerogel surface due to the porous structure [44,45]. Although a large irreversible capacity of about 130 mAh g 1 appears in the first cycle, the specific capacities of the TiO2
xerogel electrode are stabilized on following cycles, and the discharge and charge capacities in the third cycle are 140 and 135 mAh g 1, corresponding to 85 and 82 mAh cm 3, respectively. Moreover, sloped behaviors after the first discharge reaction are observed, which can be ascribed to the irreversible formation of a “nanocomposite” of crystalline grains and amorphous regions [46,47]. Fig. 8b shows the CVs of the TiO2 xerogel electrode at a scan rate of 0.1 mV s 1. The result is consistent with charge–discharge curves in Fig. 8a, there are no obvious intense peak detected during the scan process. Two broad current peaks are observed at around 1.7 V (cathodic sweep) and 2.0 V (anodic sweep), corresponding to the biphasic transition between tetragonal anatase and orthorhombic lithium titanate. The charge–discharge capacities versus cycle number for the F25 samples calcined at different temperture are shown in Fig. 8c. It can be seen that the reversible capacity of the F25 sample increases with the heat-treatment temperature in the range of 500–800 C, which is attributed to the increase of crystallinity. When the temperture increases to 900 C, the sample has a bad electrochemical performance, which might results from the collapse of porous structure. In contrast, the TiO2 xerogel electrode calcined at 800 C presents the best cycle performance, which delivers an initial discharge capacity of 180 mAh g 1 (109 mAh cm 3) and charge capacity of 135 mAh g 1 (82 mAh cm 3) with a coulombic efficiency of 75%. The coulombic efficiency increases upon cycling, which reaches to around 100% after the 10th cycle and keep unchanged during the whole cycles. Meanwhile, accompanied with the cycle number increasing, the reversible capacity decreases in the initial 10 cycles and then gradually levels off. It can still remain above 114 mAh g 1 (69 mAh cm 3) after 100 cycles, showing a low average capacity
54
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
Fig. 8. (a) Initial three charge–discharge curves (at 0.5 C, 1 C = 168 mAh g 1) and (b) CVs at a scan rate of 0.1 mV s 1of TiO2 xerogel (F25 calcined at 800 C). (c) Cycling performance (at 1 C) of TiO2 xerogels (F25) calcined at different temperatures. (d) Nyquist plots and corresponding simulation results of TiO2 xerogel (F25 calcined at 800 C), the inset is the equivalent circuit for plot fitting.
fading of 0.19% per cycle from the second cycle to 100th cycle. Furthermore, the reversible capacities of the samples calcined at 500, 700 and 900 C are 69, 76 and 67 mAh g 1, respectively, and the corresponding volumic capacities are 32, 37 and 40 mAh cm 3, respectively. Although the smallest porosity it has, the TiO2 xerogel calcined at 800 C presents the highest capacity both in mAh g 1 and mAh cm 3. The high charge–discharge capacities and superior cycle stability are due to the synergistic effect of crystallinity and hierarchically porous structure, guaranteeing the electrochemical
reaction and ensuring better contact between the electrode material and the electrolyte. Electrochemical impedance spectroscopy (EIS) was performed for the TiO2 xerogel electrode before and after the galvanostatic cycling and the corresponding Nyquist plots are shown in Fig. 8d. It can be seen that all the Nyquist plots exhibit the characteristic of a semicircle in the high frequency region and a sloping straight line in the low frequency region. The semicircle in high frequency is ascribed to the charge-transfer impedance at the electrolyte/
Fig. 9. Schematic diagram of the insertion/extraction process of the hierarchically porous TiO2 xerogel monoliths.
W. Zhu et al. / Materials Research Bulletin 73 (2016) 48–55
electrode interface, and the sloping straight line of low frequency region corresponds to the lithium ion diffusion process in the electrodes [31,48]. The inset shows the equivalent circuit for the TiO2 xerogel to simulate the experimental EIS plots, and the corresponding fitting curves are drawn in Fig. 8d, respectively. Rs is attributed to the solution resistance (including total ohmic resistance of the separator, electrolyte, and contacts between the current collector and the electrode materials), CPE (i) and Rsl (i) (i = 1, 2) represent the capacity of the layer and the migration of hydrogen ions, respectively. Rct and Cdl denote the charge-transfer resistance and a double-layer capacitance, respectively. ZW is the Warburg impedance, representing the diffusion of electrolyte ions into the electrode materials, which is associated with the inclined line at low frequencies. As indicated in Fig. 8d, a small semicircle was obtained before the cycling, indicating low resistances of Li+ diffusion in charge-transfer reaction. After cycling, the semicircle is enlarged, which suggests the resistance of Li+ diffusion increases upon cycles (Fig. 9). This is probably due to the large volume change during charge–discharge cycling that gradually deteriorates the contact between TiO2 xerogel and the collector, which resulting in the 20 % fade of capacity during the specific capacities versus cycle test in Fig. 8c. 4. Conclusions In summary, monolithic titania (TiO2) xerogels were synthesized by combining a sol–gel process with phase separation, followed by ambient pressure drying to remove the solvents. The as-prepared titania xerogel monolith exhibits an unique hierarchically porous structure constructed by large macropores and mesoporous skeletons. The monolithic shape and morphology of xerogels are not spoiled after heat-treatment, the BET surface area rapidly decreases from 444 to 61 m2 g 1 because of the formation and sintering of anatase nanoparticles. The crystallization of anatase phase and the transition of anatase-rutile are restricted due to the unique porous structure. The TiO2 xerogel calcined at 800 C presents the best electrochemical performance, the reversible capacity could remain over 114 mAh g 1 (69 mAh cm 3) at 1 C after 100 cycles with a good cycling stability. The route provides a facile strategy for the preparation of hierarchically macroporous xerogel monoliths, and we believe the hierarchically porous titania (TiO2) xerogel monolith will also extend its application fields such as catalysis, separation, absorption and energy conversion. Acknowledgements This work is supported by the National Natural Science Foundation of China (51372225), and Zhejiang Provincial Natural Science Foundation of China (LY13B010001). References [1] J.M. Tarascon, M. Armand, Nature 414 (2001) 359. [2] Y. Xia, Z. Xiao, X. Dou, H. Huang, X.H. Lu, R.J. Yan, Y.P. Gan, W.J. Zhu, J.P. Tu, W.K. Zhang, X.Y. Tao, ACS Nano 7 (2013) 7083. [3] L.X. Yuan, Z.H. Wang, W.X. Zhang, X.L. Hu, J.T. Chen, Y.H. Huang, J.B. Goodenough, Energy Environ. Sci. 4 (2011) 269. [4] H. Huang, W. Zhu, X. Tao, Y. Xia, Z. Yu, J. Fang, Y. Gan, W. Zhang, ACS Appl. Mater. Interface 4 (2012) 5974.
55
[5] W. Zhu, H. Huang, W. Zhang, X. Tao, Y. Gan, Y. Xia, H. Yang, X. Guo, Electrochim. Acta 152 (2014) 286. [6] N. Brun, S.R.S. Prabaharan, C. Surcin, M. Morcrette, H. Deleuze, M. Birot, O. Babot, M.F. Achard, R. Backov, J. Phys. Chem. C 116 (2012) 1408. [7] B.T. Zhao, S.M. Jiang, C. Su, R. Cai, R. Ran, M.O. Tade, Z.P. Shao, J. Mater. Chem. A 1 (2013) 12310. [8] X. Xin, X.F. Zhou, J.H. Wu, X.Y. Yao, Z.P. Liu, ACS Nano 6 (2012) 11035. [9] Y.G. Guo, Y.S. Hu, W. Sigle, J. Maier, Adv. Mater. 19 (2007) 2087. [10] D. Dambournet, I. Belharouak, K. Amine, Chem. Mater. 22 (2010) 1173. [11] N. Li, G. Liu, C. Zhen, F. Li, L.L. Zhang, H.M. Cheng, Adv. Funct. Mater. 21 (2011) 1717. [12] P. Kubiaka, J. Geserick, N. Husing, A. Wohfahrt-Mehrens, J. Power Sources 175 (2008) 510. [13] F. Carn, A. Colin, M.F. Achard, H. Deleuze, C. Sanchez, R. Backov, Adv. Mater. 17 (2005) 62. [14] F. Carn, M.F. Achard, O. Babot, H. Deleuze, S. Reculusa, R. Backov, J. Mater. Chem. 15 (2005) 3887. [15] J.W. Zhang, X.X. Yan, J.W. Zhang, W. Cai, Z.S. Wu, Z.J. Zhang, J. Power Sources 198 (2012) 223. [16] S.T. Myung, N. Takahashi, S. Komaba, C.S. Yoon, Y.K. Sun, K. Amine, H. Yashiro, Adv. Funct. Mater. 21 (2011) 3231. [17] Y.-H. Ding, P. Zhang, H.-M. Ren, Q. Zhuo, Z.-M. Yang, Y. Jiang, Mater Res Bull 46 (2011) 2403. [18] Z.X. Yang, G.D. Du, Q. Meng, Z.P. Guo, X.B. Yu, Z.X. Chen, T.L. Guo, R. Zeng, J. Mater. Chem. 22 (2012) 5848. [19] I. Moriguchi, R. Hidaka, H. Yamada, T. Kudo, H. Murakami, N. Nakashima, Adv Mater 18 (2006) 69. [20] V. Gentili, S. Brutti, L.J. Hardwick, A.R. Armstrong, S. Panero, P.G. Bruce, Chem Mater 24 (2012) 4468. [21] H.E. Wang, H. Cheng, C.P. Liu, X. Chen, Q.L. Jiang, Z.G. Lu, Y.Y. Li, C.Y. Chung, W.Y. Zhang, J.A. Zapien, L. Martinu, I. Bello, J. Power Sources 196 (2011) 6394. [22] G. Hasegawa, T. Sato, K. Kanamori, K. Nakanishi, T. Abe, New J Chem 38 (2014) 1380. [23] N. Brun, S.R.S. Prabaharan, M. Morcrette, C. Sanchez, G. Pecastaings, A. Derre, A. Soum, H. Deleuze, M. Birot, R. Backov, Adv. Funct. Mater. 19 (2009) 3136. [24] M. Zeller, V. Lorrmann, G. Reichenauer, M. Wiener, J. Pflaum, Adv Energy Mater 2 (2012) 598. [25] R.P. Maloney, H.J. Kim, J.S. Sakamoto, ACS Appl. Mater. Interface 4 (2012) 2318. [26] H. Yang, W.J. Zhu, S. Sun, X.Z. Guo, RSC Adv. 4 (2014) 32934. [27] X.Y. Chen, B.H. Liu, Z.P. Li, Solid State Ionics 261 (2014) 45. [28] F.B. Hao, Z.W. Zhang, L.W. Yin, ACS Appl. Mater. Interface 5 (2013) 8337. [29] R. Barreiro-Iglesias, C. Alvarez-Lorenzo, A. Concheiro, J. Control. Release 93 (2003) 319. [30] M. Kurahashi, K. Kanamori, K. Takeda, H. Kaji, K. Nakanishi, RSC Adv. 2 (2012) 7166. [31] W.J. Zhu, H. Huang, Y.P. Gan, X.Y. Tao, Y. Xia, W.K. Zhang, Electrochim. Acta 138 (2014) 376. [32] W.Y. Li, X.Z. Guo, Y. Zhu, Y. Hui, K. Kanamori, K. Nakanishi, J Sol–Gel Sci. Technol. 67 (2013) 639. [33] Z.Q. Zhao, X.L. Jiao, D.R. Chen, J. Mater. Chem. 19 (2009) 3078. [34] V. Stengl, S. Bakardjieva, J. Subrt, L. Szatmary, Microporous Mesoporous Mater. 91 (2006) 1. [35] M. Popa, D. Macovei, E. Indrea, I. Mercioniu, I.C. Popescu, V. Danciu, Microporous Mesoporous Mater. 132 (2010) 80. [36] W. Zhu, H. Yang, K. Nakanishi, K. Kanamori, X. Guo, RSC Adv. 5 (2015) 24803. [37] X.B. Wang, Y.Y. Wang, L. Yang, K. Wang, X.D. Lou, B.B. Cai, J. Power Sources 262 (2014) 72. [38] M. Depardieu, R. Janot, C. Sanchez, A. Bentaleb, R. Demir-Cakan, C. Gervais, M. Birot, M. Morcrette, R. Backov, J. Mater. Chem. A 2 (2014) 18047. [39] M. Depardieu, R. Janot, C. Sanchez, A. Bentaleb, C. Gervais, M. Birot, R. DemirCakan, R. Backov, M. Morcrette, RSC Adv. 4 (2014) 23971. [40] D.H. Kim, K.M. Min, K.S. Park, I.J. Park, I.S. Cho, W.M. Seong, D.W. Kim, K.S. Hong, Dalton Trans. 42 (2013) 4278. [41] K.S. Park, K.M. Min, Y.H. Jin, S.D. Seo, G.H. Lee, H.W. Shim, D.W. Kim, J. Mater. Chem. 22 (2012) 15981. [42] Z.S. Hong, M.D. Wei, T.B. Lan, L.L. Jiang, G.Z. Cao, Energy Environ. Sci. 5 (2012) 5408. [43] K.S. Park, J.G. Kang, Y.J. Choi, S. Lee, D.W. Kim, J.G. Park, Energy Environ. Sci. 4 (2011) 1796. [44] M. Rahman, J.Z. Wang, M.F. Hassan, D. Wexler, H.K. Liu, Adv. Energy Mater. 1 (2011) 212. [45] F. Wu, X. Li, Z. Wang, H. Guo, Nanoscale 5 (2013) 6936. [46] J. Jamnik, J. Maier, Phys. Chem. Chem. Phys. 5 (2003) 5215. [47] P. Balaya, H. Li, L. Kienle, J. Maier, Adv. Funct. Mater. 13 (2003) 621. [48] G.F. Cai, J.P. Tu, D. Zhou, J.H. Zhang, Q.Q. Xiong, X.Y. Zhao, X.L. Wang, C.D. Guts, J. Phys. Chem. C 117 (2013) 15967.