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Porous carbon-wrapped cerium oxide hollow spheres synthesized via microwave hydrothermal for long-cycle and high-rate lithium-ion batteries
Electrochimica Acta 256 (2017) 110–118
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Research paper
Porous carbon-wrapped cerium oxide hollow spheres synthesized via microwave hydrothermal for long-cycle and high-rate lithium-ion batteries Mingbo Maa,b , Hongjie Wanga,* , Sen Liangb , Shengwu Guoa , Yuan Zhangb , Xianfeng Dub,* a
State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, PR China Electronic Materials Research Laboratory, Key Laboratory of the Ministry of Education & International Center for Dielectric Research, Xi’an Jiaotong University, Xi’an 710049, PR China b
A R T I C L E I N F O
Article history: Received 5 July 2017 Received in revised form 30 September 2017 Accepted 5 October 2017 Available online 7 October 2017 Keywords: CeO2/C hollow spheres high-rate performance anodes lithium-ion batteries
A B S T R A C T
Porous hollow structures have attracted tremendous interests due to their geometrical beauty, unique structural features and fascinating physicochemical properties. In the present work, we developed a general procedure for the synthesis of hollow nanostructured carbon-wrapped cerium oxide (CeO2/C) via a microwave hydrothermal process without any surfactants or hard templates. The electrochemical performances of CeO2/C specimen were tested as anode materials in lithium ion batteries. The results showed that CeO2/C hollow spheres exhibited an exceptional cyclic stability with a high reversible capacity of 313 mA h g1 after 500 cycles at a high current density of 1000 mA g1 without signs of further degradation and also presented an excellent rate performance from 1000 to 10000 mA g1. The improved performances were attributed to the homogeneous carbon with 3D network structure in hollow spheres and the Ce3+ in the oxygen-deficient fluorite-like CeO2-x on the surface of ceria nanoparticles, which enhanced the conductivity of CeO2 hollow spheres, suppressed the aggregation of active particles and reduced the apparent activation energy to facilitate the kinetics of Li+ insertion-extraction. The unique hollow structure of CeO2 wrapped by conductive carbon have made CeO2 a promising candidate for future applications in various metal oxide electrodes to mitigate the mechanical degradation and capacity fading, critical for developing advanced electrochemical energy storage systems with long-cycle life and high-rate performance. © 2017 Elsevier Ltd. All rights reserved.
1. Introduction Lithium ion batteries (LIBs) have been widely used in portable electronics, electric tools and electric vehicles [1–3] due to their superior properties such as high energy density, long cycle life and good environmental compatibility [4,5]. In order to satisfy the ever-growing demands on high-capacity and high-power LIBs, great efforts have been made on transition metal oxides [6,7], which possess much higher capacity compared with widely-used graphite in commercial LIBs. However, the main challenges in the implementation of metal oxide electrodes are their large volume variation and poor electronic conductivity during the charge/ discharge process [8]. As a kind of important transition metal oxides, CeO2 shows fluorite structure [9,10], and the oxidation state
* Corresponding authors. E-mail addresses: [email protected] (H. Wang), [email protected] (X. Du). https://doi.org/10.1016/j.electacta.2017.10.041 0013-4686/© 2017 Elsevier Ltd. All rights reserved.
of cerium can mutate quickly between Ce(III) and Ce(IV) [8,11], leading to its wide applications in solid-state fuel cells [12], oxygen storage materials [13] and catalytic supporters [14]. The CeO2 nanoparticles displayed imperceptible volumetric and morphological changes during the lithiation/delithiation process [8], which made it a possible candidate as an anode material for LIBs in the future in case that its electronic conductivity can be improved greatly. It is well known that carbon materials have high electronic conductivity that can effectively improve the conductance of active materials and always act as buffering matrix to suppress the aggregation of active particles in the electrode [11]. So numerous efforts have been made to synthesize carbon material/metal oxides complexes for LIBs, such as SnO2@carbon composite [15], grapheme-encapsulated Fe3O4 [16] and Co3O4-graphene nanoflowers [17]. During the passed couple of years, a variety of CeO2 and carbon composites with different morphologies and microstructures, including nanocomposite [18], nanospheres [11,19] and
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nanorods [20] have been synthesized to enhance its electronic conductivity as anode materials in LIBs. Specially, CeO2-graphene nanocomposite prepared by hydrothermal method exhibited improved specific capacity of 605 mA h g1 at 50 mA g1 after 100 cycles [18]. Core-shell CeO2@C nanospheres synthesized by hydrothermal carbonization showed an initial discharge specific capacity of 863 mA h g1 in the potential range of 3.0-0.0V [11]. Graphene/CeO2/CMK-3 composites with a hierarchical structure prepared by a template method maintained the specific capacity of about 550 mA h g1 at 100 mA g1 after 100 cycles [20]. All these CeO2 materials achieved good electrochemical performances at the expense of the complex preparation process and long synthesizing time [21–23]. Moreover, electrochemical performances of CeO2 have not been investigated at large current densities for lithium ion battery in previous reports. Herein, we report a simple microwave hydrothermal method by using gas templates to synthesize hollow spheres composed by carbon-wrapped CeO2 nanoparticles (noted as CeO2/C). The structure, shape and formation process of CeO2/C were investigated by SEM, TEM, XRD, Raman and BET. And we further demonstrated their promising properties as anode materials for lithium-ion batteries. The results showed that CeO2/C hollow spheres exhibited improved discharge capacity and remarkable long-cycle stability (313 mA h g1 at 1000 mA g1 with 100% retention after 500 cycles), and excellent high-rate capacity (even at the large current density of 10000 mA h g1).
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2. Experimental 2.1. Materials All chemicals were of analytical grade and used as received without further purification. Cerium nitrate hexahydrate (Ce (NO3)36H2O), citric acid monohydrate (C6H8O7H2O), polyvinylidene fluoride (PVDF) and N-methyl-2-pyrrolidone(NMP) were purchased from Sinopharm Chemical Reagent Co., Ltd., China. And deionized (DI) water was produced by Milli-Q-Reference water system (Millipore Co., USA). 2.2. Preparation of CeO2/C hollow spheres CeO2 precursor was synthesized through microwave hydrothermal process. Typically, 0.1 mol Ce(NO3)36H2O and 0.05 mol C6H8O7H2O were dissolved in 1 L DI water at room temperature. After being stirred for 4 h, the obtained solution was transferred into quartz tube and then maintained at 150 C for 2 min under microwave condition with the wattage between 95 and 130 W (Initiator+TM Microwave System, Biotage Co., Sweden). Fig. S1 recorded the profiles of temperature (T), pressure (p) and power (P) during microwave hydrothermal process. After the system cooled totally, the solution became turbid. Then the obtained solution was filtrated, washed and dried to get the light-yellow precipitate of hollow sphere marked as CeOC. Finally, the CeOC
Fig. 1. SEM images of (a1-a3) CeOC, (b1-b3) CeO2 and (c1-c3) CeO2/C; EDS element mapping images of (a4-a6) CeOC, (b4-b6) CeO2 and (c4-c6) CeO2/C.
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samples annealed at 300 C for 1 hour in air became yellow and were noted as CeO2 while the ones pyrolyzed at 600 C for 1 hour in Ar flow became black and were denoted as CeO2/C. 2.3. Structure characterization Morphology and composition analysis of all these samples were observed using field-emission scanning electron microscopy with energy-dispersive spectrometer (FE-SEM, FEI6000, USA) and highresolution transmission electron microscopy (HR-TEM, JEM-2100F, Japan). Their crystalline structures were characterized by X-ray diffraction (XRD, PANalytical, X’pert Pro, Holland), Raman spectroscopy (Horiba HR800, France) and X-ray photoelectron spectroscopy (XPS, Thermo Fisher Scientific, ESCALAB Xi+, America). The weight changes during heat-treatment process were investigated through thermogravimetry analysis (TGA, METTLER TOLEDO, Switzerland). The electronic conductivity was estimated in a potentiostatic polarization method with the polarizing voltage of 0.5 V. The specific surface area was measured by BrunauerEmmett-Teller (BET, ASAP 2020, Micromeritics, USA). The pore size distributions were determined by Barrett-Joyner-Halenda (BJH, ASAP 2020, Micromeritics, USA) at liquid nitrogen temperature. The mass of samples was weighted by high-precision analytical balance (XS105DU, METTLER TOLEDO, Switzerland).
spectroscopy (EDS) shown in Fig. 1. It is clear to see that all these samples show hollow sphere structures with diameter distribution within 0.5–2 mm and wall thickness of about 100 nm. The free volume in hollow structure can effectively accommodate the volume changes and enhance the cyclic stability of active materials [24]. The schematic diagram of possible formation process for hollow sphere structures has been proposed in Fig. 2. Above all, Ce (III) combined with citric acid to form the complex cerium citrate (Ce(III)-citrate), of which the part was decomposed to be carbon dioxide and water at the elevated temperatures and pressure (Fig. S1). Meanwhile, Ce(III) was oxidized by oxygen and/or nitrates involved in the reaction system to synthesize poorly-crystallized CeO2[21]. Subsequently, the rest organic cerium compound and the obtained CeO2 were attached to the surface of gas templates. Two minutes later, CeOC hollow spheres can be obtained only in the Ce (NO3)3 and C6H8O7 mixed solution when the pressure reached to 20 bar (Fig. S2). The main reactions involved in the reaction system could be illustrated as follows[21]: (1)
(2)
2.4. Electrochemical measurements Electrochemical experiments were performed by using 2016 type coin cells. The test electrodes were prepared by adding active material, acetylene black and polyvinylidene fluoride (PVDF) binder into N-methy-l-2-pyrrolidone (NMP) solution with the weight ratio of 60: 30: 10, followed by grinding the mixture into homogeneous slurries. Then, the slurries were pasted on the Cu foil and dried at 120 C in a vacuum oven. The testing coin-type cells were assembled in an argon filled glove box with as-prepared electrodes, Li foil as counter and reference electrodes, a Celgard film (America) separator and 1 M LiPF6 in 1:1:1 (by volume) ethylene carbonate (EC)-propylene carbonate (PC)-dimethyl carbonate (DMC) electrolyte (Capchem Technology Co., Ltd. Shenzhen, China). All the tests were conducted on the battery tester (CT2001A, LAND, China; BT2000, Arbin, USA) within 0.01-3.00 V vs. Li/Li+ at various current densities at room temperature. The weight of electrode material was 0.7–1 mg cm2. The cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) measurements were carried out on an electrochemical workstation (VMP2, Princeton, USA) in a three-electrode configuration with CeO2/Cu as working electrode and Li foil as counter and reference electrode. 3. Results and discussion The morphologies and element mapping images of CeOC, CeO2 and CeO2/C were observed by FE-SEM with energy disperse
4HNO3 ! 4NO2 + O2 + 2H2O
(3)
Citric acid not only served as a carbon precursor for CeO2/C hollow spheres but also promoted the formation of hollow spherical shape in the process. Fig. 1a1 and a3 show the smooth morphology of CeOC with homogenously distributed C, O and Ce elements shown in Fig. 1a4 to a6. After being annealed in air, carbon in CeOC was totally burned up and the observed homogeneously-distributed carbon with small amount in Fig. 1b1 to b6 apparently came from the ambient environment. The surface of CeO2 hollow spheres became much rougher because of the removed carbon from CeOC. While C element partly remained in CeO2/C as a result of the pyrolyzed CeOC in Ar, it was quite useful to improve the electronic conductivity of CeO2. What’s more, the surface of CeO2/C hollow spheres become less rougher compared with CeO2 in Fig. 1c1 to c6. Fig. 3 shows microstructures and phase compositions of CeOC, CeO2 and CeO2/C observed by TEM. A strong contrast between paler center and darker edge in all these samples indicates their hollow interior with some differences in the surface shown in Fig. 3a, d and 3g. Specially, CeOC can not be easily penetrated by electron beam for its smooth surface and thick wall. CeO2 shows the porous morphology with thin wall and some black particles distributed on its surface. CeO2/C wrapped by carbon layer presents the middle morphology features between CeOC and
Fig. 2. Schematic diagram of the formation of CeO2 and CeO2/C hollow spheres fabricated through a microwave hydrothermal process.
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Fig. 3. HR-TEM images and SEAD patterns of (a-c) CeOC, (d-f) CeO2 and (g-i) CeO2/C.
CeO2. Then the selected area electron diffraction (SEAD) patterns were conducted to describe the crystallinity and phases of these samples. Several rings in Fig. 3b, e and h reveal their polycrystalline nature. The weak rings in Fig. 3b indicate the low crystallinity of CeOC. The rings made up of points in Fig. 3e demenstrate the wellcrystalized nanocrystalline structure of CeO2. In Fig. 3h, the rings constituted with points suggests the nanocrystalline nature of CeO2/C. The smallest ring is very broad and bright due to the overlap of (111) plane of CeO2 and (002) plane of C, further proved by the high-resolution TEM (HR-TEM) image below. HR-TEM image in Fig. 3c discloses the existence of nanocrystals and amorphous state in CeOC. The lattice fringe of these nanocrystals is 0.308 nm, corresponding to the (111) plane of CeO2. The amorphous substance may originate from the decomposition product of citric acide. The lattice fringes of 0.308 nm in Fig. 3f illustrate the existence of CeO2 nanocrystals in CeO2 samples. The nanoparticles with lattice fringes of 0.308 nm and crooked ones of 0.33 nm are corresponding to (111) plane of CeO2 and (002) plane of C respectively in Fig. 3i. More importantly, carbon wrapping up on surfaces and interfaces of CeO2 nanocrystals can favour the
improvement of the electronic conductivity of CeO2/C and promote rapid charge-transfer reactions in lithium ion batteries. Fig. 4a and b display the phase compositions of CeOC, CeO2 and CeO2/C carried out by powder XRD and Raman spectrum. Obviously, no peaks of CeOC appear in Fig. 4a, indicating its low crystallinity and small crystalline grains. On the contrary, both CeO2 and CeO2/C display the same characteristic peaks at 28.5 , 33.1, 47.5 , 56.3 , 59 , 69 , 76.7 and 79 , which match up well with the CeO2 structure. Moreover, Raman active mode at 465 cm1 in Fig. 4b demonstrates the existence of cubic CeO2 nanocrystals [25]. In addition, CeO2/C exhibits broad bands around the 1200– 1500 cm1 region, corresponding to carbon-based Rama vibration modes. The peaks at 1350 and 1560 cm1 can be attributed to the disordered D-band carbon and the graphitic G-band carbon [26], respectively. In brief, the result of Raman analysis further confirms the co-existence of carbon and CeO2 in our CeO2/C samples. In addition, XPS analysis was carried out to further confirm the phase compositions of CeO2 and CeO2/C and the results were shown in Fig. S3 and Table S1. It is obvious to see the existence of Ce3+ in the surface of these samples, which is a result of oxygen vacancies and enhanced in cerium oxide nanoparticles [27,28]. The presence of
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Fig. 4. (a) XRD patterns and (b) Raman spectra of CeO2 and CeO2/C; (c) TGA curves and (d) the evolution of current with polarization time under the polarization voltage of 0.5 V in CeO2 and CeO2/C; N2 adsorption-desorption isotherms and pore-size distribution of (e) CeO2 and (f)CeO2/C.
Ce3+ implies that the oxygen-deficient fluorite-like CeO2-x is located on the surface of ceria nanoparticles [29,30]. The transformation of Ce4+ to Ce3+ in CeO2-x driven by oxygen vacancies will result in the significant increase in the electronic conductivity of CeO2/C. In Fig. 4c, TGA analysis was carried out to determine the mass changes of CeOC in different gas and the weight percentage of carbon in CeO2/C. The mass of CeOC remains unchanged after 300 C in air and after 600 C in Ar. Obviously, 15 wt% carbon in CeO2/C may improve the electronic conductivity of CeO2/C, which has been estimated by potentialstatic polarization method. The corresponding current was measured as function of the polarization period and their relationship is shown in Fig. 4d. The electronic conductivity (se) of CeO2 and CeO2/C can be calculated from the following equation [31]:
se ¼
4Il
pD2 U
ð4Þ
Where U is the polarization voltage, l is the thickness of the sample, D is the diameter of Cu electrode, and I is the current, respectively. The result of calculation shows that the electronic conductivity is increased by almost 40 times from 1.98*109 S cm1 of CeO2 to 8.71*108 S cm1 of CeO2/C. The significantly increased electronic conductivity of CeO2/C can be attributed to the oxygendeficient fluorite-like CeO2-x and the carbon, which will favor its enhanced electrochemical performance as anode for LIBs. In order to evaluate the specific surface area and the pore structures of CeO2 and CeO2/C, the nitrogen adsorption-desorption measurements were implemented and the results are shown in Fig. 4e and f. The isotherm of CeO2 presents a distinct hysteresis loop, which belongs to type IV[32] plot on the basis of the IUPAC classification, indicating the presence of mesoporous structure. Its BET specific surface area is about 85.13 m2 g1. The inset in Fig. 4e shows the diameters of the mesopores are less than 10 nm which is calculated from the adsorption branch. The isotherm of CeO2/C exhibits a type II plot [33] with the decreased BET specific surface area of 9.56 m2 g1 due to the formation of the carbon phase that filled the partial pores and holes in CeO2/C. The pore diameters of most mesopores in Fig. 4f inset are less than 10 nm. Clearly,
variation of electronic conductivity and specific surface area play an important role in the electrochemical properties of anode materials. As for carbon-wrapped CeO2 specimen, the specific surface area decreased by 8 times while the electronic conductivity increased by 40 times. It should be mentioned that the decreased specific surface area of CeO2/C samples did not have any negative impacts on their electrochemical performances, which instead might be tremendously enhanced due to the significantly increased electronic conductivity that should make much more contributions than any other factors to the outstanding electrochemical properties of CeO2/C. Otherwise, the small specific surface area also promote the initial coulombic efficiency of CeO2/C as anode for LIBs by decreasing the amount of solid electrolyte interphase (SEI) films. All these conclusions will be further discussed and proved in the section of electrochemical analysis. The electrochemical impedance spectroscopy (EIS) was utilized to further study the charge transfer and ion diffusion kinetics of CeO2 and CeO2/C electrodes. Both Nyquist plots in Fig. 5 exhibit a semicircle in the high frequency region and a straight line in the low frequency region. The semicircle is ascribed to the SEI film resistance and charge transfer resistance [34] while the straight
Fig. 5. EIS Nyquist plots CeO2 and CeO2/C electrodes obtained at room temperature.
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Fig. 6. Electrochemical performances of CeO2 and CeO2/C. CV curves of (a) CeO2 and (b) CeO2/C at 0.1 mV s1; the initial two galvanostatic discharge-charge profiles for (c) CeO2 and (d) CeO2/C at 100 mA g1; (e) cycle performance at 1000, 4000 and 8000 mA g1; (f) rate capability tests from 1000 to 10000 mA g1.
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line (Warburg line) is attributed to Li-ion diffusion at the interface of electrolyte and electrode [34]. Compared with CeO2 electrode, the semicircle of CeO2/C electrode shows a reduced diameter, indicating the decreased charge transfer (Rct) resistance at the interface of electrode and elelctrolyte due to the homogeneous conductive carbon. The electrochemical performances of these electrodes are evaluated and presented in Fig. 6. Fig. 6a and b which show the cyclic voltammetry (CV) curves of CeO2 and CeO2/C at a scan rate of 0.1 mV s1, respectively. Three obvious cathodic peaks at about 1.16, 0.9 and 0.02 V for CeO2 during the first discharge process (Fig. 6a). In particular, the peak at 1.16 V may originate from the decomposition of the electrolyte on the partially lithiated electrode [32] and disappear in the first anodic scan. The peak at 0.9 V corresponds to the formation of SEI film and the reduction reaction of CeO2/CeO2-x to Ce as described in function (5) [19,35], while the reduction peak at 0.05 V represents the insertion of Li+ into carbon[36]. In the first anodic scan, the observed broad peak at 1.3 V on behalf of the oxidation reaction of Ce to CeO2/CeO2-x suggests the corresponding reversible reaction[35]. The oxidation peak at 0.1 V could be assigned to the extraction of Li+ from carbon. In the second cycle, the reduction peak of CeO2/CeO2-x shifts to 1.0 V and becomes broad while the oxidation peak of Ce shifts to 1.25 V, revealing that the insertion/extraction of Li+ in CeO2/CeO2-x is a reversible process. The CeO2/C electrode presents the similar electrochemical behaviors as given in Fig. 6b. CeO2 + 4Li+ + 4e = Ce + 2Li2O
(5)
CeO2-x + 2(2-x)Li+ + 2(2-x)e = Ce + (2-x)Li2O
(6)
Fig. 6c and d presents the initial discharg-charge voltage profiles of CeO2 and CeO2/C electrodes at the current density of 100 mA g1 within a cut-off voltage window of 0.01-3.0 V. The initial discharge/charge specific capacities are found to be 749/ 418 mA h g1 for CeO2/C, much higher than that of CeO2 (355/ 176 mA h g1), which benefits from the introduction of carbon and the uniform structure between C and CeO2 in CeO2/C. Futhermore, the irreversible capacity in the first cycle should be mainly attributed to the formation of SEI layer on the electroactive materials and the electrolyte decompositon. Because of the decreased specific surface area, the initial coulombic efficiency of the CeO2/C electrode is 60.5%, much higher than that of the CeO2 electrode (48%) in Fig. 6e. After several cycles, the coulombic efficiency increased to nearly 100% for both electrodes. From the second cycle onwards, both electrodes showed good reversibility. At the end of 500 discharge-charge cycles, CeO2/C electrodes still retained a reversible specific capacity of 313 mA h g1 (1000 mA g1), 210 mA h g1 (4000 mA
g1) and 169 mA h g1 (8000 mA g1) and exhibited outstanding capacity and cycle stability. Compared with the current works, CeO2/C hollow sphere shows excellent properties at large current densities from Table 1. However, CeO2 electrodes with low conductivity only retained 121 mA h g1 (1000 mA g1), 113 mA h g1 (4000 mA g1) and 47 mA h g1 (8000 mA g1) at the end of 500 discharge-charge cycles. CeO2/C electrodes also exhibit high specific capacity and good cycling stability at the current densities of 100, 200 and 500 mA g1 in Fig. S4. When changed the upper voltage limits of the voltage window, the specific capacity and the average coulombic efficiency CeO2/C at 1000 mA g1 have no obvious changes. While the lower voltage limit of the voltage window is changed, the initial discharge specific capacities and the initial coulombic efficiency CeO2/C at 1000 mA g1 are changed obviously (Fig. S5). After cycling, the CeO2/C electrode maintained its hollow sphere structure without any obvious changes in Fig. S6, which demonstrated that porous CeO2/C hollow sphere with a steady structure resulted in the excellent cycling stability as electrodes. The rate performances of CeO2 and CeO2/C were shown in Fig. 6f. Profiting from the unique structure and synergetic effect, CeO2/C presents better reversibility and higher specific capacity than CeO2 at different current rates, especially at high discharge/ charge rates. For example, CeO2/C electrode delivers specific capacity of 266, 232, 210, 188 and 169 mA h g1 at 1000, 2000, 4000, 8000 and 10000 mA g1. Whereas the values for CeO2 are 100, 84, 75, 68 and 58 mA h g1, respectively. What’s more, after being tested for 20 cycles under various current rates (even up to 10000 mA g1), the capacities of CeO2/C electrodes can recover to 300 mA h g1 at 1000 mA g1 and remain unchangeable after 500 cycles, indicating the high reversibility of lithium ion insertion/ extraction in CeO2/C. Moreover, the increased specific capacity of CeO2/C electrodes after cycling at 10000 mA h g1 due to the actived electrode enlarges the reduction peak in the CV curve and decreases the charge transfer (Rct) resistance of CeO2/C electrodes in Fig. S7. This phenomenon also appears in core–shell CeO2@C nanospheres [11]. To further illuminate the excellent performance of CeO2/C electrode, the apparent activation energies (Ea) of CeO2 and CeO2/C were calculated through the relationship between Arrhenius Equation (Equation (7)) and the exchange current (i0) as follows [42]: 0
Ea
i ¼ Ae RT
ð7Þ
0
i ¼ RT=nFRct
ð8Þ
Where R is the gas constant, T is the absolute temperature, n is the number of transferred electrons, F is the Faraday constant and A is
Table 1 List of recent works on CeO2-C composite as anode for LIBs. Material system
Current density
Discharge capacity (mAh g1)after n cycles
Capacity retention(%)
[Ref.]
Core–shell CeO2@C nanospheres brick-like CeO2 plate-like CeO2 CeO2/CMK-3 micro/nano dumbbell-shaped CeO2 shuttle-shaped CeO2 rhombus-shaped CeO2 micro/nano core–shell sphere CeO2 Core–shell CeO2 micro/nanospheres rectangular prism-like CeO2 microrods CeO2@C hollow spheres Graphite Hard carbon
0.2C 200 mA g1 200 mA g1 100 mA g1 0.2 mA cm2 0.2 mA cm2 0.2 mA cm2 0.2 mA cm2 0.6 mA cm2 0.2 mA cm2 1000 mA g1 120 mA g1 20 mA g1
355, n = 50 460, n = 100 290, n = 100 360, n = 100 590, n = 100 315.3, n = 50 374, n = 50 546.7, n = 300 192.2, n = 100 316, n = 100 308, n = 500 283, n = 100 286, n = 100
41 34 25 48 43 59.9 70.7 78.8 50.3 71.72 100 95 90
2014 [11] 2014 [35] 2014 [35] 2016 [20] 2016 [37] 2016 [38] 2016 [38] 2016 [36] 2017 [19] 2017 [39] This work 2016 [40] 2016 [41]
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reversible specific capacity of 313 mA h g1 at 1000 mA g1, 210 mA h g1 at 4000 mA g1 and 169 mA h g1 at 8000 mA g1, with amost 100% retention at the end of 500 discharge-charge cycles, which was obviously superior to CeO2 hollow spheres. The improved cycling performance was attributed to the homogeneous carbon, the Ce3+ in the oxygen-deficient fluorite-like CeO2-x on the surface of ceria nanoparticles and unique hollow sphere structure merits, which improved the conductivity of CeO2 hollow spheres, enhanced the rate of Li+ diffusion, suppressed the aggregation of active particles and reduced the apparent activation energy to enhance the kinetics of Li+ insertion-extraction. Considering the facile preparation route and the superior electrochemical performance at large current densities, CeO2/C hollow sphere may serve as a promising candidate as anode materials for lithium ion batteries. Acknowledgments
Fig. 7. Arrhenius plots of ln(i0) versus 1/T for CeO2 and CeO2/C anodes.
the temperature independent coefficient. The equations can be rearranged as below [42]: lnði0 Þ ¼
Ea 1 þ lnA RT
ð9Þ
1 ln i0 ¼ k þ b T
ð10Þ
Ea ¼ Rk
ð11Þ
EIS and the fitting results of pristine CeO2 and CeO2/C anodes were carried out at different temperatures (Fig. S8 and Table S2, Supporting Information) and the corresponding values of Rct and ln (i0) have been summarized in Table S3 in Supporting Information. Fig. 7 shows Arrhenius plots of ln(i0) as a function of 1/T of CeO2 and CeO2/C anodes. Based on the linear fitting results, the apparent activation energies of CeO2 and CeO2/C anodes are caculated as 72.33 and 59.03 kJ mol1, respectively. Therefore, the insertion and extraction of Li+ becomes easier in CeO2/C than in CeO2. Clearly, the remarkable electrochemical performance of CeO2/C electrode is originated from the homogeneous carbon and unique hollow sphere structure merits. Firstly, the pyrolyzed homogeneous carbon improves the electronic conductivity of the CeO2/C electrode, which favors the charge transfer and enhances the Li+ diffusion rate in CeO2 nanocrystals, leading to the better high-rate performance of CeO2/C. Secondly, the uniform carbon that acts as a robust three-dimensional network and the free volume in hollow structure alleviate the stress-induced material failure arising from the volume expansion in the repeating ionic intercalation process, which improves the cycling stability of CeO2/C component. More importantly, the homogeneous carbon with 3D network structure gives birth to more active sites for Li+ storage and reduces the apparent activation energy, acting as the role of enhancing the kinetics, facilitating the ions insertion-extraction and increasing the capacity of CeO2/C. 4. Conclusions Porous CeO2/C hollow spheres have been prepared via microwave hydrothermal process by using citric acid as carbon precursor and being annealed under Ar atmosphere. It was found that the presence of citric acid not only facilitated the formation of hollow sphere morphology, but also led to a homogeneous carbon network structure. The obtained CeO2/C samples were evaluated for their lithium storage properties. These electrodes delivered a
This work was financially supported by the National Natural Science Foundation of China (grant no. 51472198), the Fundamental Research Funds for Central Universities (granted no. xkjc2014009), the Scientific Research Foundation for the Returned Overseas Chinese Scholars, State Education Ministry, and the Fundamental Research Funds for the Central Universities (grant no. XJJ2016020). We thank Miss Liu Jiamei at Instrument Analysis Center of Xi’an Jiaotong University for her assistance with XPS measurement. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at https://doi.org/10.1016/j.electacta.2017.10.041. References [1] D. Larcher, J.M. Tarascon, Towards greener and more sustainable batteries for electrical energy storage, Nat. Chem. 7 (2015) 19–29. [2] J. Lee, A. Urban, X. Li, D. Su, G. Hautier, G. Ceder, Unlocking the Potential of Cation-Disordered Oxides for Rechargeable Lithium Batteries, Science 343 (2014) 519–522. [3] X. Zhang, M. van Hulzen, D.P. Singh, A. Brownrigg, J.P. Wright, N.H. van Dijk, M. Wagemaker, Direct view on the phase evolution in individual LiFePO4 nanoparticles during Li-ion battery cycling, Nat. Commun. 6 (2015). [4] Y. Wang, W.D. Richards, S.P. Ong, L.J. Miara, J.C. Kim, Y. Mo, G. Ceder, Design principles for solid-state lithium superionic conductors, Nat. Mater. 14 (2015) 1026–1031. [5] B. Dunn, H. Kamath, J.-M. Tarascon, Electrical Energy Storage for the Grid: A Battery of Choices, Science 334 (2011) 928–935. [6] X. Meng, Y. Xu, X. Sun, J. Wang, L. Xiong, X. Du, S. Mao, Graphene oxide sheetsinduced growth of nanostructured Fe3O4 for a high-performance anode material of lithium ion batteries, J. Mater. Chem. A 3 (2015) 12938–12946. [7] D. Bresser, F. Mueller, M. Fiedler, S. Krueger, R. Kloepsch, D. Baither, M. Winter, E. Paillard, S. Passerini, Transition-Metal-Doped Zinc Oxide Nanoparticles as a New Lithium-Ion Anode Material, Chem. Mater. 25 (2013) 4977–4985. [8] Q. Su, L. Chang, J. Zhang, G. Du, B. Xu, In Situ TEM Observation of the Electrochemical Process of Individual CeO2/Graphene Anode for Lithium Ion Battery, J. Phys. Chem. C 117 (2013) 4292–4298. [9] H.Y. Kim, H.M. Lee, G. Henkelman, CO Oxidation Mechanism on CeO2Supported Au Nanoparticles, J. Am. Chem. Soc. 134 (2012) 1560–1570. [10] X. Wang, D. Liu, S. Song, H. Zhang, Pt@CeO2 Multicore@Shell Self-Assembled Nanospheres: Clean Synthesis, Structure Optimization, and Catalytic Applications, J. Am. Chem. Soc. 135 (2013) 15864–15872. [11] X. Wu, H. Niu, S. Fu, J. Song, C. Mao, S. Zhang, D. Zhang, C. Chen, Core-shell CeO2@C nanospheres as enhanced anode materials for lithium ion batteries, J. Mater. Chem. 2 (2014) 6790–6795. [12] G. Shen, Q. Wang, Z. Wang, Y. Chen, Hydrothermal synthesis of CeO2 nanooctahedrons, Mater. Lett. 65 (2011) 1211–1214. [13] N. Izu, S. Nishizaki, W. Shin, T. Itoh, M. Nishibori, I. Matsubara, Resistive Oxygen Sensor Using Ceria-Zirconia Sensor Material and Ceria-Yttria Temperature Compensating Material for Lean-Burn Engine, Sensors 9 (2009) 8884–8895. [14] M. Shen, F. Lin, G. Wei, J. Wang, S. Zhu, Improved sulfur-resistant ability on CO oxidation of Pd/Ce0.75Zr0.25O2 over Pd/CeO2-TiO2 and Pd/CeO2, J. Rare Earths 33 (2015) 56–61.
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Electrochimica Acta journal homepage: www.elsevier.com/locate/electacta
Research paper
Porous carbon-wrapped cerium oxide hollow spheres synthesized via microwave hydrothermal for long-cycle and high-rate lithium-ion batteries Mingbo Maa,b , Hongjie Wanga,* , Sen Liangb , Shengwu Guoa , Yuan Zhangb , Xianfeng Dub,* a
State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, PR China Electronic Materials Research Laboratory, Key Laboratory of the Ministry of Education & International Center for Dielectric Research, Xi’an Jiaotong University, Xi’an 710049, PR China b
A R T I C L E I N F O
Article history: Received 5 July 2017 Received in revised form 30 September 2017 Accepted 5 October 2017 Available online 7 October 2017 Keywords: CeO2/C hollow spheres high-rate performance anodes lithium-ion batteries
A B S T R A C T
Porous hollow structures have attracted tremendous interests due to their geometrical beauty, unique structural features and fascinating physicochemical properties. In the present work, we developed a general procedure for the synthesis of hollow nanostructured carbon-wrapped cerium oxide (CeO2/C) via a microwave hydrothermal process without any surfactants or hard templates. The electrochemical performances of CeO2/C specimen were tested as anode materials in lithium ion batteries. The results showed that CeO2/C hollow spheres exhibited an exceptional cyclic stability with a high reversible capacity of 313 mA h g1 after 500 cycles at a high current density of 1000 mA g1 without signs of further degradation and also presented an excellent rate performance from 1000 to 10000 mA g1. The improved performances were attributed to the homogeneous carbon with 3D network structure in hollow spheres and the Ce3+ in the oxygen-deficient fluorite-like CeO2-x on the surface of ceria nanoparticles, which enhanced the conductivity of CeO2 hollow spheres, suppressed the aggregation of active particles and reduced the apparent activation energy to facilitate the kinetics of Li+ insertion-extraction. The unique hollow structure of CeO2 wrapped by conductive carbon have made CeO2 a promising candidate for future applications in various metal oxide electrodes to mitigate the mechanical degradation and capacity fading, critical for developing advanced electrochemical energy storage systems with long-cycle life and high-rate performance. © 2017 Elsevier Ltd. All rights reserved.
1. Introduction Lithium ion batteries (LIBs) have been widely used in portable electronics, electric tools and electric vehicles [1–3] due to their superior properties such as high energy density, long cycle life and good environmental compatibility [4,5]. In order to satisfy the ever-growing demands on high-capacity and high-power LIBs, great efforts have been made on transition metal oxides [6,7], which possess much higher capacity compared with widely-used graphite in commercial LIBs. However, the main challenges in the implementation of metal oxide electrodes are their large volume variation and poor electronic conductivity during the charge/ discharge process [8]. As a kind of important transition metal oxides, CeO2 shows fluorite structure [9,10], and the oxidation state
* Corresponding authors. E-mail addresses: [email protected] (H. Wang), [email protected] (X. Du). https://doi.org/10.1016/j.electacta.2017.10.041 0013-4686/© 2017 Elsevier Ltd. All rights reserved.
of cerium can mutate quickly between Ce(III) and Ce(IV) [8,11], leading to its wide applications in solid-state fuel cells [12], oxygen storage materials [13] and catalytic supporters [14]. The CeO2 nanoparticles displayed imperceptible volumetric and morphological changes during the lithiation/delithiation process [8], which made it a possible candidate as an anode material for LIBs in the future in case that its electronic conductivity can be improved greatly. It is well known that carbon materials have high electronic conductivity that can effectively improve the conductance of active materials and always act as buffering matrix to suppress the aggregation of active particles in the electrode [11]. So numerous efforts have been made to synthesize carbon material/metal oxides complexes for LIBs, such as SnO2@carbon composite [15], grapheme-encapsulated Fe3O4 [16] and Co3O4-graphene nanoflowers [17]. During the passed couple of years, a variety of CeO2 and carbon composites with different morphologies and microstructures, including nanocomposite [18], nanospheres [11,19] and
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nanorods [20] have been synthesized to enhance its electronic conductivity as anode materials in LIBs. Specially, CeO2-graphene nanocomposite prepared by hydrothermal method exhibited improved specific capacity of 605 mA h g1 at 50 mA g1 after 100 cycles [18]. Core-shell CeO2@C nanospheres synthesized by hydrothermal carbonization showed an initial discharge specific capacity of 863 mA h g1 in the potential range of 3.0-0.0V [11]. Graphene/CeO2/CMK-3 composites with a hierarchical structure prepared by a template method maintained the specific capacity of about 550 mA h g1 at 100 mA g1 after 100 cycles [20]. All these CeO2 materials achieved good electrochemical performances at the expense of the complex preparation process and long synthesizing time [21–23]. Moreover, electrochemical performances of CeO2 have not been investigated at large current densities for lithium ion battery in previous reports. Herein, we report a simple microwave hydrothermal method by using gas templates to synthesize hollow spheres composed by carbon-wrapped CeO2 nanoparticles (noted as CeO2/C). The structure, shape and formation process of CeO2/C were investigated by SEM, TEM, XRD, Raman and BET. And we further demonstrated their promising properties as anode materials for lithium-ion batteries. The results showed that CeO2/C hollow spheres exhibited improved discharge capacity and remarkable long-cycle stability (313 mA h g1 at 1000 mA g1 with 100% retention after 500 cycles), and excellent high-rate capacity (even at the large current density of 10000 mA h g1).
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2. Experimental 2.1. Materials All chemicals were of analytical grade and used as received without further purification. Cerium nitrate hexahydrate (Ce (NO3)36H2O), citric acid monohydrate (C6H8O7H2O), polyvinylidene fluoride (PVDF) and N-methyl-2-pyrrolidone(NMP) were purchased from Sinopharm Chemical Reagent Co., Ltd., China. And deionized (DI) water was produced by Milli-Q-Reference water system (Millipore Co., USA). 2.2. Preparation of CeO2/C hollow spheres CeO2 precursor was synthesized through microwave hydrothermal process. Typically, 0.1 mol Ce(NO3)36H2O and 0.05 mol C6H8O7H2O were dissolved in 1 L DI water at room temperature. After being stirred for 4 h, the obtained solution was transferred into quartz tube and then maintained at 150 C for 2 min under microwave condition with the wattage between 95 and 130 W (Initiator+TM Microwave System, Biotage Co., Sweden). Fig. S1 recorded the profiles of temperature (T), pressure (p) and power (P) during microwave hydrothermal process. After the system cooled totally, the solution became turbid. Then the obtained solution was filtrated, washed and dried to get the light-yellow precipitate of hollow sphere marked as CeOC. Finally, the CeOC
Fig. 1. SEM images of (a1-a3) CeOC, (b1-b3) CeO2 and (c1-c3) CeO2/C; EDS element mapping images of (a4-a6) CeOC, (b4-b6) CeO2 and (c4-c6) CeO2/C.
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samples annealed at 300 C for 1 hour in air became yellow and were noted as CeO2 while the ones pyrolyzed at 600 C for 1 hour in Ar flow became black and were denoted as CeO2/C. 2.3. Structure characterization Morphology and composition analysis of all these samples were observed using field-emission scanning electron microscopy with energy-dispersive spectrometer (FE-SEM, FEI6000, USA) and highresolution transmission electron microscopy (HR-TEM, JEM-2100F, Japan). Their crystalline structures were characterized by X-ray diffraction (XRD, PANalytical, X’pert Pro, Holland), Raman spectroscopy (Horiba HR800, France) and X-ray photoelectron spectroscopy (XPS, Thermo Fisher Scientific, ESCALAB Xi+, America). The weight changes during heat-treatment process were investigated through thermogravimetry analysis (TGA, METTLER TOLEDO, Switzerland). The electronic conductivity was estimated in a potentiostatic polarization method with the polarizing voltage of 0.5 V. The specific surface area was measured by BrunauerEmmett-Teller (BET, ASAP 2020, Micromeritics, USA). The pore size distributions were determined by Barrett-Joyner-Halenda (BJH, ASAP 2020, Micromeritics, USA) at liquid nitrogen temperature. The mass of samples was weighted by high-precision analytical balance (XS105DU, METTLER TOLEDO, Switzerland).
spectroscopy (EDS) shown in Fig. 1. It is clear to see that all these samples show hollow sphere structures with diameter distribution within 0.5–2 mm and wall thickness of about 100 nm. The free volume in hollow structure can effectively accommodate the volume changes and enhance the cyclic stability of active materials [24]. The schematic diagram of possible formation process for hollow sphere structures has been proposed in Fig. 2. Above all, Ce (III) combined with citric acid to form the complex cerium citrate (Ce(III)-citrate), of which the part was decomposed to be carbon dioxide and water at the elevated temperatures and pressure (Fig. S1). Meanwhile, Ce(III) was oxidized by oxygen and/or nitrates involved in the reaction system to synthesize poorly-crystallized CeO2[21]. Subsequently, the rest organic cerium compound and the obtained CeO2 were attached to the surface of gas templates. Two minutes later, CeOC hollow spheres can be obtained only in the Ce (NO3)3 and C6H8O7 mixed solution when the pressure reached to 20 bar (Fig. S2). The main reactions involved in the reaction system could be illustrated as follows[21]: (1)
(2)
2.4. Electrochemical measurements Electrochemical experiments were performed by using 2016 type coin cells. The test electrodes were prepared by adding active material, acetylene black and polyvinylidene fluoride (PVDF) binder into N-methy-l-2-pyrrolidone (NMP) solution with the weight ratio of 60: 30: 10, followed by grinding the mixture into homogeneous slurries. Then, the slurries were pasted on the Cu foil and dried at 120 C in a vacuum oven. The testing coin-type cells were assembled in an argon filled glove box with as-prepared electrodes, Li foil as counter and reference electrodes, a Celgard film (America) separator and 1 M LiPF6 in 1:1:1 (by volume) ethylene carbonate (EC)-propylene carbonate (PC)-dimethyl carbonate (DMC) electrolyte (Capchem Technology Co., Ltd. Shenzhen, China). All the tests were conducted on the battery tester (CT2001A, LAND, China; BT2000, Arbin, USA) within 0.01-3.00 V vs. Li/Li+ at various current densities at room temperature. The weight of electrode material was 0.7–1 mg cm2. The cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) measurements were carried out on an electrochemical workstation (VMP2, Princeton, USA) in a three-electrode configuration with CeO2/Cu as working electrode and Li foil as counter and reference electrode. 3. Results and discussion The morphologies and element mapping images of CeOC, CeO2 and CeO2/C were observed by FE-SEM with energy disperse
4HNO3 ! 4NO2 + O2 + 2H2O
(3)
Citric acid not only served as a carbon precursor for CeO2/C hollow spheres but also promoted the formation of hollow spherical shape in the process. Fig. 1a1 and a3 show the smooth morphology of CeOC with homogenously distributed C, O and Ce elements shown in Fig. 1a4 to a6. After being annealed in air, carbon in CeOC was totally burned up and the observed homogeneously-distributed carbon with small amount in Fig. 1b1 to b6 apparently came from the ambient environment. The surface of CeO2 hollow spheres became much rougher because of the removed carbon from CeOC. While C element partly remained in CeO2/C as a result of the pyrolyzed CeOC in Ar, it was quite useful to improve the electronic conductivity of CeO2. What’s more, the surface of CeO2/C hollow spheres become less rougher compared with CeO2 in Fig. 1c1 to c6. Fig. 3 shows microstructures and phase compositions of CeOC, CeO2 and CeO2/C observed by TEM. A strong contrast between paler center and darker edge in all these samples indicates their hollow interior with some differences in the surface shown in Fig. 3a, d and 3g. Specially, CeOC can not be easily penetrated by electron beam for its smooth surface and thick wall. CeO2 shows the porous morphology with thin wall and some black particles distributed on its surface. CeO2/C wrapped by carbon layer presents the middle morphology features between CeOC and
Fig. 2. Schematic diagram of the formation of CeO2 and CeO2/C hollow spheres fabricated through a microwave hydrothermal process.
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Fig. 3. HR-TEM images and SEAD patterns of (a-c) CeOC, (d-f) CeO2 and (g-i) CeO2/C.
CeO2. Then the selected area electron diffraction (SEAD) patterns were conducted to describe the crystallinity and phases of these samples. Several rings in Fig. 3b, e and h reveal their polycrystalline nature. The weak rings in Fig. 3b indicate the low crystallinity of CeOC. The rings made up of points in Fig. 3e demenstrate the wellcrystalized nanocrystalline structure of CeO2. In Fig. 3h, the rings constituted with points suggests the nanocrystalline nature of CeO2/C. The smallest ring is very broad and bright due to the overlap of (111) plane of CeO2 and (002) plane of C, further proved by the high-resolution TEM (HR-TEM) image below. HR-TEM image in Fig. 3c discloses the existence of nanocrystals and amorphous state in CeOC. The lattice fringe of these nanocrystals is 0.308 nm, corresponding to the (111) plane of CeO2. The amorphous substance may originate from the decomposition product of citric acide. The lattice fringes of 0.308 nm in Fig. 3f illustrate the existence of CeO2 nanocrystals in CeO2 samples. The nanoparticles with lattice fringes of 0.308 nm and crooked ones of 0.33 nm are corresponding to (111) plane of CeO2 and (002) plane of C respectively in Fig. 3i. More importantly, carbon wrapping up on surfaces and interfaces of CeO2 nanocrystals can favour the
improvement of the electronic conductivity of CeO2/C and promote rapid charge-transfer reactions in lithium ion batteries. Fig. 4a and b display the phase compositions of CeOC, CeO2 and CeO2/C carried out by powder XRD and Raman spectrum. Obviously, no peaks of CeOC appear in Fig. 4a, indicating its low crystallinity and small crystalline grains. On the contrary, both CeO2 and CeO2/C display the same characteristic peaks at 28.5 , 33.1, 47.5 , 56.3 , 59 , 69 , 76.7 and 79 , which match up well with the CeO2 structure. Moreover, Raman active mode at 465 cm1 in Fig. 4b demonstrates the existence of cubic CeO2 nanocrystals [25]. In addition, CeO2/C exhibits broad bands around the 1200– 1500 cm1 region, corresponding to carbon-based Rama vibration modes. The peaks at 1350 and 1560 cm1 can be attributed to the disordered D-band carbon and the graphitic G-band carbon [26], respectively. In brief, the result of Raman analysis further confirms the co-existence of carbon and CeO2 in our CeO2/C samples. In addition, XPS analysis was carried out to further confirm the phase compositions of CeO2 and CeO2/C and the results were shown in Fig. S3 and Table S1. It is obvious to see the existence of Ce3+ in the surface of these samples, which is a result of oxygen vacancies and enhanced in cerium oxide nanoparticles [27,28]. The presence of
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Fig. 4. (a) XRD patterns and (b) Raman spectra of CeO2 and CeO2/C; (c) TGA curves and (d) the evolution of current with polarization time under the polarization voltage of 0.5 V in CeO2 and CeO2/C; N2 adsorption-desorption isotherms and pore-size distribution of (e) CeO2 and (f)CeO2/C.
Ce3+ implies that the oxygen-deficient fluorite-like CeO2-x is located on the surface of ceria nanoparticles [29,30]. The transformation of Ce4+ to Ce3+ in CeO2-x driven by oxygen vacancies will result in the significant increase in the electronic conductivity of CeO2/C. In Fig. 4c, TGA analysis was carried out to determine the mass changes of CeOC in different gas and the weight percentage of carbon in CeO2/C. The mass of CeOC remains unchanged after 300 C in air and after 600 C in Ar. Obviously, 15 wt% carbon in CeO2/C may improve the electronic conductivity of CeO2/C, which has been estimated by potentialstatic polarization method. The corresponding current was measured as function of the polarization period and their relationship is shown in Fig. 4d. The electronic conductivity (se) of CeO2 and CeO2/C can be calculated from the following equation [31]:
se ¼
4Il
pD2 U
ð4Þ
Where U is the polarization voltage, l is the thickness of the sample, D is the diameter of Cu electrode, and I is the current, respectively. The result of calculation shows that the electronic conductivity is increased by almost 40 times from 1.98*109 S cm1 of CeO2 to 8.71*108 S cm1 of CeO2/C. The significantly increased electronic conductivity of CeO2/C can be attributed to the oxygendeficient fluorite-like CeO2-x and the carbon, which will favor its enhanced electrochemical performance as anode for LIBs. In order to evaluate the specific surface area and the pore structures of CeO2 and CeO2/C, the nitrogen adsorption-desorption measurements were implemented and the results are shown in Fig. 4e and f. The isotherm of CeO2 presents a distinct hysteresis loop, which belongs to type IV[32] plot on the basis of the IUPAC classification, indicating the presence of mesoporous structure. Its BET specific surface area is about 85.13 m2 g1. The inset in Fig. 4e shows the diameters of the mesopores are less than 10 nm which is calculated from the adsorption branch. The isotherm of CeO2/C exhibits a type II plot [33] with the decreased BET specific surface area of 9.56 m2 g1 due to the formation of the carbon phase that filled the partial pores and holes in CeO2/C. The pore diameters of most mesopores in Fig. 4f inset are less than 10 nm. Clearly,
variation of electronic conductivity and specific surface area play an important role in the electrochemical properties of anode materials. As for carbon-wrapped CeO2 specimen, the specific surface area decreased by 8 times while the electronic conductivity increased by 40 times. It should be mentioned that the decreased specific surface area of CeO2/C samples did not have any negative impacts on their electrochemical performances, which instead might be tremendously enhanced due to the significantly increased electronic conductivity that should make much more contributions than any other factors to the outstanding electrochemical properties of CeO2/C. Otherwise, the small specific surface area also promote the initial coulombic efficiency of CeO2/C as anode for LIBs by decreasing the amount of solid electrolyte interphase (SEI) films. All these conclusions will be further discussed and proved in the section of electrochemical analysis. The electrochemical impedance spectroscopy (EIS) was utilized to further study the charge transfer and ion diffusion kinetics of CeO2 and CeO2/C electrodes. Both Nyquist plots in Fig. 5 exhibit a semicircle in the high frequency region and a straight line in the low frequency region. The semicircle is ascribed to the SEI film resistance and charge transfer resistance [34] while the straight
Fig. 5. EIS Nyquist plots CeO2 and CeO2/C electrodes obtained at room temperature.
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Fig. 6. Electrochemical performances of CeO2 and CeO2/C. CV curves of (a) CeO2 and (b) CeO2/C at 0.1 mV s1; the initial two galvanostatic discharge-charge profiles for (c) CeO2 and (d) CeO2/C at 100 mA g1; (e) cycle performance at 1000, 4000 and 8000 mA g1; (f) rate capability tests from 1000 to 10000 mA g1.
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line (Warburg line) is attributed to Li-ion diffusion at the interface of electrolyte and electrode [34]. Compared with CeO2 electrode, the semicircle of CeO2/C electrode shows a reduced diameter, indicating the decreased charge transfer (Rct) resistance at the interface of electrode and elelctrolyte due to the homogeneous conductive carbon. The electrochemical performances of these electrodes are evaluated and presented in Fig. 6. Fig. 6a and b which show the cyclic voltammetry (CV) curves of CeO2 and CeO2/C at a scan rate of 0.1 mV s1, respectively. Three obvious cathodic peaks at about 1.16, 0.9 and 0.02 V for CeO2 during the first discharge process (Fig. 6a). In particular, the peak at 1.16 V may originate from the decomposition of the electrolyte on the partially lithiated electrode [32] and disappear in the first anodic scan. The peak at 0.9 V corresponds to the formation of SEI film and the reduction reaction of CeO2/CeO2-x to Ce as described in function (5) [19,35], while the reduction peak at 0.05 V represents the insertion of Li+ into carbon[36]. In the first anodic scan, the observed broad peak at 1.3 V on behalf of the oxidation reaction of Ce to CeO2/CeO2-x suggests the corresponding reversible reaction[35]. The oxidation peak at 0.1 V could be assigned to the extraction of Li+ from carbon. In the second cycle, the reduction peak of CeO2/CeO2-x shifts to 1.0 V and becomes broad while the oxidation peak of Ce shifts to 1.25 V, revealing that the insertion/extraction of Li+ in CeO2/CeO2-x is a reversible process. The CeO2/C electrode presents the similar electrochemical behaviors as given in Fig. 6b. CeO2 + 4Li+ + 4e = Ce + 2Li2O
(5)
CeO2-x + 2(2-x)Li+ + 2(2-x)e = Ce + (2-x)Li2O
(6)
Fig. 6c and d presents the initial discharg-charge voltage profiles of CeO2 and CeO2/C electrodes at the current density of 100 mA g1 within a cut-off voltage window of 0.01-3.0 V. The initial discharge/charge specific capacities are found to be 749/ 418 mA h g1 for CeO2/C, much higher than that of CeO2 (355/ 176 mA h g1), which benefits from the introduction of carbon and the uniform structure between C and CeO2 in CeO2/C. Futhermore, the irreversible capacity in the first cycle should be mainly attributed to the formation of SEI layer on the electroactive materials and the electrolyte decompositon. Because of the decreased specific surface area, the initial coulombic efficiency of the CeO2/C electrode is 60.5%, much higher than that of the CeO2 electrode (48%) in Fig. 6e. After several cycles, the coulombic efficiency increased to nearly 100% for both electrodes. From the second cycle onwards, both electrodes showed good reversibility. At the end of 500 discharge-charge cycles, CeO2/C electrodes still retained a reversible specific capacity of 313 mA h g1 (1000 mA g1), 210 mA h g1 (4000 mA
g1) and 169 mA h g1 (8000 mA g1) and exhibited outstanding capacity and cycle stability. Compared with the current works, CeO2/C hollow sphere shows excellent properties at large current densities from Table 1. However, CeO2 electrodes with low conductivity only retained 121 mA h g1 (1000 mA g1), 113 mA h g1 (4000 mA g1) and 47 mA h g1 (8000 mA g1) at the end of 500 discharge-charge cycles. CeO2/C electrodes also exhibit high specific capacity and good cycling stability at the current densities of 100, 200 and 500 mA g1 in Fig. S4. When changed the upper voltage limits of the voltage window, the specific capacity and the average coulombic efficiency CeO2/C at 1000 mA g1 have no obvious changes. While the lower voltage limit of the voltage window is changed, the initial discharge specific capacities and the initial coulombic efficiency CeO2/C at 1000 mA g1 are changed obviously (Fig. S5). After cycling, the CeO2/C electrode maintained its hollow sphere structure without any obvious changes in Fig. S6, which demonstrated that porous CeO2/C hollow sphere with a steady structure resulted in the excellent cycling stability as electrodes. The rate performances of CeO2 and CeO2/C were shown in Fig. 6f. Profiting from the unique structure and synergetic effect, CeO2/C presents better reversibility and higher specific capacity than CeO2 at different current rates, especially at high discharge/ charge rates. For example, CeO2/C electrode delivers specific capacity of 266, 232, 210, 188 and 169 mA h g1 at 1000, 2000, 4000, 8000 and 10000 mA g1. Whereas the values for CeO2 are 100, 84, 75, 68 and 58 mA h g1, respectively. What’s more, after being tested for 20 cycles under various current rates (even up to 10000 mA g1), the capacities of CeO2/C electrodes can recover to 300 mA h g1 at 1000 mA g1 and remain unchangeable after 500 cycles, indicating the high reversibility of lithium ion insertion/ extraction in CeO2/C. Moreover, the increased specific capacity of CeO2/C electrodes after cycling at 10000 mA h g1 due to the actived electrode enlarges the reduction peak in the CV curve and decreases the charge transfer (Rct) resistance of CeO2/C electrodes in Fig. S7. This phenomenon also appears in core–shell CeO2@C nanospheres [11]. To further illuminate the excellent performance of CeO2/C electrode, the apparent activation energies (Ea) of CeO2 and CeO2/C were calculated through the relationship between Arrhenius Equation (Equation (7)) and the exchange current (i0) as follows [42]: 0
Ea
i ¼ Ae RT
ð7Þ
0
i ¼ RT=nFRct
ð8Þ
Where R is the gas constant, T is the absolute temperature, n is the number of transferred electrons, F is the Faraday constant and A is
Table 1 List of recent works on CeO2-C composite as anode for LIBs. Material system
Current density
Discharge capacity (mAh g1)after n cycles
Capacity retention(%)
[Ref.]
Core–shell CeO2@C nanospheres brick-like CeO2 plate-like CeO2 CeO2/CMK-3 micro/nano dumbbell-shaped CeO2 shuttle-shaped CeO2 rhombus-shaped CeO2 micro/nano core–shell sphere CeO2 Core–shell CeO2 micro/nanospheres rectangular prism-like CeO2 microrods CeO2@C hollow spheres Graphite Hard carbon
0.2C 200 mA g1 200 mA g1 100 mA g1 0.2 mA cm2 0.2 mA cm2 0.2 mA cm2 0.2 mA cm2 0.6 mA cm2 0.2 mA cm2 1000 mA g1 120 mA g1 20 mA g1
355, n = 50 460, n = 100 290, n = 100 360, n = 100 590, n = 100 315.3, n = 50 374, n = 50 546.7, n = 300 192.2, n = 100 316, n = 100 308, n = 500 283, n = 100 286, n = 100
41 34 25 48 43 59.9 70.7 78.8 50.3 71.72 100 95 90
2014 [11] 2014 [35] 2014 [35] 2016 [20] 2016 [37] 2016 [38] 2016 [38] 2016 [36] 2017 [19] 2017 [39] This work 2016 [40] 2016 [41]
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reversible specific capacity of 313 mA h g1 at 1000 mA g1, 210 mA h g1 at 4000 mA g1 and 169 mA h g1 at 8000 mA g1, with amost 100% retention at the end of 500 discharge-charge cycles, which was obviously superior to CeO2 hollow spheres. The improved cycling performance was attributed to the homogeneous carbon, the Ce3+ in the oxygen-deficient fluorite-like CeO2-x on the surface of ceria nanoparticles and unique hollow sphere structure merits, which improved the conductivity of CeO2 hollow spheres, enhanced the rate of Li+ diffusion, suppressed the aggregation of active particles and reduced the apparent activation energy to enhance the kinetics of Li+ insertion-extraction. Considering the facile preparation route and the superior electrochemical performance at large current densities, CeO2/C hollow sphere may serve as a promising candidate as anode materials for lithium ion batteries. Acknowledgments
Fig. 7. Arrhenius plots of ln(i0) versus 1/T for CeO2 and CeO2/C anodes.
the temperature independent coefficient. The equations can be rearranged as below [42]: lnði0 Þ ¼
Ea 1 þ lnA RT
ð9Þ
1 ln i0 ¼ k þ b T
ð10Þ
Ea ¼ Rk
ð11Þ
EIS and the fitting results of pristine CeO2 and CeO2/C anodes were carried out at different temperatures (Fig. S8 and Table S2, Supporting Information) and the corresponding values of Rct and ln (i0) have been summarized in Table S3 in Supporting Information. Fig. 7 shows Arrhenius plots of ln(i0) as a function of 1/T of CeO2 and CeO2/C anodes. Based on the linear fitting results, the apparent activation energies of CeO2 and CeO2/C anodes are caculated as 72.33 and 59.03 kJ mol1, respectively. Therefore, the insertion and extraction of Li+ becomes easier in CeO2/C than in CeO2. Clearly, the remarkable electrochemical performance of CeO2/C electrode is originated from the homogeneous carbon and unique hollow sphere structure merits. Firstly, the pyrolyzed homogeneous carbon improves the electronic conductivity of the CeO2/C electrode, which favors the charge transfer and enhances the Li+ diffusion rate in CeO2 nanocrystals, leading to the better high-rate performance of CeO2/C. Secondly, the uniform carbon that acts as a robust three-dimensional network and the free volume in hollow structure alleviate the stress-induced material failure arising from the volume expansion in the repeating ionic intercalation process, which improves the cycling stability of CeO2/C component. More importantly, the homogeneous carbon with 3D network structure gives birth to more active sites for Li+ storage and reduces the apparent activation energy, acting as the role of enhancing the kinetics, facilitating the ions insertion-extraction and increasing the capacity of CeO2/C. 4. Conclusions Porous CeO2/C hollow spheres have been prepared via microwave hydrothermal process by using citric acid as carbon precursor and being annealed under Ar atmosphere. It was found that the presence of citric acid not only facilitated the formation of hollow sphere morphology, but also led to a homogeneous carbon network structure. The obtained CeO2/C samples were evaluated for their lithium storage properties. These electrodes delivered a
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