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Red phosphorus confined in N-doped multi-cavity mesoporous carbon for ultrahigh-performance sodium-ion batteries
Journal of Power Sources 450 (2020) 227696
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Red phosphorus confined in N-doped multi-cavity mesoporous carbon for ultrahigh-performance sodium-ion batteries Yuhao Zhang a, Beihou Liu a, Timur Borjigin a, Shubiao Xia b, ***, Xiaofei Yang d, Shuhui Sun c, **, Hong Guo a, * a
School of Materials Science and Engineering, Green Energy Key Laboratory of All-Solid Ion Battery in Yunnan Province University, Yunnan University, No. 2, Green Lake North Road, Kunming, 650091, China Center for Yunnan-Guizhou Plateau Chemical Functional Materials and Pollution Control, Qujing Normal University, Qujing, 655011, China c � Centre-Energie Mat�eriaux T�el�ecommunications, Institut National de la Recherche Scientifique (INRS), 1650 Boul. Lionel Boulet, J3X 1S2, Varennes, Quebec, Canada b
d
Nanomaterials and Energy Lab, Department of Mechanical and Materials Engineering, Western University, London, Ontario, Canada
H I G H L I G H T S
� N-doped multi-cavity carbon is prepared via the unique self-assembly strategy. � The amorphous red P with high loading over 50 wt% is encapsulated in the NMC. � P@NMC exhibits ultra-long cycling life with 923.7 mAh g 1 after 1000 cycles. � Sodium-ion storage mechanism of P@NMC is dynamically investigated. A R T I C L E I N F O
A B S T R A C T
Keywords: Sodium-ion batteries Red P N-doped multi-cavity mesoporous carbon Electrochemical mechanism
Phosphorus-based materials are one of the most promising anodes due to their high theoretical capacity (2596 mAh g 1) and safe working potential for sodium-ion batteries (SIBs). However, they suffer from severe volume changes (�293–487%) as well as poor electronic conductivity (10 14 S cm 1), resulting in fast capacity decay during cycling. To tackle the aforementioned issues, in this work, N-doped multi-cavity carbon (NMC) connected by porous structured walls is prepared via the surface energy-driven self-assembly strategy. The internal space is divided into multiple small grids connected by multiple porous carbon walls, which helps accommodate the volume change and build interconnected electronic conductive networks. The amorphous red P is encapsulated in the hollow multi-cavity of carbon and forms P@NMC anode. Benefitting from the unique structure in alle viating volume change and enabling fast electron transport, the SIBs assembled with P@NMC anode exhibit excellent cycling stability and outstanding rate capability as well, retaining a high discharge capacity of 923.7 mAh g 1 after 1000 cycles at 0.5 A g 1 and keeping the Coulombic Efficiencies (CEs) as over 98.5%. Further more, the mechanism of sodium-ion storage in P@NMC is dynamically investigated, providing a new perspective for figuring out the electrochemical cycling behavior.
1. Introduction Sodium-ion batteries (SIBs) as the leading alternatives to lithium-ion batteries (LIBs) have been paid a wide range of attention in the field of advanced energy storage because of the abundant and cost-effective sodium source [1–3]. However, compared to that of lithium ions (0.76 Å), the large radius (1.02 Å) of sodium ions makes sodium storage
more difficult for common anode materials. For example, graphite and crystalline silicon are not electrochemically available for SIBs [4–6]. Besides, there are some pervading issues in anode materials for SIBs, such as low electronic conductivity, large volume expansion, pedal re action kinetics, unsatisfactory cycling life, and low capacity output [7–9]. Thereby, despite the big achievements have been made in the anodes of LIBs and have demonstrated that they are compatible to that
* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (S. Xia), [email protected] (S. Sun), [email protected] (H. Guo). https://doi.org/10.1016/j.jpowsour.2019.227696 Received 21 October 2019; Received in revised form 18 December 2019; Accepted 31 December 2019 Available online 29 January 2020 0378-7753/© 2020 Elsevier B.V. All rights reserved.
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of LIBs, seeking for promising next-generation anode materials with high reversible capacity, good rate performance, stability, and long cycle life for SIBs is still challenging. Phosphorus (P) possesses a safe working potential (0.45 V for vs. Naþ/Na) and a high theoretical capacity (2596 mAh g 1) based on the high reaction of phosphorus with sodium to form the Na3P phase, making it stand out from various anode materials [10–14]. Red P ex hibits superior chemical stability at room temperature and lower cost compared with other allotropes. A fly in the ointment is that most of the red P electrodes own poor rate performance, severe capacity decay and poor electrochemical reversibility, which mainly resulted from the huge volume expansion exceeded to 300%, during the repeated trans formation between the red P and the Na3P phase. Furthermore, the repeated volume changes lead to instability of the electronically insu lating solid electrolyte interface (SEI) layer on the surface of the red P electrode, which decelerates the reaction kinetics, resulting in lower CEs and rapid capacity decay. On the other perspective, the ultralow elec tronic conductivity of amorphous red P (10 14 S cm 1) is responsible for the large polarizability, low utilization of active materials, and reduced high rate performance [15–17]. To achieve better performance of red P anodes, coupling with elec trically conductive carbon is an effective way, which allows the multielectrochemical reaction to proceed in a hybrid matrix of distinct ma terial systems. For instance, Kim et al. prepared P/C composite via a ballmilling method and achieved a high discharge capacity of 1260 mAh g 1 at 100 mA g 1 and stably run for 30 cycles [17]. Furthermore, doping heteroatoms such as B, N, S, P, etc. to the carbon architecture, can further increase the electronic conductivity and create abundant active sites [18]. On the other hand, developing the molecular architecture of red P with functional nanostructures under accurate control has also been verified to be effective to improve the electrochemical perfor mance. Particularly, assembling 3D spatial structures can highly over come the shortcomings of bulk materials by raising electrochemical reaction sites and buffering volume change effectively during cycling [19–21]. For instance, Zhou et al. designed a self-standing P/CNTS@rGO hierarchical structure by ball milling method as anode materials for SIB, delivering a high capacity of about 1000 mAh g 1 at 1C after 500 cycles [22]. Nevertheless, the active materials are easy to aggregate during the charging/discharging process by simply mixing the large-sized red P with hollow materials via the ball-milling method, where the P is coated outside of the hollow materials. As an alternative, a vaporization-condensation strategy was developed by infusing the P vapor into hollow structures under high temperatures and formed ho mogenous P@C hybrid materials. However, it should be mentioned that the P contents in most reported red P@C hybrid materials are less than 35 wt%, such as P-CNTs and P-rGO, which can’t meet the demand of energy density [23,24]. In this regard, further exploration of novel architectural carbon for achieving high P contents is of significance. Recently, a yolk-shell interlinked carbon, which contains cavities, mesoporous shells, and microporous cores anchored to the shell, was developed and provided large space for P accommodation and enabled the P@C composites high P contents. Moreover, the interconnect mes oporous shells built continuous electronic conduction networks to make sure fast electron transport through the P/C interface and improved the electrical conductivity of the P@C electrode. Equally important, the abundant porous structure is beneficial for alleviating the large volume change during cycling. Generally speaking, such a unique structure enables the P@C anode with high P content, high electronic conduc tivity and able to accommodate the volume change during cycling as well. Whereas, synthesis of such multi-cavity hollow structure is typi cally related to the template fabrication or multistep/high-cost pro cedures, which is highly related to result of the collapse of hollow structures. Thus, the rational design of multi-cavity hollow structures as P hosts with a low-cost and scalable procedure is urgent and necessary to the development of high-performance SIBs. In this work, we developed a facile, scalable and controllable
fabrication of N-doped C-encapsulated multi-cavity mesoporous frame (labeled as P@NMC) as illustrated in Scheme 1. Based on the afore mentioned merits, the SIBs assembled with as-prepared P@NMC anode perform out-standing charge storage capacities, rate performance as well as cycling stability. The kinetic properties of the P@NMC are investigated to understand the Na storage process in pseudo-capacitance properties. The capacitive contribution gradually increases as the scan rate increases, indicating P@NMC with a larger capacitive contribution is good for Naþ storage and renders a higher rate performance. 2. Experimental section The raw materials, preparation of nano emulsion, synthesis of multicavity polymer spheres (MCP) precursor are shown in the Supporting Information. Synthesis of N-doped multi-cavity nanocarbon (NMC). The surface of the NMC material was coated with a silica shell by a modified Stober process. 0.96 g of CTAB was stirred vigorously for 1 h by adding 30 mL of water. Then, the CTAB solution was added to a mixed alcohol-aqueous solution of 0.24 g of MCP, 150 mL of water, 60 mL of ethanol, and 2.4 mL of ammonia (25 wt%) for 30 min, and then 1.68 mL of tetraethyl orthosilicate was added drop by drop, subsequently stirred at a constant temperature in a 30 � C water bath for 16 h. MCP@mSiO2 was collected by centrifugation and washed several times with an HClalcohol solution to remove CTAB from the silica shell, then dried at 60 � C overnight. NMC@mSiO2 was prepared by heating the product at 800 � C for 2 h in an Ar gas atmosphere. Finally, NMC@mSiO2 was stirred at 50 � C for 24 h in aqueous NaOH to remove the mSiO2 shell to obtain NMC. 2.1. Synthesis of carbon-red phosphorus composites P@NMC The carbon-red phosphorus composites P@NMC was prepared based on a vaporization/condensation method. The commercial red P and the shower-like NMC were uniformly mixed at a mass ratio of 1:1 and placed in one sealed quartz tube under vacuum. The sealed quartz tube was then heated to 500 � C and kept at the temperature for 3 h, after cooling to 250 � C, it kept at the temperature for 24 h to convert white P to red P, and finally after cooling to room temperature, the product was washed several times with CS2 to remove white P and drying under vacuum. 3. Results and discussion The crystal structure of P confined in the NMC is analyzed by X-ray diffraction (XRD). As shown in Fig. 1a, the P@NMC exhibits a similar XRD pattern compared with the NMC and only two characteristic peaks are observed at 23.6� and 44.2� , corresponding to the carbon. No obvious peaks at 15.3� and 33.5� and 57.2� (PDF No. 44-0906), belonging to the red P, are detected, suggesting that the P is wellconfined in the porous structure of NMC and exists as amorphous structure [25,26]. The corresponding Raman spectra are shown in Fig. 1b. As can be seen, there are two Raman peaks of NMC appear at 1336.49 and 1159.65 cm 1, belonging to sp3 type disordered carbon and sp2 type graphitized carbon, respectively, which are also known as D band and the G band. The peak area ratio of D band to G band is adopted to evaluate the graphitization degree of carbon. NMC presents the value of 1.28, indicating the high graphitization degree of NMC. After P impregnation, the D band and G band exhibit no obvious shift and the value shows negligible change, proving that the structure of NMC maintains well, which coincides well with the XRD results. Moreover, no scattering peaks are detected at 300–500 cm 1, further confirming that the P is well-confined in the pores of NMC. Fourier transform infrared (FTIR) spectroscopy is hired to study the surface chemical states be tween the NMC and red P (Fig. 1c). The peak at 1510 cm 1 for the – C. The peak sample of NMC and P@NMC is assigned to the band of C– 2
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Scheme 1. Representative illustration of red P enveloped multi-cavity carbon spheres composite.
Fig. 1. XRD patterns (a), Raman spectra (b), FT-IR spectra (c) and TGA curves (d) of red P, NMC and P@NMC. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
appears at 1062 cm 1 for pristine red P corresponding to the charac teristic P–O due to the exposure of P nanoparticles in the air. For about P@NMC, the intensity of this position shows a little weak, suggesting the amorphous red P is confined in the NMC system. Furthermore, it can be noticed the formed peak at the position of 1001 cm 1 for the P@NMC sample should be attributed to P–O–C bond, suggesting the strong interaction between NMC and red P. There is no obvious characteristic peak of P in the XRD, Raman and FTIR spectra of P@NMC, revealing the
most of amorphous red P is encapsulated in the NMC frame. Meanwhile, static dissolution of red P, NMC and P@NMC is designed to simulate the working environment in batteries. We disperse the red P, NMC, and P@NMC into the liquid electrolyte and stand for 720 h. As shown in Fig. S1 (Supporting Information), the red P and NMC solutions still maintain their intrinsic colors, while the P@NMC solution appears stratification and the supernatant is colorless and transparent. The striking visual evidence illustrates the P@NMC composite can inhabit 3
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the free-diffusion of P from into the electrolyte and avoid unnecessary mass loss of red P during electrochemical cycling, which is beneficial for increasing the active materials utilization and improving the cycling stability. The P content in the P@NMC composite is determined by the thermal gravimetric analysis (TGA), as shown in Fig. 1d. As it can be seen a mass loss of 50.1 wt% is observed between 400 � C and 700 � C. It can be attributed to the P sublimation, suggesting a high P content of 50.1 wt%, which is much higher than most of current reported P@C composite [14, 23,24,27,28]. X-ray photoelectron spectroscopy (XPS) is taken to further analyze the interaction between red P and NMC. The full-spectra scan of NMC and P@NMC are shown in Fig. S2 (a). The deconvolution of the N1s spectrum identifies the surface functionalities of NMC and P@NMC
(Fig. S2 (b)). There are three fitted component peaks at 399.6, 401.1 and 402.2 eV for NMC and P@NMC, which are designated for pyridinic (N-6), pyrrolic/pyridine (N-5) and quaternary nitrogen (N-Q), respec tively. This order is in accordance with the previous reports [29]. The proper N-doping is essential for improving the electronic conductivity of the NMC and make further efforts to optimize the electrochemical per formance of the P@NMC. For the red P, two peaks located at 129.9 and 130.8 eV are observed in the P–P 2p spectrum due to the 2p3/2 and 2p1/2 spin-orbit splitting of P–P bonds [30]. In addition, red P presents a characteristic peak at 134.6 eV, belonging to the P–O bond, which can be explained as the partial oxidization by the air (Fig. S2 (c)). In the case of P@NMC, the signals of P–P bond peaks disappeared, which can be attributed to the encapsulation of P by the NMC and limited penetration
Fig. 2. SEM (a, b), TEM (c-f) images of the prepared NMC and P@NMC; EDX mapping images (g, the element of C, N, P) of the fabricated P@NMC samples. 4
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depth of X-ray. In particular, there is a significant intensity increasing for the peak located at 134.6 eV, demonstrating the formation of P–O–C bond during the vaporization-adsorption process due to the chemical interactions between red P and NMC [31]. Morphology of the NMC is examined by SEM (Fig. 2a). The carbon spheres exhibit smooth outer surfaces with an approximate diameter of 100–200 nm. After P impregnation, the morphology of NMC shows negligible change and the as-prepared P@NMC composite present almost the same morphology compared with NMC (Fig. 2b). The inner hollow space and porous structure of NMC and P@NMC are further demonstrated by the TEM images (Fig. 2c–f). As it can be seen, the surface of the carbon wall is smooth and uniform in thickness about 15 nm for the NMC (Fig. 2c). The interior of the carbon sphere is crosslinked by ultra-thin carbon walls and uniformly divided into a plural ity of small cavities with a size of 7 nm. After loading red P, a large shade area is observed on the TEM image of P@NMC (Fig. 2d). Meanwhile, the grid structure of NMC becomes inconspicuous and the crystal lattice is not detected at all, according to Fig. 2f. The element mappings of C, N, and P clearly demonstrate the homogeneous distribution of red P in NMC (Fig. 2g). Therefore, the amorphous red P is confined in the multicavities of carbon effectively and forms multi-cavity yolk-shell struc tured P@NMC composite. To further study the existing state of red P in NMC, the specific surface area and pore size distribution of red P, NMC and P@NMC composites are tested by N2 adsorption-desorption isotherm test. As shown in Fig. S3, NMC presents a mixture of type-I and type-IV curves, indicating abundant micro-pores (less than 2 nm) and meso-pores (around 7 nm), confirmed by Fig. S3b. Owing to the abundant microand meso-pores, the NMC delivers a large surface area of 730.7 m2 g 1 and an average pore size of 4.2 nm. The presence of these pores can facilitate electrolyte infiltration and rapid transport of Naþ. After P loaded, the specific surface area of the obtained P@NMC composite decreases to 19.4 m2 g 1 and almost all the micro- and meso-pores in P@NMC disappear, further proving the successful P infusion. The electrochemical performance of the P@NMC composite, as an anode, is firstly evaluated by the cyclic voltammetry (CV) testing at a scan rate of 0.1 mV s 1 from 0.01 to 3 V. As shown in Fig. 3a, during the first negative scanning, two peaks are observed at 0.81 V and 0.44 V, corresponding to the electrolyte decomposition to form the SEI, reduc tion of red phosphide to NaxP, respectively. When the voltage is further reduced to near 0.01 V, the current increases sharply, which can correspond to the further reaction of carbon with sodium. Following the positive scanning, the peaks at about 0.39 V can be detected, corre sponding to the stepwise de-sodiation reaction. For the second cycle, the reduction peak at 0.77 V and oxidization peak at 0.39 V are maintained, suggesting the reversible electrochemical reaction between P and Na3P [10–14,23,24]. Fig. 3b shows the Charge/Discharge profiles of the SIBs assembled with P@NMC anode at the 1st, 2nd, 5th and 50th cycles operating at a current density of 0.5 A g 1. As can be seen, the cell de livers a high initial discharge capacity of 2146.4 mAh g 1. Similar to the CV results, there is a large irreversible capacity between the 1st cycle and the 2nd cycle. The capacity sharply drops to 1210.4 mAh g 1 at the 2nd cycle and result in a low initial CE of 56.4%. The low CE results from the decomposition of the electrolyte forming an irreversible SEI film and the generation of Na2C60, which caused rapid capacity decay. Except for the first cycle, the SIBs assembled with P@NMC anode exhibit excellent cycling stability, proving by the overlapped charge/discharge profiles in Fig. 3b. Additionally, high CEs of over 98.7% are achieved. The discharge profile includes the three sloped regions of 1.5 to 0.75 V (vs. Naþ/Na), 0.75 to 0.39 V (vs. Naþ/Na) and 0.39 to 0.01 V (vs. Naþ/Na). The charging curve also exhibits the inclined platform at 0.39, and 0.75 V, respectively. The process corresponds to sodium ion intercalation and de-intercalation process, which is in agreement with the CV analysis. To clarify the Na-ion insertion mechanism of P@NMC electrode, exsitu XRD analysis is performed to investigate the structural evolution at different (dis)charge states (Fig. 4a) with a scan rate of 0.5� min 1.
Fig. 3. Electrochemical performance of the prepared P@NMC for SIBs. (a) CV curves with a scan rate of 0.1 mV s 1. (b) Charge/Discharge curves for the 1st, 2nd, 5th and 50th cycles at current density of 0.5 A g 1.
When P@NMC electrode is continuously discharged from the initial state of 2.0 V–0.75 V, the characteristic peaks of Na2C60 (ICSD: 47-1577) appear in the XRD patterns, and their intensity increases gradually with the decrease of potential. Some peaks of Na2C60 also exist in the charging process without significant weakening such as (111) and (220) plane, while other peaks’ intensity of lattice planes (222) and (440) shows a little declining. This result means the Naþ reacts with NMC to form Na2C60, and some Naþ cannot emerge from Na2C60 during the electrochemical cycle. However, these inert Na2C60 contents may act as buffer medium and contribute greatly to the cycle stability of P@NMC. In fact, according to our test, the pure NMC shows stable sodium storage capacity of about 140 mAh g 1. When discharged to 0.75 V, the peaks belong to Na3P (102), Na3P (110) and Na3P (103) of (ICSD:04-0764) are detected, indicating that the embedded Na-ions are alloyed with P and form Na3P initially. As the depth of discharge increases, the intension of Na3P peaks gradually sharped and arrives at maximum value when the potential reduces to 0.01 V. Subsequently, its intension gradually re cedes with increasing depth of charge, and the peaks of (110) and (103) disappear while the potential reaches 0.75 V. Another interesting phe nomenon also should be attention, the NaP (ICSD: 14009) appears during the discharge process since the potential declines at about 0.4 V 5
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capacities of no more than 50 mAh g 1, while NMC and P@NMC elec trodes deliver a specific discharge capacity of 145.3 and 1076.9 mAh g 1 at 0.5 A g-1, respectively. This result is consistent with the discussion of electrochemical reaction mechanism above. Despite the initial irre versible capacity loss, both electrodes exhibit a steady reversible ca pacity in subsequent cycles. In addition, the rate performance of P@NMC electrode is evaluated under different current densities, as shown in Fig. 5b. P@NMC electrode renders superior rate performance. It delivers an average reversible capacity of 1076.5, 1063.6, 898.3 and 786.2 mAh g 1 at 0.5, 1, 2 and 3 A g 1, respectively. Even under a high current rate of 5 A g 1, an average reversible capacity of 646.3 mAh g 1 is achieved. When the current density is recovered to 0.5 A g 1, the specific capacity returns to 1042.5 mAh g 1. Cyclic stability testing is an important indicator for the electrode. To evaluate the long-term cyclic behavior of P@NMC, the cyclic performance is examined at 0.5 A g 1 for 1000 cycles (Fig. 5c). The electrode delivers a reversible capacity of 923.7 mAh g 1 after 1000 cycles and achieves high CEs over 98%, demonstrating its excellent cyclic stability. The P@NMC exhibits remarkable improvement compared with most recently reported Pbased materials (Table S1) [32–38]. The structure of P@NMC after cycling is detected by TEM (Fig. 5d). Promisingly, after the long-term cycling, the P@NMC maintains its structure well and no breakage is observed, which suggests the strong capability of NMC in accommoda ting the volume change during cycling and clarifies the reason for the excellent cycling stability of P@NMC anode. To further explain the excellent sodium storage performance of P@NMC electrode, the kinetic features of P@NMC electrode are analyzed by the CV test for the survey on pseudo-capacitance properties during the charge-discharge process. As the scan rate from 0.2 increase to 2 mV s 1 (Fig. 6a), the rela tionship between the measured peak currents related to Naþ storage in P@NMC and sweep rates can be presented by the power law [39–41]. i ¼ avb
(1)
Generally, the i and v are the peak current and scan rate, respectively. The value of b confirms the type of Naþ insertion/extraction. The b-value of 0.5 indicates a diffusion-controlled behavior (battery), whereas the bvalue of 1.0 represents a capacitive-controlled behavior (capacitor) [40]. According to the linear relationship between log i and log v plots (Fig. 6b), the slopes for cathodic peaks 1, cathodic peaks 2 and anodic peaks are 0.688, 0.540 and 0.541, respectively, revealing a pseudoca pacitive process between typical behaviors of battery and capacitors. Especially, cathodic peak 1 exhibits higher pseudo capacitive behaviors, which is attributed to the contribution of the porous carbon skeleton to capacitance in P@NMC materials [42]. The capacity contribution from pseudocapacitive and diffusion-controlled charge can be quantified by using the given relationship: i(V) ¼ k1 υ þ k2 υ1/2(2).
Fig. 4. Structure evolution of the electrochemical process: (a) The Ex situ XRD of the discharge-charge profile of P@NMP at different potentials for the 1st cycle at current density of 0.5 A g 1. (b) The contour and response surface analysis corresponding to discharge-charge process. (c) Representative illus tration of the electrochemical reaction process for the P@NMC.
k2 υ1/2 represents diffusion-controlled reactions and k1 υ represents capacitance control [43,44]. Therefore, the relative contributions from the capacitive process and Naþ intercalation can be decoupled (Fig. 6c). At the scan rate of 0.2 mV s 1, the capacitive charge contribution is calculated to be 22.6%, which increases with scan rate and found to be 54.5% at a scan rate of 2 mV s 1. This behavior is not surprising because the pseudo-capacitive contribution should play a critical role in the case of carbon composites, which possesses stable mesopore structures. It should be noted that increasing tendency refers to the contribution ratio rather than the absolute value because pseudo-capacitance is an intrinsic characteristic of a material. The typical voltage profile for the capacitive current (green region) in comparison with the total current is shown in the Fig. 6d, at scan rate of 0.2 mV s 1, a dominating capacitive contribution (~22.6%) is obtained for the P@NMC anode, implying it has a good capacitive contribution, and thus results of the improvement of high capacity and rate performance [44]. The electrochemical impedance spectroscopy (EIS) for three
and it exhibits the same trend as that of Na3P. The highest intensity of NaP peak presents at 0.01 V during the discharge process, while its in tensity becomes weak gradually and then fades away when the potential rise at 0.75 V during the charging process. Figs. S4–S6 exhibit the XRD patterns of Na2C60, Na3P, and NaP when the P@NMC anode discharges to 0.01 V, respectively. Moreover, the contour and response surface analysis corresponding to the discharge-charge process is shown in Fig. 4b, which reflects the (de)alloying process of Na3P and NaP during cycling visually. Therefore, high activity reversible Na3P and NaP ren ders the electrode with a high specific capacity, and the electrochemical reaction process is listed in Fig. 4c. The cyclic performance of commercial red P, NMC and P@NMC electrodes is presented in Fig. 5a. The pure P delivers low specific 6
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Fig. 5. (a) Cycling performance of commercial red P, NMC and P@NMC at a constant current density of 0.5 A g 1. (b) Rate capability of P@NMC from 0.5 A g 1 to 5 A g 1. (c) Cycling capability of P@NMC with prolonged cycle life (1000 cycles) at 0.5 A g 1. (d) TEM image of P@NMC electrode after 1000 cycles. (For inter pretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
Fig. 6. Kinetics analysis of P@NMC electrode. (a) CV curves of P@NMC at different scan rates (0.2–2 mV s 1). (b) Relationship between logarithm redox peak current and logarithm scan rate of the P@NMC. (c) Normalized contribution ratio of capacitive at different scan rates. (d) CV profile of the capacitive contribution at 0.2 mV s 1.
different electrode materials: red P, NMC, and P@NMC are shown in Fig. S7. Table S2 shows the impedance parameters of each impedance component obtained by Z-Simpwin simulation of the equivalent circuit, including the ohmic resistance (Rs), charge transfer resistance (Rct) and Warburg impedance (W). The ohmic resistance (Rs) of P@NMC is simulated as 19.17Ω, which is slightly lower than the pristine red P (20.87Ω). The Rct value of the P@NMC composite electrode is 183.1 Ω,
which is much lower than the charge transfer resistance of red P (616.1 Ω), indicating the enhanced charge transfer dynamics. The improved electrochemical kinetics can be attributed to the interconnected elec tronic conductive network in NMC that enhanced the electronic con ductivity and confined the P as nanoparticles. Overall, such a unique structure of NMC can facilitate electron transport and alleviate the volume change, enabling the high P loading NIBs with faster 7
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electrochemical performance.
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kinetics,
excellent
cycling
stability,
and
rate
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4. Conclusions In summary, a P@NMC composite with high red P content (over 50 wt%) is successfully developed. The multi-cavity NMC acts as an effec tive P host by structural monomers driven by surface free energy. The mesh-like multi-cavity interior provides a large space for P accommo dation as well as alleviate volume change during cycling. Moreover, the interconnected porous carbon walls build an interconnected electronic conductive network for fast electron transport. Nitrogen doping in car bon spheres enhances the conductivity and also provides more active sites for electrochemical reactions. Benefitting from the aforementioned merits of the unique structure, the NIBs assembled with P@NMC anode present excellent electrochemical performance, which maintains a high capacity of 923.7 mAh g 1 after 1000 cycles at a current density of 0.5 A g 1. The electrochemical performance of P@NMC is superior to most of the current phosphorus-carbon anodes. This general strategy can be extended to further explore other advanced energy materials. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements Y.H.Z. and B.H.L. contributed equally to this work. The authors would like to acknowledge financial support provided by Key National Natural Science Foundation of Yunnan Province (No.2018FA028 and No.2019FY003023), National Natural Science Foundation of China (No.51474191 and No.21467030), Major State Basic Research Devel opment Program of China (973 Program, No. 2014CB643406), and the Program for Outstand Young Talents (2018) of Yunnan University. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227696. References [1] W. Chen, T. Li, Q. Hu, C. Li, H. Guo, J. Power Sources 286 (2015) 159–165. [2] W. Wang, Y. Guo, L. Liu, S. Wang, X. Yang, H. Guo, J. Power Sources 245 (2014) 624–629. [3] H. Liu, B. Liu, H. Guo, M. Lianga, Y. Zhang, T. Borjigina, X. Yang, L. Wang, X. Sun, Nano Energy 51 (2018) 639–648. [4] X. Song, X. Li, Z. Bai, B. Yan, D. Li, X. Sun, Nano Energy 26 (2016) 533–540. [5] Z. Liu, X. Yu, X.W. Lou, U. Paik, Energy Environ. Sci. 9 (2016) 2314–2318. [6] S. Wang, L. Xia, L. Yu, L. Zhang, H. Wang, X.W. Lou, Adv. Energy Mater. 6 (2016) 1502217. [7] H. Liu, H. Guo, B. Liu, M. Liang, Z. Lv, K.R. Adair, X. Sun, Adv. Funct. Mater. 28 (2018) 1707480.
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Red phosphorus confined in N-doped multi-cavity mesoporous carbon for ultrahigh-performance sodium-ion batteries Yuhao Zhang a, Beihou Liu a, Timur Borjigin a, Shubiao Xia b, ***, Xiaofei Yang d, Shuhui Sun c, **, Hong Guo a, * a
School of Materials Science and Engineering, Green Energy Key Laboratory of All-Solid Ion Battery in Yunnan Province University, Yunnan University, No. 2, Green Lake North Road, Kunming, 650091, China Center for Yunnan-Guizhou Plateau Chemical Functional Materials and Pollution Control, Qujing Normal University, Qujing, 655011, China c � Centre-Energie Mat�eriaux T�el�ecommunications, Institut National de la Recherche Scientifique (INRS), 1650 Boul. Lionel Boulet, J3X 1S2, Varennes, Quebec, Canada b
d
Nanomaterials and Energy Lab, Department of Mechanical and Materials Engineering, Western University, London, Ontario, Canada
H I G H L I G H T S
� N-doped multi-cavity carbon is prepared via the unique self-assembly strategy. � The amorphous red P with high loading over 50 wt% is encapsulated in the NMC. � P@NMC exhibits ultra-long cycling life with 923.7 mAh g 1 after 1000 cycles. � Sodium-ion storage mechanism of P@NMC is dynamically investigated. A R T I C L E I N F O
A B S T R A C T
Keywords: Sodium-ion batteries Red P N-doped multi-cavity mesoporous carbon Electrochemical mechanism
Phosphorus-based materials are one of the most promising anodes due to their high theoretical capacity (2596 mAh g 1) and safe working potential for sodium-ion batteries (SIBs). However, they suffer from severe volume changes (�293–487%) as well as poor electronic conductivity (10 14 S cm 1), resulting in fast capacity decay during cycling. To tackle the aforementioned issues, in this work, N-doped multi-cavity carbon (NMC) connected by porous structured walls is prepared via the surface energy-driven self-assembly strategy. The internal space is divided into multiple small grids connected by multiple porous carbon walls, which helps accommodate the volume change and build interconnected electronic conductive networks. The amorphous red P is encapsulated in the hollow multi-cavity of carbon and forms P@NMC anode. Benefitting from the unique structure in alle viating volume change and enabling fast electron transport, the SIBs assembled with P@NMC anode exhibit excellent cycling stability and outstanding rate capability as well, retaining a high discharge capacity of 923.7 mAh g 1 after 1000 cycles at 0.5 A g 1 and keeping the Coulombic Efficiencies (CEs) as over 98.5%. Further more, the mechanism of sodium-ion storage in P@NMC is dynamically investigated, providing a new perspective for figuring out the electrochemical cycling behavior.
1. Introduction Sodium-ion batteries (SIBs) as the leading alternatives to lithium-ion batteries (LIBs) have been paid a wide range of attention in the field of advanced energy storage because of the abundant and cost-effective sodium source [1–3]. However, compared to that of lithium ions (0.76 Å), the large radius (1.02 Å) of sodium ions makes sodium storage
more difficult for common anode materials. For example, graphite and crystalline silicon are not electrochemically available for SIBs [4–6]. Besides, there are some pervading issues in anode materials for SIBs, such as low electronic conductivity, large volume expansion, pedal re action kinetics, unsatisfactory cycling life, and low capacity output [7–9]. Thereby, despite the big achievements have been made in the anodes of LIBs and have demonstrated that they are compatible to that
* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (S. Xia), [email protected] (S. Sun), [email protected] (H. Guo). https://doi.org/10.1016/j.jpowsour.2019.227696 Received 21 October 2019; Received in revised form 18 December 2019; Accepted 31 December 2019 Available online 29 January 2020 0378-7753/© 2020 Elsevier B.V. All rights reserved.
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of LIBs, seeking for promising next-generation anode materials with high reversible capacity, good rate performance, stability, and long cycle life for SIBs is still challenging. Phosphorus (P) possesses a safe working potential (0.45 V for vs. Naþ/Na) and a high theoretical capacity (2596 mAh g 1) based on the high reaction of phosphorus with sodium to form the Na3P phase, making it stand out from various anode materials [10–14]. Red P ex hibits superior chemical stability at room temperature and lower cost compared with other allotropes. A fly in the ointment is that most of the red P electrodes own poor rate performance, severe capacity decay and poor electrochemical reversibility, which mainly resulted from the huge volume expansion exceeded to 300%, during the repeated trans formation between the red P and the Na3P phase. Furthermore, the repeated volume changes lead to instability of the electronically insu lating solid electrolyte interface (SEI) layer on the surface of the red P electrode, which decelerates the reaction kinetics, resulting in lower CEs and rapid capacity decay. On the other perspective, the ultralow elec tronic conductivity of amorphous red P (10 14 S cm 1) is responsible for the large polarizability, low utilization of active materials, and reduced high rate performance [15–17]. To achieve better performance of red P anodes, coupling with elec trically conductive carbon is an effective way, which allows the multielectrochemical reaction to proceed in a hybrid matrix of distinct ma terial systems. For instance, Kim et al. prepared P/C composite via a ballmilling method and achieved a high discharge capacity of 1260 mAh g 1 at 100 mA g 1 and stably run for 30 cycles [17]. Furthermore, doping heteroatoms such as B, N, S, P, etc. to the carbon architecture, can further increase the electronic conductivity and create abundant active sites [18]. On the other hand, developing the molecular architecture of red P with functional nanostructures under accurate control has also been verified to be effective to improve the electrochemical perfor mance. Particularly, assembling 3D spatial structures can highly over come the shortcomings of bulk materials by raising electrochemical reaction sites and buffering volume change effectively during cycling [19–21]. For instance, Zhou et al. designed a self-standing P/CNTS@rGO hierarchical structure by ball milling method as anode materials for SIB, delivering a high capacity of about 1000 mAh g 1 at 1C after 500 cycles [22]. Nevertheless, the active materials are easy to aggregate during the charging/discharging process by simply mixing the large-sized red P with hollow materials via the ball-milling method, where the P is coated outside of the hollow materials. As an alternative, a vaporization-condensation strategy was developed by infusing the P vapor into hollow structures under high temperatures and formed ho mogenous P@C hybrid materials. However, it should be mentioned that the P contents in most reported red P@C hybrid materials are less than 35 wt%, such as P-CNTs and P-rGO, which can’t meet the demand of energy density [23,24]. In this regard, further exploration of novel architectural carbon for achieving high P contents is of significance. Recently, a yolk-shell interlinked carbon, which contains cavities, mesoporous shells, and microporous cores anchored to the shell, was developed and provided large space for P accommodation and enabled the P@C composites high P contents. Moreover, the interconnect mes oporous shells built continuous electronic conduction networks to make sure fast electron transport through the P/C interface and improved the electrical conductivity of the P@C electrode. Equally important, the abundant porous structure is beneficial for alleviating the large volume change during cycling. Generally speaking, such a unique structure enables the P@C anode with high P content, high electronic conduc tivity and able to accommodate the volume change during cycling as well. Whereas, synthesis of such multi-cavity hollow structure is typi cally related to the template fabrication or multistep/high-cost pro cedures, which is highly related to result of the collapse of hollow structures. Thus, the rational design of multi-cavity hollow structures as P hosts with a low-cost and scalable procedure is urgent and necessary to the development of high-performance SIBs. In this work, we developed a facile, scalable and controllable
fabrication of N-doped C-encapsulated multi-cavity mesoporous frame (labeled as P@NMC) as illustrated in Scheme 1. Based on the afore mentioned merits, the SIBs assembled with as-prepared P@NMC anode perform out-standing charge storage capacities, rate performance as well as cycling stability. The kinetic properties of the P@NMC are investigated to understand the Na storage process in pseudo-capacitance properties. The capacitive contribution gradually increases as the scan rate increases, indicating P@NMC with a larger capacitive contribution is good for Naþ storage and renders a higher rate performance. 2. Experimental section The raw materials, preparation of nano emulsion, synthesis of multicavity polymer spheres (MCP) precursor are shown in the Supporting Information. Synthesis of N-doped multi-cavity nanocarbon (NMC). The surface of the NMC material was coated with a silica shell by a modified Stober process. 0.96 g of CTAB was stirred vigorously for 1 h by adding 30 mL of water. Then, the CTAB solution was added to a mixed alcohol-aqueous solution of 0.24 g of MCP, 150 mL of water, 60 mL of ethanol, and 2.4 mL of ammonia (25 wt%) for 30 min, and then 1.68 mL of tetraethyl orthosilicate was added drop by drop, subsequently stirred at a constant temperature in a 30 � C water bath for 16 h. MCP@mSiO2 was collected by centrifugation and washed several times with an HClalcohol solution to remove CTAB from the silica shell, then dried at 60 � C overnight. NMC@mSiO2 was prepared by heating the product at 800 � C for 2 h in an Ar gas atmosphere. Finally, NMC@mSiO2 was stirred at 50 � C for 24 h in aqueous NaOH to remove the mSiO2 shell to obtain NMC. 2.1. Synthesis of carbon-red phosphorus composites P@NMC The carbon-red phosphorus composites P@NMC was prepared based on a vaporization/condensation method. The commercial red P and the shower-like NMC were uniformly mixed at a mass ratio of 1:1 and placed in one sealed quartz tube under vacuum. The sealed quartz tube was then heated to 500 � C and kept at the temperature for 3 h, after cooling to 250 � C, it kept at the temperature for 24 h to convert white P to red P, and finally after cooling to room temperature, the product was washed several times with CS2 to remove white P and drying under vacuum. 3. Results and discussion The crystal structure of P confined in the NMC is analyzed by X-ray diffraction (XRD). As shown in Fig. 1a, the P@NMC exhibits a similar XRD pattern compared with the NMC and only two characteristic peaks are observed at 23.6� and 44.2� , corresponding to the carbon. No obvious peaks at 15.3� and 33.5� and 57.2� (PDF No. 44-0906), belonging to the red P, are detected, suggesting that the P is wellconfined in the porous structure of NMC and exists as amorphous structure [25,26]. The corresponding Raman spectra are shown in Fig. 1b. As can be seen, there are two Raman peaks of NMC appear at 1336.49 and 1159.65 cm 1, belonging to sp3 type disordered carbon and sp2 type graphitized carbon, respectively, which are also known as D band and the G band. The peak area ratio of D band to G band is adopted to evaluate the graphitization degree of carbon. NMC presents the value of 1.28, indicating the high graphitization degree of NMC. After P impregnation, the D band and G band exhibit no obvious shift and the value shows negligible change, proving that the structure of NMC maintains well, which coincides well with the XRD results. Moreover, no scattering peaks are detected at 300–500 cm 1, further confirming that the P is well-confined in the pores of NMC. Fourier transform infrared (FTIR) spectroscopy is hired to study the surface chemical states be tween the NMC and red P (Fig. 1c). The peak at 1510 cm 1 for the – C. The peak sample of NMC and P@NMC is assigned to the band of C– 2
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Scheme 1. Representative illustration of red P enveloped multi-cavity carbon spheres composite.
Fig. 1. XRD patterns (a), Raman spectra (b), FT-IR spectra (c) and TGA curves (d) of red P, NMC and P@NMC. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
appears at 1062 cm 1 for pristine red P corresponding to the charac teristic P–O due to the exposure of P nanoparticles in the air. For about P@NMC, the intensity of this position shows a little weak, suggesting the amorphous red P is confined in the NMC system. Furthermore, it can be noticed the formed peak at the position of 1001 cm 1 for the P@NMC sample should be attributed to P–O–C bond, suggesting the strong interaction between NMC and red P. There is no obvious characteristic peak of P in the XRD, Raman and FTIR spectra of P@NMC, revealing the
most of amorphous red P is encapsulated in the NMC frame. Meanwhile, static dissolution of red P, NMC and P@NMC is designed to simulate the working environment in batteries. We disperse the red P, NMC, and P@NMC into the liquid electrolyte and stand for 720 h. As shown in Fig. S1 (Supporting Information), the red P and NMC solutions still maintain their intrinsic colors, while the P@NMC solution appears stratification and the supernatant is colorless and transparent. The striking visual evidence illustrates the P@NMC composite can inhabit 3
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the free-diffusion of P from into the electrolyte and avoid unnecessary mass loss of red P during electrochemical cycling, which is beneficial for increasing the active materials utilization and improving the cycling stability. The P content in the P@NMC composite is determined by the thermal gravimetric analysis (TGA), as shown in Fig. 1d. As it can be seen a mass loss of 50.1 wt% is observed between 400 � C and 700 � C. It can be attributed to the P sublimation, suggesting a high P content of 50.1 wt%, which is much higher than most of current reported P@C composite [14, 23,24,27,28]. X-ray photoelectron spectroscopy (XPS) is taken to further analyze the interaction between red P and NMC. The full-spectra scan of NMC and P@NMC are shown in Fig. S2 (a). The deconvolution of the N1s spectrum identifies the surface functionalities of NMC and P@NMC
(Fig. S2 (b)). There are three fitted component peaks at 399.6, 401.1 and 402.2 eV for NMC and P@NMC, which are designated for pyridinic (N-6), pyrrolic/pyridine (N-5) and quaternary nitrogen (N-Q), respec tively. This order is in accordance with the previous reports [29]. The proper N-doping is essential for improving the electronic conductivity of the NMC and make further efforts to optimize the electrochemical per formance of the P@NMC. For the red P, two peaks located at 129.9 and 130.8 eV are observed in the P–P 2p spectrum due to the 2p3/2 and 2p1/2 spin-orbit splitting of P–P bonds [30]. In addition, red P presents a characteristic peak at 134.6 eV, belonging to the P–O bond, which can be explained as the partial oxidization by the air (Fig. S2 (c)). In the case of P@NMC, the signals of P–P bond peaks disappeared, which can be attributed to the encapsulation of P by the NMC and limited penetration
Fig. 2. SEM (a, b), TEM (c-f) images of the prepared NMC and P@NMC; EDX mapping images (g, the element of C, N, P) of the fabricated P@NMC samples. 4
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depth of X-ray. In particular, there is a significant intensity increasing for the peak located at 134.6 eV, demonstrating the formation of P–O–C bond during the vaporization-adsorption process due to the chemical interactions between red P and NMC [31]. Morphology of the NMC is examined by SEM (Fig. 2a). The carbon spheres exhibit smooth outer surfaces with an approximate diameter of 100–200 nm. After P impregnation, the morphology of NMC shows negligible change and the as-prepared P@NMC composite present almost the same morphology compared with NMC (Fig. 2b). The inner hollow space and porous structure of NMC and P@NMC are further demonstrated by the TEM images (Fig. 2c–f). As it can be seen, the surface of the carbon wall is smooth and uniform in thickness about 15 nm for the NMC (Fig. 2c). The interior of the carbon sphere is crosslinked by ultra-thin carbon walls and uniformly divided into a plural ity of small cavities with a size of 7 nm. After loading red P, a large shade area is observed on the TEM image of P@NMC (Fig. 2d). Meanwhile, the grid structure of NMC becomes inconspicuous and the crystal lattice is not detected at all, according to Fig. 2f. The element mappings of C, N, and P clearly demonstrate the homogeneous distribution of red P in NMC (Fig. 2g). Therefore, the amorphous red P is confined in the multicavities of carbon effectively and forms multi-cavity yolk-shell struc tured P@NMC composite. To further study the existing state of red P in NMC, the specific surface area and pore size distribution of red P, NMC and P@NMC composites are tested by N2 adsorption-desorption isotherm test. As shown in Fig. S3, NMC presents a mixture of type-I and type-IV curves, indicating abundant micro-pores (less than 2 nm) and meso-pores (around 7 nm), confirmed by Fig. S3b. Owing to the abundant microand meso-pores, the NMC delivers a large surface area of 730.7 m2 g 1 and an average pore size of 4.2 nm. The presence of these pores can facilitate electrolyte infiltration and rapid transport of Naþ. After P loaded, the specific surface area of the obtained P@NMC composite decreases to 19.4 m2 g 1 and almost all the micro- and meso-pores in P@NMC disappear, further proving the successful P infusion. The electrochemical performance of the P@NMC composite, as an anode, is firstly evaluated by the cyclic voltammetry (CV) testing at a scan rate of 0.1 mV s 1 from 0.01 to 3 V. As shown in Fig. 3a, during the first negative scanning, two peaks are observed at 0.81 V and 0.44 V, corresponding to the electrolyte decomposition to form the SEI, reduc tion of red phosphide to NaxP, respectively. When the voltage is further reduced to near 0.01 V, the current increases sharply, which can correspond to the further reaction of carbon with sodium. Following the positive scanning, the peaks at about 0.39 V can be detected, corre sponding to the stepwise de-sodiation reaction. For the second cycle, the reduction peak at 0.77 V and oxidization peak at 0.39 V are maintained, suggesting the reversible electrochemical reaction between P and Na3P [10–14,23,24]. Fig. 3b shows the Charge/Discharge profiles of the SIBs assembled with P@NMC anode at the 1st, 2nd, 5th and 50th cycles operating at a current density of 0.5 A g 1. As can be seen, the cell de livers a high initial discharge capacity of 2146.4 mAh g 1. Similar to the CV results, there is a large irreversible capacity between the 1st cycle and the 2nd cycle. The capacity sharply drops to 1210.4 mAh g 1 at the 2nd cycle and result in a low initial CE of 56.4%. The low CE results from the decomposition of the electrolyte forming an irreversible SEI film and the generation of Na2C60, which caused rapid capacity decay. Except for the first cycle, the SIBs assembled with P@NMC anode exhibit excellent cycling stability, proving by the overlapped charge/discharge profiles in Fig. 3b. Additionally, high CEs of over 98.7% are achieved. The discharge profile includes the three sloped regions of 1.5 to 0.75 V (vs. Naþ/Na), 0.75 to 0.39 V (vs. Naþ/Na) and 0.39 to 0.01 V (vs. Naþ/Na). The charging curve also exhibits the inclined platform at 0.39, and 0.75 V, respectively. The process corresponds to sodium ion intercalation and de-intercalation process, which is in agreement with the CV analysis. To clarify the Na-ion insertion mechanism of P@NMC electrode, exsitu XRD analysis is performed to investigate the structural evolution at different (dis)charge states (Fig. 4a) with a scan rate of 0.5� min 1.
Fig. 3. Electrochemical performance of the prepared P@NMC for SIBs. (a) CV curves with a scan rate of 0.1 mV s 1. (b) Charge/Discharge curves for the 1st, 2nd, 5th and 50th cycles at current density of 0.5 A g 1.
When P@NMC electrode is continuously discharged from the initial state of 2.0 V–0.75 V, the characteristic peaks of Na2C60 (ICSD: 47-1577) appear in the XRD patterns, and their intensity increases gradually with the decrease of potential. Some peaks of Na2C60 also exist in the charging process without significant weakening such as (111) and (220) plane, while other peaks’ intensity of lattice planes (222) and (440) shows a little declining. This result means the Naþ reacts with NMC to form Na2C60, and some Naþ cannot emerge from Na2C60 during the electrochemical cycle. However, these inert Na2C60 contents may act as buffer medium and contribute greatly to the cycle stability of P@NMC. In fact, according to our test, the pure NMC shows stable sodium storage capacity of about 140 mAh g 1. When discharged to 0.75 V, the peaks belong to Na3P (102), Na3P (110) and Na3P (103) of (ICSD:04-0764) are detected, indicating that the embedded Na-ions are alloyed with P and form Na3P initially. As the depth of discharge increases, the intension of Na3P peaks gradually sharped and arrives at maximum value when the potential reduces to 0.01 V. Subsequently, its intension gradually re cedes with increasing depth of charge, and the peaks of (110) and (103) disappear while the potential reaches 0.75 V. Another interesting phe nomenon also should be attention, the NaP (ICSD: 14009) appears during the discharge process since the potential declines at about 0.4 V 5
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capacities of no more than 50 mAh g 1, while NMC and P@NMC elec trodes deliver a specific discharge capacity of 145.3 and 1076.9 mAh g 1 at 0.5 A g-1, respectively. This result is consistent with the discussion of electrochemical reaction mechanism above. Despite the initial irre versible capacity loss, both electrodes exhibit a steady reversible ca pacity in subsequent cycles. In addition, the rate performance of P@NMC electrode is evaluated under different current densities, as shown in Fig. 5b. P@NMC electrode renders superior rate performance. It delivers an average reversible capacity of 1076.5, 1063.6, 898.3 and 786.2 mAh g 1 at 0.5, 1, 2 and 3 A g 1, respectively. Even under a high current rate of 5 A g 1, an average reversible capacity of 646.3 mAh g 1 is achieved. When the current density is recovered to 0.5 A g 1, the specific capacity returns to 1042.5 mAh g 1. Cyclic stability testing is an important indicator for the electrode. To evaluate the long-term cyclic behavior of P@NMC, the cyclic performance is examined at 0.5 A g 1 for 1000 cycles (Fig. 5c). The electrode delivers a reversible capacity of 923.7 mAh g 1 after 1000 cycles and achieves high CEs over 98%, demonstrating its excellent cyclic stability. The P@NMC exhibits remarkable improvement compared with most recently reported Pbased materials (Table S1) [32–38]. The structure of P@NMC after cycling is detected by TEM (Fig. 5d). Promisingly, after the long-term cycling, the P@NMC maintains its structure well and no breakage is observed, which suggests the strong capability of NMC in accommoda ting the volume change during cycling and clarifies the reason for the excellent cycling stability of P@NMC anode. To further explain the excellent sodium storage performance of P@NMC electrode, the kinetic features of P@NMC electrode are analyzed by the CV test for the survey on pseudo-capacitance properties during the charge-discharge process. As the scan rate from 0.2 increase to 2 mV s 1 (Fig. 6a), the rela tionship between the measured peak currents related to Naþ storage in P@NMC and sweep rates can be presented by the power law [39–41]. i ¼ avb
(1)
Generally, the i and v are the peak current and scan rate, respectively. The value of b confirms the type of Naþ insertion/extraction. The b-value of 0.5 indicates a diffusion-controlled behavior (battery), whereas the bvalue of 1.0 represents a capacitive-controlled behavior (capacitor) [40]. According to the linear relationship between log i and log v plots (Fig. 6b), the slopes for cathodic peaks 1, cathodic peaks 2 and anodic peaks are 0.688, 0.540 and 0.541, respectively, revealing a pseudoca pacitive process between typical behaviors of battery and capacitors. Especially, cathodic peak 1 exhibits higher pseudo capacitive behaviors, which is attributed to the contribution of the porous carbon skeleton to capacitance in P@NMC materials [42]. The capacity contribution from pseudocapacitive and diffusion-controlled charge can be quantified by using the given relationship: i(V) ¼ k1 υ þ k2 υ1/2(2).
Fig. 4. Structure evolution of the electrochemical process: (a) The Ex situ XRD of the discharge-charge profile of P@NMP at different potentials for the 1st cycle at current density of 0.5 A g 1. (b) The contour and response surface analysis corresponding to discharge-charge process. (c) Representative illus tration of the electrochemical reaction process for the P@NMC.
k2 υ1/2 represents diffusion-controlled reactions and k1 υ represents capacitance control [43,44]. Therefore, the relative contributions from the capacitive process and Naþ intercalation can be decoupled (Fig. 6c). At the scan rate of 0.2 mV s 1, the capacitive charge contribution is calculated to be 22.6%, which increases with scan rate and found to be 54.5% at a scan rate of 2 mV s 1. This behavior is not surprising because the pseudo-capacitive contribution should play a critical role in the case of carbon composites, which possesses stable mesopore structures. It should be noted that increasing tendency refers to the contribution ratio rather than the absolute value because pseudo-capacitance is an intrinsic characteristic of a material. The typical voltage profile for the capacitive current (green region) in comparison with the total current is shown in the Fig. 6d, at scan rate of 0.2 mV s 1, a dominating capacitive contribution (~22.6%) is obtained for the P@NMC anode, implying it has a good capacitive contribution, and thus results of the improvement of high capacity and rate performance [44]. The electrochemical impedance spectroscopy (EIS) for three
and it exhibits the same trend as that of Na3P. The highest intensity of NaP peak presents at 0.01 V during the discharge process, while its in tensity becomes weak gradually and then fades away when the potential rise at 0.75 V during the charging process. Figs. S4–S6 exhibit the XRD patterns of Na2C60, Na3P, and NaP when the P@NMC anode discharges to 0.01 V, respectively. Moreover, the contour and response surface analysis corresponding to the discharge-charge process is shown in Fig. 4b, which reflects the (de)alloying process of Na3P and NaP during cycling visually. Therefore, high activity reversible Na3P and NaP ren ders the electrode with a high specific capacity, and the electrochemical reaction process is listed in Fig. 4c. The cyclic performance of commercial red P, NMC and P@NMC electrodes is presented in Fig. 5a. The pure P delivers low specific 6
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Fig. 5. (a) Cycling performance of commercial red P, NMC and P@NMC at a constant current density of 0.5 A g 1. (b) Rate capability of P@NMC from 0.5 A g 1 to 5 A g 1. (c) Cycling capability of P@NMC with prolonged cycle life (1000 cycles) at 0.5 A g 1. (d) TEM image of P@NMC electrode after 1000 cycles. (For inter pretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
Fig. 6. Kinetics analysis of P@NMC electrode. (a) CV curves of P@NMC at different scan rates (0.2–2 mV s 1). (b) Relationship between logarithm redox peak current and logarithm scan rate of the P@NMC. (c) Normalized contribution ratio of capacitive at different scan rates. (d) CV profile of the capacitive contribution at 0.2 mV s 1.
different electrode materials: red P, NMC, and P@NMC are shown in Fig. S7. Table S2 shows the impedance parameters of each impedance component obtained by Z-Simpwin simulation of the equivalent circuit, including the ohmic resistance (Rs), charge transfer resistance (Rct) and Warburg impedance (W). The ohmic resistance (Rs) of P@NMC is simulated as 19.17Ω, which is slightly lower than the pristine red P (20.87Ω). The Rct value of the P@NMC composite electrode is 183.1 Ω,
which is much lower than the charge transfer resistance of red P (616.1 Ω), indicating the enhanced charge transfer dynamics. The improved electrochemical kinetics can be attributed to the interconnected elec tronic conductive network in NMC that enhanced the electronic con ductivity and confined the P as nanoparticles. Overall, such a unique structure of NMC can facilitate electron transport and alleviate the volume change, enabling the high P loading NIBs with faster 7
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electrochemical performance.
Journal of Power Sources 450 (2020) 227696
kinetics,
excellent
cycling
stability,
and
rate
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4. Conclusions In summary, a P@NMC composite with high red P content (over 50 wt%) is successfully developed. The multi-cavity NMC acts as an effec tive P host by structural monomers driven by surface free energy. The mesh-like multi-cavity interior provides a large space for P accommo dation as well as alleviate volume change during cycling. Moreover, the interconnected porous carbon walls build an interconnected electronic conductive network for fast electron transport. Nitrogen doping in car bon spheres enhances the conductivity and also provides more active sites for electrochemical reactions. Benefitting from the aforementioned merits of the unique structure, the NIBs assembled with P@NMC anode present excellent electrochemical performance, which maintains a high capacity of 923.7 mAh g 1 after 1000 cycles at a current density of 0.5 A g 1. The electrochemical performance of P@NMC is superior to most of the current phosphorus-carbon anodes. This general strategy can be extended to further explore other advanced energy materials. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements Y.H.Z. and B.H.L. contributed equally to this work. The authors would like to acknowledge financial support provided by Key National Natural Science Foundation of Yunnan Province (No.2018FA028 and No.2019FY003023), National Natural Science Foundation of China (No.51474191 and No.21467030), Major State Basic Research Devel opment Program of China (973 Program, No. 2014CB643406), and the Program for Outstand Young Talents (2018) of Yunnan University. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.jpowsour.2019.227696. References [1] W. Chen, T. Li, Q. Hu, C. Li, H. Guo, J. Power Sources 286 (2015) 159–165. [2] W. Wang, Y. Guo, L. Liu, S. Wang, X. Yang, H. Guo, J. Power Sources 245 (2014) 624–629. [3] H. Liu, B. Liu, H. Guo, M. Lianga, Y. Zhang, T. Borjigina, X. Yang, L. Wang, X. Sun, Nano Energy 51 (2018) 639–648. [4] X. Song, X. Li, Z. Bai, B. Yan, D. Li, X. Sun, Nano Energy 26 (2016) 533–540. [5] Z. Liu, X. Yu, X.W. Lou, U. Paik, Energy Environ. Sci. 9 (2016) 2314–2318. [6] S. Wang, L. Xia, L. Yu, L. Zhang, H. Wang, X.W. Lou, Adv. Energy Mater. 6 (2016) 1502217. [7] H. Liu, H. Guo, B. Liu, M. Liang, Z. Lv, K.R. Adair, X. Sun, Adv. Funct. Mater. 28 (2018) 1707480.
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