Rational Design of Three-Dimensional Graphene Encapsulated with

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Rational Design of Three Dimensional Graphene Encapsulated with Hollow FeP@carbon Nanocomposite as Outstanding Anode Material for Lithium Ion and Sodium Ion Batteries Xiujuan Wang, Kai Chen, Gang Wang, Xiaojie Liu, and Hui Wang ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.7b06625 • Publication Date (Web): 19 Oct 2017 Downloaded from http://pubs.acs.org on October 20, 2017

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GO Carbon-coating

H-Fe3O4

Phosphatizing

Stirring

NaH2PO2

Calcination

H-Fe3O4@C

H-FeP@C

H-FeP@C@GR

Three dimensional graphene encapsulated with hollow FeP@carbon was successfully fabricated via a combination of hydrothermal reaction, carbon coating process, phosphidation treatment, and carbothermic reaction. They exhibit excellent cycle stability and superior rate capability when used for LIBs and SIBs.

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Rational Design of Three Dimensional Graphene Encapsulated with Hollow FeP@carbon Nanocomposite as Outstanding Anode Material for Lithium Ion and Sodium Ion Batteries Xiujuan Wanga, Kai Chenb, Gang Wangb, Xiaojie Liua*, Hui Wanga* a

Key Laboratory of Synthetic and Natural Functional Molecule Chemistry (Ministry

of Education), College of Chemistry & Materials Science, Northwest University, Xi'an 710069, PR China b

National Key Laboratory of Photoelectric Technology and Functional Materials

(Culture Base), National Photoelectric Technology and Functional Materials & Application

International

Cooperation

Base,

Institute

of

Photonics

Photon-Technology, Northwest University, Xi’an 710069, PR China

*

Corresponding author:

Tel.: +86 029 88363115 Fax: +86 29 88302571 E-mail address: [email protected] (H. Wang). [email protected]


&

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Abstract Transition metal phosphides (TMPs) have been extensively investigated owing to their high theoretical capacities and relatively low intercalation potentials vs. Li/Li+, but their practical applications have been hindered by low electrical conductivity and dramatic volume variation during cycling. In this work, an interesting strategy for rational design of graphene (GR) encapsulated with hollow FeP@carbon nanocomposite (H-FeP@C@GR) via a combination of hydrothermal route, carbon coating process, phosphidation treatment, and carbothermic reaction is reported. The hollow FeP (H-FeP) nanospheres shelled with thin carbon layers are wonderfully incorporated into the GR matrix, interconnecting to form a three dimensional (3D) hierarchical architecture. Such a design offers distinct advantages for FeP-based anode materials for both lithium-ion batteries (LIBs) and sodium-ion batteries (SIBs). For example, the 3D omnibearing conductive networks from the GR skeleton and outer coating carbon can provide opened freeway for electron/ion transport, promoting the electrode reaction kinetics. In addition, the wrapping of H-FeP nanosphere in a thin carbon layer enables the formation of solid electrolyte interphase (SEI) on the carbon layer surface instead of on the individual H-FeP surface, preventing the continual reforming of the SEI. When used as anode materials for LIBs and SIBs, H-FeP@C@GR exhibited excellent electrochemistry performances. Keywords: FeP, carbon coating, graphene, batteries, DFT calculations


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Rechargeable lithium ion batteries (LIBs) have been widely used in portable electronics, electric vehicles and stationary energy storage because of their advantages of high energy, long lifespan and environment benignity over other alternatives.1-5 However, considering the limited and unevenly distributed availability of Li deposits on the earth, LIBs have not been able to meet tremendous demand for large scale storage systems, leading to the reality that its price will become an important issue for future LIBs development. To address these concerns, rechargeable sodium ion batteries (SIBs), as alternative to LIBs, have received increasing attention in recent years, owning to the wide availability of sodium resources ubiquitous around the world, potential low cost, along with similar chemistry to LIBs. Nevertheless, compared to LIBs, the more sluggish kinetics occurs in insertion/extraction of SIBs due to the larger sodium ionic radius, resulting in poor electrochemical performance, which becomes the specific hindrance for future SIBs advancement.6,7 Therefore, it is imperative for scientists to advance the battery storage technology in both LIBs and SIBs so as to support certain applications. One practical approach in this direction is to design and synthesize appropriate electrode materials that possess the capability of playing a dual storage role for both lithium and sodium. Among various negative electrode materials, transition metal phosphides (TMPs) have been extensively investigated owing to their high theoretical capacities and relatively low intercalation potentials vs. Li/Li+. So far, many transition metal phosphides, such as Sn4P3,8 Cu3P,9 Ni2P,10 MoP,11 CoP,12 VP,13 and ZnP214 have displayed good electrochemical performances in lithium and/or sodium storage. Recently, iron phosphide (FeP), as one kind in transition metal phosphide family, owning theoretical capacity of ～926 mAh g 1, has been proven to be one of the most −

promising candidates as anode materials for LIBs and SIBs enlisting a reversible


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two-step insertion/conversion process.15-17 In the detailed reaction route, FeP with orthorhombic structure react fully with lithium or sodium through a conversion reaction: FeP + 3Li/Na → intermediate tetragonal LiFeP/NaFeP phase → Li3P/Na3P + Fe0 in its first discharge cycle, resulting in nanocomposite expressed by nanosized Fe0 particles embedded in Li3P or Na3P matrix (illustrated in Figure 1a). Moreover, within FeP unit cell, Fe-P distances are in between 2.186 and 2.447 Å,18 larger than lithium ion diameter 1.08 Å and sodium ion diameter 2.04 Å, which favors the immigration of Li+ and Na+ ions in the FeP crystal structure, ameliorating the ionic conductivity and reducing the volume change to some degree during ion insertion and desertion. Despite FeP has several obvious advantages over conventional oxides and sulfides materials, the most intractable obstacles are still hindering the future development of anode materials, including: (1) inevitable volume variation during charge and discharge, which not only leads to serious electrode material pulverization and loss of the electrical contact between active materials and the current collector but also makes it difficult to coat a stable protection layer to maintain the solid electrolyte interphase (SEI), and eventually results in serious capacity fading during long-term cycling.19-23 (2) low electronic conductivity, hampering the fast transfer of electrons inside the active materials and confining their electrochemical reactions.24,25 To cope with these problems, the most commonly used approach is to design a nanoscale FeP coated with carbon layer. The protective carbon shell acts as a physical buffering layer for the large volume change (cushion effect), improves the electronic conductivity through intimate contact between carbon and FeP and prevents FeP pulverization during Li+/Na+ insertion-extraction process.26-30 For instance, recently, Han et al. reported that a nanoplate FeP shelled with carbon layer had a high reversible lithium storage capacity up to 720 mAh g 1 and a capacity retention rate of 96% after 100 cycles at a current −


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density of 200 mA g 1.15 He and his co-workers also synthesized an amorphous and −

mesoporous FeP modified with a dual-carbon phase (carbon coating and CNTs) used as an anode material for SIBs, which demonstrated a utilization rate of 78% for the active material and a reversible capacity of 415 mAh g 1.31 Even though rate capability and −

cycle life of the nanostructures have been markedly enhanced by coating amorphous carbon on the surface of FeP as an outer conductive layer, the electrochemical performance improved by single carbon coating are still far from meeting the requirements for large scale energy storage. For example, a single introduced carbon matrix into nanostructured FeP cannot maintain the structural stability for a very long cycling, and the individual outer carbon coating cannot fully ensure a fast ion/electron transfer across the interfaces. Hence, an upgrade strategy combining a design of hollow nanostructure FeP and a double carbon network modification consisting of amorphous carbon layer and 2D Graphene (GR) is proposed by ourselves to solve the aforementioned challenges. Compared to the previously reported work, the hollow void space of FeP can accommodate the huge volume change during cycling, delaying capacity fading. More importantly, GR, a two-dimensional (2D) carbon nanostructure, has been recently considered as a promising material for energy storage due to its superior electrical conductivity, excellent mechanical flexibility, large specific surface area, and good thermal and structural flexibility.32-34 Herein, we attempt to design and fabricate a GR encapsulated with hollow FeP@carbon nanocomposite (H-FeP@C@GR) via a combination of hydrothermal reaction, carbon coating process, phosphidation treatment, and carbothermic reaction (illustrated in Figure 1b). The hollow FeP (H-FeP) nanospheres shelled with thin carbon layers are well incorporated into the GR matrix, interconnecting to form a three dimensional (3D) hierarchical architecture. The crucial roles of GR on upgrading a


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simple FeP@carbon nanocomposite to a complex 3D network structure nanocomposite are as follows: (1) The 3D omnibearing conductive networks from the GR skeleton and outer carbon coating can provide rapid diffusion channels for electron/ion transport, promoting the electrode reaction kinetics. (2) The GR, as a cage-like framework around the H-FeP@C nanospheres, is of importance to prevent the agglomeration of active materials and cracking of the electrode. (3) The large volume change of H-FeP@C during Li+/Na+ insertion-extraction process can be buffered through accommodation by the elastic GR along with available internal voids, ensuring the structural integrity of electrode. As a result, when used as anode materials for LIBs and SIBs, H-FeP@C@GR nanocomposite exhibited high reversible capacities, ultra-long cycle life and superior rate performance. Results and Discussion As illustrated in Figure 1b, the detailed preparation procedure of 3D H-FeP@C@GR nanocomposite

comprises

hydrothermal

reaction,

carbon

coating

process,

phosphidation treatment and annealing process. Thanks to the hydrophobic nature of graphene, the anode materials with addition of graphene are favorable for the contact with organic LiPF6 electrolyte. Surface wettability of LiPF6 droplet on the H-FeP, H-FeP@C@GR, and GR were shown in Figure S1. The H-FeP membrane exhibited a large contact angle of 24.7° at the LiPF6/H-FeP interface and this angle decreased slightly after 8s, indicating its poor surface wettability. Whereas, in the case of H-FeP@C@GR and GR membranes, wetting behaviors were significantly changed due to the existence of GR. LiPF6 droplet quickly spread out on the surface within 8s, demonstrating their high hydrophobicities, which is beneficial for FeP contact with nonpolar LiPF6 electrolyte. Moreover, high electrical conductivities were also demonstrated in the H-FeP@C@GR nanocomposite. As illustrated in Figure S2, due


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to the 3D omnibearing conductive networks from the GR skeleton and outer coating carbon, a higher electrical conductivity (1.4 × 102 S m−1) was obtained for the H-FeP@C@GR, compared to those of pure H-FeP (4.8 × 10-3 S m−1) and H-FeP@C (0.86 S m−1), which can impart greater reversibility to reaction. SEM images in Figure S3a and b demonstrated that the obtained H-Fe3O4 nanospheres as starting material were highly uniform with an average diameter of 300-400 nm. They had very high homogeneity and insignificant particle aggregation. The purity of H-Fe3O4 nanospheres was corroborated by XRD (JCPDS: 19-0629) in Figure S4. The surfaces of these nanospheres were uneven, clearly indicating that the nanospheres were assembled by ultra-nanoparticles driven by Ostwald ripening. TEM image in Figure S3c confirmed that the nanospheres consisted of ultra-nanocrystals and displayed that they had hollow nanostructure. According to Scherrer equation based on XRD, the crystallite size was 4.1 nm. In Figure 2a-c, it was seen that H-FeP, obtained by the phosphidation of H-Fe3O4, remained spherical shape with high uniformity and without obvious aggregation. H-FeP nanospheres, the same as H-Fe3O4, were assembled by small nanocrystals, suggesting that phosphidation took place on a single H-Fe3O4 ultra-nanocrystal. The purity of H-FeP nanospheres was confirmed by XRD (JCPDS: 65-2595) in Figure 4a. Based on Scherrer equation, the H-FeP crystallite had the size of 9.6 nm, higher than that of H-Fe3O4, which might be due to the slight sintering in phosphidation process. Meanwhile, the hollow nanostructure and self-assembly of H-FeP nanosphere were verified by TEM images in Figure 3a and b. HRTEM image of H-FeP was shown in Figure 3c, which suggested that these small nanoparticles had sizes of tens of nanometers. The lattice resolved TEM image taken over the area marked by a square in Figure 3c displayed a highly ordered fringe with a d-spacing 0.251 nm, which was attributed to the (102) reflection of FeP phase.


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Subsequently, as-prepared H-Fe3O4 product was carbonized by using glucose as a carbon source to prepare H-Fe3O4@C. The success of this strategy could be confirmed by SEM and TEM images and XRD analysis (Figure S3d-f and S4), where hollow Fe3O4 coated with carbon maintained well the hollow structure, spherical shape as well as unchanged size. These resultant H-Fe3O4@C nanospheres were then converted into H-FeP@C nanospheres via a solid/gas-phase phosphorization strategy. In this process, NaH2PO2 was used as a precursor to generate PH3 gas, which could react with the H-Fe3O4@C nanospheres to form H-FeP@C nanospheres.15 SEM and TEM images (Figure 2d-f and Figure 3d) showed that the spherical morphology was preserved after phosphidation step. By observing HRTEM image (Figure 3f) magnified from the edge of the nanosphere in Figure 3e, it was found that an obvious amorphous carbon layer was formed on the surface of H-FeP nanosphere. The thickness of the carbon layer was around 4 nm. XRD of H-FeP@C in Figure 4a showed the high purity of single FeP phase and the crystallite size was 9.8 nm based on Scherrer equation. Finally, H-FeP@C@GR nanocomposite was obtained by mixing H-FeP@C with GR. The purity of H-FeP nanospheres in H-FeP@C@GR nanocomposite was confirmed by XRD in Figure 4a. In Figure 2g and h, SEM and mapping images indicated that H-FeP nanospheres coated with thin carbon layers were incorporated into the GR matrix. In this progress, GO was successfully transformed to GR, the succession of this strategy could be further confirmed by SEM images and XRD analysis (Figure S5 and 6), where layered structures, ruffled, and clear layer spacing appearance of graphene-based sheets can be clearly observed. EDX spectrum in Figure 2i revealed that the atomic ratio of Fe:P was 1:1, matching with XRD result. Figure 3g and h exhibited that H-FeP nanospheres remained spherical shape and that H-FeP@C@GR nanocomposite maintained integrity of 3D nanostructure. The corresponding selected-area electron


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diffraction (SAED) pattern clearly showed four bright diffraction rings, which could be indexed to (111), (202), (211), and (212) planes of orthorhombic FeP,35,36 revealing the formation of polycrystalline FeP (Figure 3i). In order to further evaluate the actual carbon content and its quality, Thermal analysis and Raman spectroscopy were both performed. In Figure 4b, the small weight loss of about 2 wt% below 200 °C was ascribed to the removal of adsorbed water on the surface of H-FeP@C@GR nanocomposite, while the weight increase between 200 and 380 °C was attributed to the gradual oxidation of FeP to Fe2O3 and P2O5.37 Finally, the drastic weight loss from 400 to 600 °C indicated that the carbon component in H-FeP@C@GR nanocomposite was completely burned. By calculating the TGA results of H-FeP, H-FeP@C, and H-FeP@C@GR based on oxidation reaction, the weight fractions were 67.6 wt%, 10.7 wt% and 21.7 wt% for active material FeP, coating carbon layer and GR, respectively, in H-FeP@C@GR nanocomposite. In the Raman spectra of Figure 4c, two broad peaks located at 1355 and 1593 cm 1 could be −

observed for H-FeP@C@GR nanocomposite and H-FeP@C, corresponding to the sp3-type disordered carbon form (D band) and sp2-type ordered graphitic carbon form (G band), respectively.38-40 The ID/IG ratio of H-FeP@C@GR nanocomposite was calculated to be 1.07, slightly higher than that of H-FeP@C 0.72, suggesting that the H-FeP@C@GR nanocomposite possessed more defects and disorders in the carbon component, which was favorable to improve the diffusion rate of Li+/Na+ ions and electrons.41,42 After making a complete literature survey, we find that the ID/IG ratio of the H-FeP@C@GR nanocomposite in our work is higher than those of the majority of double carbon network modification anode materials reported in the literatures,34,43-46 which indicates that carbon layer and graphene network with a certain number of defects and disorders in the H-FeP@C@GR could enhance the electrical and ionic


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transport significantly. XPS was employed to probe the surface electronic states and chemical composition of H-FeP@C@GR nanocomposite. The survey scan XPS spectrum (Figure 5a) revealed the presence of Fe, P, C, and O elements on the surface of the sample, which was consistent with the above mapping results. The high-resolution Fe 2p XPS spectrum was shown in Figure 5b. The peaks centered at 707.7 and 720.9 eV could be attributed to Fe (ш) 2p3/2 and 2p1/2 of FeP. While the peaks with binding energy values of 713.3 and 727.2 eV could be assigned to the oxidized form of Fe deriving from the prolonged exposure of H-FeP@C@GR to air.25,47 For the high-resolution P 2p XPS spectrum (Figure 5c), two peaks at 129.3 and 130.1 eV corresponded to the binding energy of P 2p3/2 and P 2p1/2, respectively. The other P 2p peak located at 133.6 eV could be assigned to the partial surface oxidation of P species during the process of XPS measurement when the sample was exposed to ambient air.35,48 The high-resolution XPS spectrum of C 1s (Figure 5d) was fitted with four components. In addition to main peak located at 284.4 eV attributed to sp2-hybridized graphitic C atoms and the peak at 285.6 eV ascribed to sp3-hybridized C atoms, the other two peaks at 286.4 and 288.8 eV were attributed to the functional groups on GR sheets, such as C=O, and O-C=O.49,50 Electrochemical characterizations for LIBs The electrochemical performance of H-FeP@C@GR nanocomposite was firstly investigated by CV at the scan rate of 0.1 mV s-1 in the 0-3.0 V voltage range (Figure 6a). An obvious difference between the first and the subsequent cycles was clearly seen. In first cathodic sweep, an apparent wide reduction peak appeared at approximately 1.02 V, which could be attributed to the intercalation of Li+ ions into FeP to form LixFeP (x=0-3) phase (eqn (1)).18 The following peak at 0.44 V could be explained by


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the formation of SEI film on the electrode (eqn (2)) and the reduction of LixFeP into Fe and Li3P phases (eqn (3)).51 The following anodic scan gave a broad oxidation peak at 1.08 V, which related to the reverse processes of eqn (1) and (3). In the subsequent cycles, it is noteworthy that the reduction peak shifted to 0.78 V while the oxidation peak stayed at the same position as in the first cycle, indicating an excellent reversible nature of the electrochemical reactions in H-FeP@C@GR nanocomposite anode. For comparison, CV result of H-FeP electrode was shown in Figure S7a, which apparently displayed that the current density decreased gradually from 1st to 3 rd cycle, indicating a relatively poor reversibility of H-FeP electrode. FeP + xLi+ + xe → LixFeP −

(x=0-3)

(1)

Li+ + e + electrolyte → SEI (Li)

(2)

LixFeP + (3-x)Li+ + (3-x)e → Fe + Li3P

(3)

−

−

The first three discharge-charge curves of H-FeP@C@GR nanocomposite anode at a current density of 0.2 A g 1 in a voltage window of 0.005-3 V were also shown in Figure −

6b. The specific capacities were calculated based on the total mass of H-FeP@C and GR. In the first cycle, H-FeP@C@GR electrode yielded an initial discharge and charge capacity of 1566 and 1154 mAh·g–1, respectively, with a Coulombic efficiency (CE) of 74%. These values were higher than those of H-FeP nanospheres (1269 mAh g 1 at 0.2 −

A g 1 with an initial CE of 66%) (Figure S7b). The initial irreversible capacity losses for −

the two electrodes might be attributed to the formation of SEI layer on electrode surface and electrolyte decomposition, which occurs commonly in all of anode materials.52-54 This characteristic was exactly consistent with the CV results that the reduction peaks, attributed to SEI formation, were present in the first cathodic sweep while absent afterward. In the subsequent two cycles, the specific capacities of H-FeP@C@GR


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exhibited only a slight decay, while the capacities of H-FeP decreased to a large extent. The excellent performance and high stability of H-FeP@C@GR electrode clearly benefited from the introduction of coating carbon and encapsulated GR. Unexpectedly, no obvious plateaus at approximately 0.78 V for discharge and 1.08 V for charge, matching with CV results, were found in the discharge-charge curves, which might be due to that the cathodic reaction of Li+ ions into FeP to form LixFeP (x=0-3) phase at 0.78 V and anodic reaction of delithiation of LixFeP (x=0-3) to reform FeP at 1.08 V were insufficiently active to build the noticeable plateaus in those curves.9,41 The rate capabilities of all as-prepared electrodes from 0.2 to 8 A·g–1 were evaluated shown in Figure 6c. The H-FeP electrode delivered discharge capacities of 529 (10th), 221 (20th), 88 (30th), 42 (40th), 21 (50th), and 10 (60th) mAh g 1 at current densities of 0.2, 0.5, 1.0, −

2.0, 4.0, and 8.0 A g 1, respectively. By comparison, at the same rate, the corresponding −

discharge capacities for H-FeP@C electrode were 859, 570, 441, 318, 222, and 187 mAh g 1. Notably, the discharge capacities of two electrodes were unsatisfactory. In the −

case of H-FeP@C@GR electrode, it offered obviously enhanced specific capacities of roughly 1030, 876, 755, 657, 577, and 482 mAh g 1, respectively, demonstrating the −

architecture advantage of 3D H-FeP@C@GR nanocomposite for high capacity and high rate lithium storage. Figure 6d manifested the cycling performance of H-FeP, GR, H-FeP@C nanospheres, and H-FeP@C@GR nanocomposite at a current density of 0.2 A g 1. It can be seen that H-FeP@C@GR nanocomposite greatly outperformed H-FeP −

and H-FeP@C anodes in terms of cycling stability and capacity, which demonstrated that the construction of coating carbon and encapsulated GR could progressively optimize the electrochemical performance of H-FeP nanospheres. Regarding H-FeP electrode, its discharge capacity faded rapidly from 1270 to 207 mAh g 1 after 100 −


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cycles, which was primarily ascribed to the poor conductivity, serious aggregation and pulverization of H-FeP nanospheres during the cycling. Once H-FeP nanospheres were coated with carbon layer to form H-FeP@C, the electrode exhibited decent cycling stability with enhanced capacity (586 mAh g 1 after 100 cycles). In the case of −

H-FeP@C@GR, owning to excellent electrical conductivity and 3D continuously interconnected macroporous structures of graphene with large surface area, it could deliver a discharge capacity of 771 mAh g 1 after 100 cycles at 0.2 A g 1, which was −

−

approximately 3.7 times higher than that of H-FeP nanospheres. Moreover, the cycling performance of GR was also measured for comparison. The discharge capacity of GR was about 240 mAh g 1 after 100 cycles. By taking into account of low mass percent −

and low capacity of GR in H-FeP@C@GR nanocomposite, it was concluded that GR contributed slightly to the lithium storage capacity of H-FeP@C@GR nanocomposite but played a crucial role on maintaining the structural integrity and enhancing conductivity of the active materials. The long-term cycling performance of the H-FeP@C@GR nanocomposite at a current density of 0.5 A g−1 was also evaluated (Figure 6e). It was noted that the discharge capacity of H-FeP@C@GR nanocomposite was 678 mAh g 1 in 50th cycle and kept at 542 mAh g 1 even after the 300th cycle. But −

−

for H-FeP@C nanospheres (Figure S8), the corresponding specific capacities were 340 mAh g 1 in 50th cycle and 115 mAh g 1 after 200th cycle, respectively. The rapid −

−

capacity decay behavior of H-FeP@C nanospheres could be attributed to the irreversible aggregation of nanospheres due to absence of 3D continuously interconnected structures GR, which markedly deteriorate full lithiation of their internal portions. To get insights into the differences in rate capabilities of H-FeP, H-FeP@C nanospheres, and H-FeP@C@GR nanocomposite, the electrochemical impedance spectroscopy (EIS) before and after the rate tests was carried out (Figure 7a


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and b). The EIS was fitted by an equivalent electrical circuit, as indicated in Figure 7c. For the equivalent electrical circuit, the intercept of the high-frequency semicircle on the Z’ axis can be attributed to the resistance of electrolyte (Rs). The semicircle in the high and middle frequency regions respectively represent the SEI layer resistance (Rf) and charge-transfer impedance on electrode-electrolyte interface (Rct), while the slope line at low frequency is related to the Warburg impedance (Zw) of the lithium-ion diffusion.55 Those values were also collected in Table S1. For H-FeP@C@GR electrode, it is noted that there was a slight change in its Rf after the rate tests, demonstrating that the electrode had a stabilized SEI layer. By contrast, an obvious change in Rf could be observed for H-FeP nanospheres anode. These results implied that repetitive breaks and growths of the SEI occurred during the rate tests of H-FeP nanospheres, which acted as a barrier to hinder Li+ diffusion and eventually resulted in poor rate capability. Like Rf, Rct value of H-FeP@C@GR nanocomposite was also slight changed after the rate tests, thus further confirming the stability of the structure. As for H-FeP@C nanospheres, although the Rct and Rf values showed small changes, the absence of 3D omnibearing conductive networks led to the aggregation of H-FeP nanospheres, inducing the destruction of electrical connections during prolonged cycling. The EIS results were in good agreement with the electrochemical behaviors of all the materials as anodes for LIBs. To further study structural evolution and morphology change at different discharge-charge stages in LIBs, ex-situ TEM/HRTEM images and corresponding SAED patterns of H-FeP@C@GR nanocomposite were displayed in Figure S9. After charged in 100th cycle to 3 V, shown in Figure S9a and b, H-FeP@C@GR remained spherical shape and porous structure without obvious aggregation, suggesting that carbon and graphene matrix into nanostructured FeP nanospheres could effectively


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maintain the structural stability for a very long cycling. In Figure S9c, a few lattice fringes were observed with d-spacing at 0.251 and 0.386 nm, assigned to the (102) and (101) planes of FeP (JCPDS: 65-2595), respectively. Furthermore, diffraction rings of FeP (111) and (211) were present in the SAED patterns shown in Figure S9d, consistent with ex-situ HRTEM, which proved complete delithiation of Li3P to reform highly crystalline FeP nanocrystals. After initial discharge to 0.8 V, HRTEM image in Figure S9e showed a few lattice fringes with a d-spacing at 0.380 nm, probably attributed to LixFeP (x=0-3) phase. Besides, the unclear patterns of inset SAED might be due to the intermediate LixFeP phase. After full discharge to 0.005V, in Figure S9f, lattice fringe with 0.209 nm corresponded to Li3P phase, and diffraction rings of SAED were consistent with Fe and Li3P phases. Electrochemical properties analysis of SIBs Inspired by the significant lithium storage performance of H-FeP@C@GR nanocomposite, all the samples were also examined as anode materials for SIBs as comparison. The electrochemical performance of H-FeP@C@GR nanocomposite was firstly investigated by CV at the scan rate of 0.1 mV s-1 in the 0-3.0 V voltage range (Figure 8a). An obvious difference between the first and the subsequent cycles was clearly seen. In first cathodic sweep, an apparent wide reduction peak appeared at approximately 0.59V, which could be attributed to the intercalation of Na+ ions into FeP to form NaxFeP (x=0-3) phase. The following peak at 0.23 V could be explained by the formation of SEI film on the electrode and the reduction of NaxFeP into Fe and Na3P phases. For the anodic curves, there were no apparent peaks, suggesting that in the whole oxidation process, the de-intercalation of Na+ from NaxFeP (x=0-3) phase to FeP was rather smooth without a fast reaction at a given voltage. In the subsequent cycles, it was noteworthy that the reduction peaks shifted to 1.0 V while the oxidation curves


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were overlapped, indicating an excellent reversible nature of the electrochemical reactions in the H-FeP@C@GR nanocomposite anode. Figure 8b showed the charge-discharge curves of H-FeP@C@GR nanocomposite at a current density of 0.1 A g 1. Clearly, the initial discharge and charge capacities achieved up to 995 and 656 −

mAh g 1, respectively, with an initial CE of 66 %. The large initial capacity loss was −

ascribed to the formation of irreversible Na2O and SEI layer. In the subsequent cycles, the discharge and charge capacities of H-FeP@C@GR nanocomposite decreased slightly, further indicating its high capacity and excellent reversibility. Besides, in discharge curves, there were small plateaus at around 1.16 V, matching with CV results. But no plateaus could be seen in the charge curves. The rate capabilities of three electrodes were evaluated at a series of current densities shown in Figure 8c. Clearly, the discharge capacities of H-FeP@C@GR nanocomposite were 620 (10th), 487 (20th), 366 (30th), 285 (40th), and 237 (50th) mAh g 1 at current densities of 0.1, 0.2, 0.4, 0.8, and 1.6 A g 1, respectively. As the rate −

−

decreased back to 0.1 A g 1, the capacity could be as high as 480 mAh g 1 and remained −

−

stable in the following cycles. The discharge capacities of H-FeP@C and H-FeP at various current densities were listed in Table S2. Owing to structural features, the as-obtained H-FeP@C@GR electrode featured more superior cycle stability and higher rate capacity than H-FeP@C and H-FeP electrodes. Figure 8d showed long term cycling performance of H-FeP@C@GR electrode at a current density of 0.1 A g 1. It −

was obvious that the capacity of H-FeP@C@GR nanocomposite faded significantly during the course of the first few cycles, and showed a slight decay from the 40th cycle onwards. Whereas, H-FeP nanospheres showed continuous capacity decay to 141 mAh g 1 under the same conditions, revealed in Figure S10. After 250 cycles, the discharge −


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capacity of H-FeP@C@GR electrode was retained at 400 mAh g 1. As for H-FeP@C, −

the first and the second discharge capacities were 943 and 633 mAh g 1 at a current −

density of 0.1 A g 1, respectively, whereas it was only 184 mAh g 1 after 100 cycles −

−

(Figure S10). Clearly, the cycling stability and specific capacity of H-FeP@C@GR nanocomposite were higher than those of H-FeP@C and H-FeP nanospheres, mainly due to the wrapping of GR that acted as a cage-like framework around H-FeP@C nanospheres so as to prevent the agglomeration of active materials and cracking of the electrode. The reversible capacities of H-FeP@C@GR electrode for SIBs were much lower than those of their counterparts in LIBs (Table S3), which was because of the large radius of Na+ (2.04 Å), unavoidably leading to a large volume change and poor kinetics of sodiation/desodiation processes upon cycling. This could be further confirmed by the comparison of the EIS spectra of H-FeP@C@GR nanocomposite in SIBs and LIBs. As shown in Figure 8e, it is evident that the charge-transfer resistance of LIBs was much lower than that of SIBs, which was consistent with the electrochemical behaviors. To further study structural evolution and morphology change at different discharge-charge stages in SIBs, ex-situ TEM/HRTEM images and corresponding SAED patterns of H-FeP@C@GR were displayed in Figure 9. After H-FeP@C@GR was charged to 3 V in 100th cycle, ex-situ TEM image (Figure 9a) showed that it remained spherical shape and porous structure and that the size was almost unchanged, suggesting that carbon layer and graphene into nanostructured FeP can effectively maintain the structural stability for a very long cycling. In Figure 9b, a few lattice fringes were observed with d-spacing at 0.251 and 0.24 nm, assigned to the (102) and (101) planes of FeP (JCPDS: 65-2595), respectively. Moreover, diffraction rings of FeP (111), (202), (211), and (212) were present in the SAED patterns shown in Figure


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9c, consistent with ex-situ HRTEM, proving the complete desodiation of Na3P to reform highly polycrystalline FeP nanocrystals. Ex-situ XRD and IR were also used to further investigate the H-FeP@C@GR nanocomposite after being charged to 3 V in 100th cycle. As shown in Figure S11a, ex-situ XRD pattern (I) indicated that the diffraction peaks at 47.0o and 48.4o could be assigned to (202) and (211) plane of FeP (JCPDS: 65-2595), respectively, in addition to the strong peak of collector Ni, which corroborated that the Na3P were desodiated fully to reform polycrystalline FeP nanocrystals. Moreover, in Figure S11b, ex-situ IR spectrum (I) showed that a broad absorption peak in the range of 500-700 cm-1, ascribed to Fe-P vibration, was observed, consistent with the results above. Some other absorptions might be due to the existence of polymer binder in the sample. After it was initially discharged to 0.8V, shown in Figure 9d, ex-situ HRTEM of nanosphere edge displayed that a few lattice planes were found with a d-spacing at 0.202 nm corresponding to Fe (002) plane, and at 0.308 and 0.430 nm corresponding to (102) and (100) planes of Na3P, respectively, which elaborated that edge portion of FeP nanosphere was fully sodiated to form Fe and Na3P phases at this condition. Meanwhile, a single FeP nanosphere was characterized by SAED. The diffraction rings of Fe (101) and Na3P (102) and (110) were present in the SAED patterns shown in Figure 9e, verifying the existence of Fe and Na3P phases. Another unknown diffraction ring might be due to NaxFeP (x=0-3) phase, which was mainly present in internal portion of FeP nanosphere. Likewise, after H-FeP@C@GR was initially discharged to 0.8V, ex-situ XRD pattern (II) showed that diffraction peak at 49.2o could be indexed to (102) plane of Fe (JCPDS: 50-1275). From ex-IR spectrum (II) of Figure S11b, the absorption peak could be attributed to NaxFeP (x=0-3) phase. Furthermore, after it was fully discharge to 0.005 V, shown in Figure 9f and g, lattice planes matching with Fe and Na3P phases were found. SAED patterns of a single


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nanosphere showed only Fe and Na3P phases without presence of NaxFeP (x=0-3), demonstrating the full sodiation of FeP, which is verified by the ex-situ XRD and IR in Figure S11. Additionally, after 100th cycle of charging to 3 V, H-FeP@C@GR nanocomposite was analyzed by SEM and EDX, shown in Figure S12. It can be seen that the FeP nanospheres were unchanged without collapse and aggregation in term of shape, size and hollow structure. From EDX spectrum, the atomic ratio of Na:Fe:P was obtained 0.26:1:1.07, implying that a certain amount of Na+ ions was stored in FeP nanocrystal without desodiation even after full charging, which might explain the capacity loss after cycling. More importantly, in the case of H-FeP@C@GR nanocomposite, SEM equipped with energy-dispersive X-ray (EDX) was used to investigate the amount of Na+ ions intercalated into FeP nanocrystals, i.e. x variation in NaxFeP (x=0-3), at different discharging voltage in its first discharge cycle, the result of which was displayed in Figure 10. SEM images demonstrated that FeP nanospheres remained the spherical morphology and the same size after discharging. EDX spectra were collected from corresponding SEM images. The calculated atomic ratios of Na:Fe:P were 0.86:1:1.06, 2.09:1:1.07, 2.97:1:1.07 at discharge voltage 1.5, 0.8 and 0.005 V, respectively. The curve of voltage as a function of x variation in NaxFeP (x=0-3) was plotted, the trend of which was consistent with that of discharge curve. From above assessments, it has been established that H-FeP@C@GR showed more enhanced specific capacity, rate capability, and cycling stability in LIBs than in SIBs, which is due to that the more sluggish kinetics occurs in insertion/extraction of SIBs than that of LIBs. To demonstrate this, we carried out density functional theory (DFT) calculations and simulated the adsorption of Li+ ions and Na+ ions on the FeP surface as well as Li+/Na+ ions diffusion in FeP. By comparing all the possible adsorption sites, we concluded that the adsorptions of Li+ and Na+ ions at the sites of P atoms were the


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most energetically stable. The Li+ ions adsorption energy was calculated to be 0.21 eV, while the Na+ ions adsorption energy was calculated to be -0.12 eV, indicating that Na-FeP adsorption interactions were more favorable than Li-FeP interactions, which could be explained by the fact that Na has more electrons and larger radius than Li. Additionally, the simulation on Li+/Na+ ions diffusion in FeP crystal was carried out. To simulate the diffusion processes of Li+ and Na+ in bulk FeP, firstly, 2a×b×2c supercell of orthorhombic FeP cell (space group Pnma(62)) was built using the experimental parameters and one Li impurity was interstitially doped into it to find the most stable doped position. We considered the system of Li+ located on the most stable doped position as a reactant and the system of Li+ located on the neighboring equivalent position as a product. Ten images were inserted between the reactant and product, and the climbing image scheme was used to search the transient state. Moreover, to compare the diffusion process of Li+ and Na+ in bulk FeP, we computed the diffusion process of Na+ using the diffusion path of the Li+. As schematically shown in the inset of Figure 11a and b, the diffusion barrier energy of lithium ion in the optimized process was 11.133 eV, while the diffusion barrier energy of sodium ion was 16.711 eV. This indicated that the Li+ ion with smaller diffusion barrier energy can achieve faster charge and discharge than Na+ ion, consequently boosting the power rate performance of the lithium ion batteries. Meanwhile, our DFT calculations showed that the large and advantage cavity structure of bulk FeP were more favorable for the storage of lithium ions. Based on the above simulation, we can see that compared to high diffusion barrier energy (11.133 eV for Li+ and 16.711 eV for Na+), the adsorption energy (0.21 eV for Li+ and -0.12 eV for Na+) could be neglected, suggesting that ion diffusion barrier energy was much more crucial parameter for the battery performance. Overall, diffusion barrier energy of Li+ was


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lower than that of Na+, leading to that FeP demonstrates more excellent performance in LIBs than in SIB, matching with the experimental results. Conclusions In conclusion, we successfully designed a 3D GR encapsulated with hollow FeP@carbon nanocomposite via a combination of hydrothermal reaction, carbon coating process, phosphidation treatment, and carbothermic reaction. SEM and TEM confirmed that H-FeP@C nanospheres were strongly anchored on the surface of the stacked structure as well as in the parallel layers of the GR. This 3D network structure can not only naturally accommodate the lithiation/sodiation induced volume change, but also offer more active sites for lithium/sodium ion insertion and electron pathway for electron and ion transportation. As a result, when tested in LIBs and SIBs, H-FeP@C@GR electrode yielded an excellent electrochemical performance in terms of reversibility, rate capacity and cycling performance. For example, H-FeP@C@GR electrode showed a specific capacity about 771 and 446 mAh g 1 after 100 cycles for −

LIBs and SIBs (much higher than those of H-FeP and H-FeP@C), respectively. Besides, H-FeP@C@GR electrode could also afford the high current density. When H-FeP@C@GR nanocomposite was used as an anode material for LIBs at different current density in charge/discharge processes (0.2, 0.5, 1, 2, 4, and 8 A g 1), the capacity −

still can be recovered to 953 mAh g 1. −

Methods Materials. Synthesis of hollow Fe3O4 (H-Fe3O4) nanospheres: The nanospheres were synthesized by solvothermal method. Briefly, 2.70 g of FeCl3·6H2O and 7.20 g of sodium acetate were dissolved in 100 mL of ethylene glycol under magnetic stirring. After stirring for 1 h, the resulting homogeneous yellow solution was transferred to a Teflon-lined stainless-steel autoclave. The autoclave was heated at 200 °C for 8 h and


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then naturally cooled to ambient temperature. Finally, the resulting H-Fe3O4 nanospheres were washed seven times with ethanol and dried in vacuum at 80 °C for 10 h. Synthesis of H-Fe3O4@C nanospheres: Typically, 0.05g as-synthesized H-Fe3O4 and 0.5 g glucose were dissolved in a solution containing 17.5 mL of deionized water and 5 mL of ethanol with a following vigorous stirring for 30 min. Then, the solution was placed in a 50 mL Teflon-sealed autoclave and maintained at 180 °C for 2 h. The precipitates were centrifuged and washed with deionized water and ethanol and dried under vacuum at 80 °C for 10 h. Finally, the resulting samples were calcined at 500 °C for 4 h in flowing argon. Synthesis of H-FeP and H-FeP@C nanospheres: The H-FeP nanospheres were fabricated by the phosphidationof H-Fe3O4. In a typical procedure, the obtained H-Fe3O4 nanospheres and NaH2PO2 were placed at two separate positions in one closed porcelain crucible with NaH2PO2 at the upstream side of the furnace. The weight ratio of the H-Fe3O4 nanospheres to NaH2PO2 was 1:20. Subsequently, the samples were heated at 300 ºC for 3 h with a heating speed of 2 ºC min−1 in a static Ar atmosphere. At last, the black H-FeP nanospheres were obtained after cooled to ambient temperature under Ar flow. The H-FeP@C nanospheres were synthesized by the same procedure using H-Fe3O4@C nanospheres as starting material. Synthesis

of

3D

H-FeP@C@GR

nanocomposite:

The

H-FeP@C@GR

nanocomposite was prepared through direct annealing of H-FeP@C nanospheres that were simply mixed with GO sheets obtained by our previous work.56 Typically, 80 mg of H-FeP@C nanospheres were firstly dispersed in 100 mL of distilled water and ultrasonically treated for 2 h. Then, 20 mL of homogeneous GO (7.5 mg mL−1) aqueous suspension was added into the H-FeP@C nanospheres suspension under magnetic


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stirring for 6 h. After that, the mixed suspension was collected by freeze-drying and then annealed at 600 °C for 2 h in Ar atmosphere to finally obtain 3D H-FeP@C@GR nanocomposite. Material Characterization. The phase purity and crystal structure of samples were measured by a Bruker D8 ADVANCE X-ray powder diffractometer (XRD) with Cu Kα radiation (λ = 0.15418 nm) at a scanning rate of 0.02o s−1 in the 2θ range from 10 to 80o. Scanning electron microscopy (SEM) images of products were obtained using a FEI Quanta 400 ESEM-FEG instrument with an accelerating voltage of 20 kV. Transmission electron microscopy (TEM) images and high-resolution transmission electron microscopy (HRTEM) images were obtained by JEOL JEM-3010 instrument. The X-ray photoelectron spectroscopy (XPS) experiments were carried out on a PHI-5400 electron spectrometer. The Raman spectrum was performed by Raman spectrometer with 532 nm laser excitation. Thermogravimetry analysis (TGA) of the product was conducted from room temperature up to 800 oC with a heating rate of 10 oC min−1 under flowing air (TGA, Q 600). Fourier transform infrared (FT-IR) spectra were recorded on a Bruker Optics Tensor 27 FT-IR spectrometer (Germany). Electrochemical measurements. The electrochemical behavior was performed with 2025 coin-type cells assembled in a glove box in argon atmosphere. The working electrode was prepared by mixing active materials (80 wt%) with acetylene black (10 wt%) and polytetrafluoroethylene (PVDF, 10 wt%) to form a slurry. And the obtained slurry was spread uniformly on a circular piece of nickel foam with 14 mm diameter. The nickel foam was pressed at 20 MPa so as to obtain good contact between nickel foam and then dried at 80 °C in a vacuum oven for 12 h. The active material loading in each electrode was about 1.2 mg cm−2. The electrolyte involves LiPF6 (1 M) dissolved in a mixture of dimethyl carbonate, diethyl carbonate and ethylene carbonate (1:1:1 by


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volume), and the separator was microporous polypropylene film. As for sodium batteries, the electrolyte consisted of a solution of 1 M NaClO4 in ethylene carbonate/propylene carbonate (1/1; in volume). The charge-discharge tests were carried out on a LAND battery program-control test system in a cut-off potential window of 0-3.0 V after aging for more than 10 h. The cyclic voltammetry (CV) between 0.005 and 3.0 V vs at a scan rate of 0.1 mV s−1 and electrochemical impedance spectroscopy (EIS) with the frequency range from 0.01 to 100 kHz were carried out on a CHI 660D electrochemical work station. Computational Details. To simulate the diffusion process of Li+ in bulk FeP, the climbing image nudged elastic band (CI-NEB) method was used. All calculations were performed based on the density functional theory (DFT) within the open-source code Quantum ESPRESSO,57 with the GBRV ultrasoft pseudopotentials (USPP)58 to describe the electron-ionic core interaction. We adopted generalized gradient approximation (GGA) of Perdew-Burke-Ernzerhof for solid (PBEsol)59 to describe the exchange-correlation interaction of electrons and the wave functions were expanded in a plane-wave basis set with an energy cutoff of 40 Ry to ensure accurate results. The force on each ion was converged to be less than 0.001 Ry/a.u. and all the geometric structures were fully relaxed to minimize the total energy of the system until a precision of 10-4 Ry is reached. The Fe 3s23p63d64s2, P 3s23p3 and Li 1s22s1 electrons were treated as valence electrons. A 2×4×4 k-point grid in reciprocal space is used to ensure the convergence for the total energy self-consistent calculations. Acknowledgements This work was supported by the National Natural Science Foundation of China (No. 51672213), the Key Project of Research and Development of Shaanxi Province (No. 2017ZDCXL-GY-08-01), the Key Science and Technology Innovation Team Project


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of Natural Science Foundation of Shaanxi Province (No. 2017KCT-01), and the Natural Science Foundation of Shaanxi Province (No. 2017JM2025). Associated Content Supporting Information Characterizations of XRD, SEM, TEM, HRTEM, SAED, surface wettability, and electrical conductivities for prepared materials; electrochemical measurements including CV curves, cycling performances and rate capabilities of FeP-based materials (PDF) References (1) Chao, D. L.; Zhu, C. R.; Xia, X. H.; Liu, J. L.; Zhang, X.; Wang, J.; Liang, P.; Lin, J. Y.; Zhang, H.; Shen, Z. X. Graphene Quantum Dots Coated VO2 Arrays for Highly Durable Electrodes for Li and Na Ion Batteries. Nano Lett. 2015, 15, 565-573. (2) Wang, X. J.; Liu, X. J.; Wang, G.; Xia, Y.; Wang, H. One-Dimensional Hybrid Nanocomposite of High-Density Monodispersed Fe3O4 Nanoparticles and Carbon Nanotubes for High-Capacity Storage of Lithium and Sodium. J. Mater. Chem. A 2016, 4, 18532-18542. (3) Fu, K.; Xue, L. G.; Yildiz, O.; Li, S.; Lee, H.; Li, Y.; Xu, G. J.; Zhou, L.; Bradford, P. D.; Zhang, X. W. Effect of CVD Carbon Coatings on Si@CNF Composite as Anode for Lithium-ion Batteries. Nano Energy 2013, 2, 976-986. (4) Hwang, S. M.; Kim, S. Y.; Kim, J. G.; Kim, K. J.; Lee, J. W.; Park, M. S.; Kim, Y. Jun.; Mohammed, S.; Yamauchi, Y.; Kim, J. H. Electrospun Manganese-Cobalt Oxide Hollow Nanofibres Synthesized Combustion Reactions and Their Lithium Storage Performance. Nanoscale 2015, 7, 8351-8355. (5) Xue, H. R.; Zhao, J. Q.; Tang, J.; Gong, H.;

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Porous Carbon for Superior Lithium-Ion Batteries. Chem. Eur. J. 2016, 22, 4915-4923. (6) Zuo, X. X.; Xia, Y. G.; Ji, Q.; Gao, X.; Yin, S. S.; Wang, M. M.; Wang, X. Y.; Qiu, B.; Wei, A. X.; Sun, Z. C.; Liu, Z. P.; Zhu, J.; Cheng, Y. J. Self-Templating Construction of 3D Hierarchical Macro-/Mesoporous Silicon from 0 D Silica Nanoparticles. ACS Nano 2017, 11, 889-899. (6) Wang, B.; Li, X. L.; Zhang, X. F.; Luo, B.; Jin, M. H.; Liang, M. H.; Dayeh, S. A.; Picraux, S. T.; Zhi, L. J. Adaptable Silicon Carbon Nanocables Sandwiched between Reduced Graphene Oxide Sheets as Lithium Ion Battery Anodes. ACS Nano 2013, 7, 1437-1445. (8) Li, Q.; Li, Z. Q.; Zhang, Z. W.; Li, C. X.; Ma, J. Y.; Wang, C. X.; Ge, X. L.; Dong, S. H.; Yin, L. W. Low-Temperature Solution-Based Phosphorization Reaction Route to Sn4P3/Reduced Graphene Oxide Nanohybrids as Anodes for Sodium Ion Batteries. Adv. Energy Mater. 2016, 6, 1600376-1600385. (9) Wu, C.; Kopold, P.; Aken, P. A.; Maier, J.; Yu, Y. High Performance Graphene/Ni2P Hybrid Anodes for Lithium and Sodium Storage through 3D Yolk-Shell-Like Nanostructural Design. Adv. Mater. 2017, 29, 1604015-1604022. (10) Fan, M. P.; Chen, Y.; Xie, Y. H.; Yang, T. Z.; Shen, X. W.; Xu, N.; Yu, H. Y.; Yan, C. L. Half-Cell and Full-Cell Applications of Highly Stable and Binder-Free Sodium Ion Batteries Based on Cu3P Nanowire Anodes. Adv. Funct. Mater. 2016, 26, 5019-5027. (11) Wang, X.; Sun, P. P.; Qin, J. W.; Wang, J. Q.; Xiao, Y.; Cao, M. H. A Three-Dimensional Porous MoP@C Hybrid as a High-Capacity, Long-Cycle Life Anode Material for Lithium-Ion Batteries. Nanoscale 2016, 8, 10330-10338. (12) Cui, Y. H.; Xue, M. Z.; Fu, Z. W.; Wang, X. L.; Liu, X. J. Nanocrystalline CoP


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Sodium-Ion Batteries. Energy Environ. Sci. 2015, 8, 3531-3538. (21) Lu, Y.; Tu, J. P.; Gu, C. D.; Wang, X. L.; Mao, S. X. In Situ Growth and Electrochemical Characterization Versuslithium of a Core/Shell-Structured Ni2P@C Nanocomposite Synthesized by a Facile Organic-Phase Strategy. J. Mater. Chem. 2011, 21, 17988-17997. (22) Fullwnwarth, J.; Darwiche, A.; Soares, A.; Donnadieu, B.; Monconduit, L. NiP3: a Promising Negative Electrode for Li- and Na-Ion Batteries. J. Mater. Chem. A 2014, 2, 2050-2059. (23) Hwang, S. M.; Lim, Y. G.; Kim, J. G.; Heo, Y. U.; Lim, J. H.; Yamauchi, Y.; Park, M. S.; Kim, Y. J.; Dou, S. X.; Kim, J. H. A Case Study on Fibrous Porous SnO2 Anode for Robust, High-Capacity Lithium-Ion Batteries. Nano Energy 2014, 10, 53-62. (24) Liu, Q.; Tian, J.; Cui, W.; Jiang, P.; Cheng, N.; Asiri, A. M.; Sun, X. Carbon Nanotubes Decorated with CoP Nanocrystals: A Highly Active Non-Noble-Metal Nanohybrid Electrocatalyst for Hydrogen Evolution. Angew. Chem. Int. Ed. 2014, 53, 6710-6714. (25) Han, F.; Tan, C. Y. J.; Gao, Z. Improving the Specific Capacity and Cyclability of Sodium-Ion Batteries by Engineering a Dual-Carbon Phase-Modified Amorphous and Mesoporous Iron Phosphide. ChemElectroChem 2016, 3, 1054-1062. (26) Wen, Z. H.; Wang, Q.; Zhang, Q.; Li, J. H. In Situ Growth of Mesoporous SnO2 on Multiwalled Carbon Nanotubes: A Novel Composite with Porous-Tube Structure as Anode for Lithium Batteries. Adv. Funct. Mater. 2007, 17, 2772-2778. (27) Liu, Z.; Yu, X. Y.; Paik, U. Etching-in-a-Box: A Novel Strategy to Synthesize Unique Yolk-Shelled Fe3O4@Carbon with an Ultralong Cycling Life for Lithium Storage. Adv. Energy Mater. 2016, 6, 1502318-1502326.


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(28) Ng, S. H.; Wang, J. Z.; Wexler, D.; Konstantinov, K.; Guo, Z. P.; Liu, H. K. Highly

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Lithium

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Spheroidal

Carbon-Coated

Silicon

Nanocomposites as Anodes for Lithium-Ion Batteries. Angew. Chem. Int. Ed. 2006, 45, 6896-6899. (29) Li, D.; Wang, H.; Liu, H. K.; Guo, Z. A New Strategy for Achieving a High Performance

Anode

for

Lithium

Ion

Batteries-Encapsulating

Germanium

Nanoparticles in Carbon Nanoboxes. Adv. Energy Mater. 2016, 6, 1501666-1501672. (30) Derrien, G.; Hassoun, J.; Panero, S.; Scrosati, B. Nanostructured Sn-C Composite as an Advanced Anode Material in High-Performance Lithium-Ion Batteries. Adv. Mater. 2007, 19, 2336-2340. (31)

Wang,

B.;

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Polyethyleneimine-Mediated Synthesis of Ultrathin Hexagonal Co3O4Nanosheets with Reactive Facets for Lithium-Ion Batteries. ChemElectroChem 2016, 3, 55-65. (32) Stankovich, S.; Dikin, D. A.; Dommett, G. H. B.; Kohlhaas, K. M.; Zimney, E. J.; Stach, E. A.; Piner, R. D.; Nguyen, S. T.; Ruoff, R. S. Graphene-Based Composite Materials. Nature 2006, 442, 282-286. (33) Zhang, Y.; Chen, P. H.; Gao, X.; Wang, B.; Liu, H.; Wu, H.; Liu, H. K.; Dou, S. X. Nitrogen-Doped Graphene Ribbon Assembled Core-Sheath MnO@Graphene Scrolls as Hierarchically Ordered 3D Porous Electrodes for Fast and Durable Lithium Storage. Adv. Funct. Mater. 2016, 26, 7754-7765. (34) Shen, L. F.; Zhang, X. G.; Li, H. S.; Yuan, C. Z.; Cao, G. Z. Design and Tailoring of a Three-Dimensional TiO2-Graphene-Carbon Nanotube Nanocomposite for Fast Lithium Storage. Phys J. Chem. Lett. 2011, 2, 3096-3101. (35) Jiang, P.; Liu, Q.; Liang, Y.; Tian, J.; Asiri, A. M.; Sun, X. A Cost-Effective 3D Hydrogen Evolution Cathode with High Catalytic Activity: FeP Nanowire Array as the


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Active Phase. Angew. Chem. Int. Ed. 2014, 53, 12855-12859. (36) Yan, Y.; Xia, B. Y.; Ge, X.; Liu, Z.; Fisher, A.; Wang, X. A Flexible Electrode Based on Iron Phosphide Nanotubes for Overall Water Splitting. Chem. Eur. J. 2015, 21, 18062-18067. (37) Jiang, J.; Wang, C. D.; Liang, J. W.; Zuo, J.; Yang, Q. Synthesis of Nanorod-FeP@C Composites with Hysteretic Llithiation in Lithium-Ion Batteries. Dalton Trans. 2015, 44, 10297-10303. (38) Toprakci, O.; Ji, L.; Lin, Z.; Toprakci, H. A.; Zhang, X.; Fabrication and Electrochemical Characteristics of Electrospun LiFePO4/Carbon Composite Fibers for Lithium-Ion Batteries. J. Power Sources 2011, 196, 7692-7699. (39) Kudin, K. N.; Ozbas, B.; Schniepp, H. C.; Prudhomme, R. K.; Aksay, I. A.; Car, R. Raman Spectra of Graphite Oxide and Functionalized Graphene Sheets. Nano Lett. 2008, 8, 36-41. (40) Liu, J. L.; Feng, H. B.; Wang, X. P.; Qian, D.; Jiang, J. B.; Li, J. H.; Peng, S. J.; Deng, M.; Liu, Y. C. Self-Assembly of Nano/Micro-Structured Fe3O4 Microspheres among 3D rGO/CNTs Hierarchical Networks with Superior Lithium Storage Performances. Nanotechnology 2014, 25, 225401-225410. (41) Wang, B.; Abdulla, W. A.; Wang, D. L.; Zhao, X. S. AThree-Dimensional Porous LiFePO4 Cathode Material Modified with a Nitrogen-Doped Graphene Aerogel for High-Power Lithium Ion Batteries. Energy Environ. Sci. 2015, 8, 869-875. (42) Thomsen, C.; Reich, S. Double Resonant Raman Scattering in Graphite. Phys. Rev. Lett. 2000, 85, 5214-5225. (43) Li, Z. Q.; Zhang, L. Y.; Ge, X. L.; Li, C. X.; Dong, S. H.; Wang, C. X.; Yin, L. W. Core-Shell Structured CoP/FeP Porous Microcubes Interconnected by Reduced Ggraphene Oxide as High Performance Anodes for Sodium Ion Batteries. Nano


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Energy 2017, 32, 494-502. (44) Cheng, J. L.; Wang, B.; Park, C. M.; Wu, Y. P.; Huang, H.; Nie, F. D. CNT@Fe3O4@C Coaxial Nanocables: One-Pot, Additive-Free Synthesis and Remarkable Lithium Storage Behavior. Chem. Eur. J. 2013, 19, 9866-9874. (45) Park, D. Y.; Myung, S. T. Carbon-Coated Magnetite Embedded on Carbon Nanotubes for Rechargeable Lithium and Sodium Batteries. ACS Appl. Mater. Interfaces 2014, 6, 11749-11757. (46)

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(52) Hu, X. B.; Ma, M. H.; Zeng, M. Q.; Sun, Y. Y.; Chen, L. F.; Xue, Y. H.; Zhang, T.; Ai, X. P.; Mendes, R. G.; Rummeli, M. H.; Fu, L. Supercritical Carbon Dioxide Anchored Fe3O4 Nanoparticles on Graphene Foam and Lithium Battery Performance. ACS Appl. Mater. Interfaces 2014, 6, 22527-22533. (53) Wang, B. B.; Wang, G.; Wang, H. Hybrids of Mo2C Nanoparticles Anchored on Ggraphene Sheets as Anode Materials for High Performance Lithium-ion Batteries. J. Mater. Chem. A 2015, 3, 17403-17411. (54) Su, Y. Z.; Li, S.; Wu, D. Q.; Zhang, F.; Liang, H. W.; Gao, P. F.; Cheng, C.; Feng, X. L. Two-Dimensional Carbon-Coated Graphene/Metal Oxide Hybrids for Enhanced Lithium Storage. ACS Nano 2012, 6, 8349- 8356. (55) Xia, Y.; Xiao, Z.; Dou, X.; Huang, H.; Lu, X. H.; Yan, R. J.; Gan, Y. P.; Zhu, W. J.; Tu, J. P.; Zhang, W. K.; Tao, X. Y. Green and Facile Fabrication of Hollow Porous MnO/C Microspheres from Microalgaes for Lithium-Ion Batteries. ACS Nano 2013, 7, 7083-7079. (56) Wang, J. X.; Wang, H. Synthesis of Free-Standing Reduced Graphene Oxide Membranes with Different Thicknesses and Comparison of Their Electrochemical Performance as Anodes for Lithium-Ion Batteries. RSC Adv. 2015, 5, 30084-30091. (57) Giannozzi, P.; Baroni, S.; Bonini, N.; Calandra, M.; Car, R.; Cavazzoni, C.; Ceresoli, D.; Guido, L. C.; Cococcioni, M.; Dabo, I.; Corso, A. D.; Fabris, S.; Fratesi, G.; de Gironcoli, S.; Gebauer, R.; Gerstmann, U.; Gougoussis, C.; Kokalj, A.; Lazzeri, M.; Martin-Samos, L.; Marzari, N.; Mauri, F.; Mazzarello, R.; Paolini, S.; Pasquarello, A.; Paulatto, L.; Sbraccia, C.; Scandolo, S.; Sclauzero, G.; Seitsonen, A. P.; Smogunov, A.; Umari, P.; Wentzcovitch, R. M. QUANTUM ESPRESSO: A Modular and Open-Source Software Project for Quantum Simulations of Materials. J. Phys.: Condens. Matter 2009, 21, 395502-395521.


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(58) Garrity, K. F.; Bennett, J. W.; Rabe, K. M.; Vanderbilt, D. Pseudopotentials for High-throughput DFT Calculations. Comput. Mater. Sci. 2014, 81, 446-452. (59) Perdew, J. P.; Ruzsinszky, A.; Csonka, G. I.; Vydrov, O. A.; Scuseria, G. E.; Constantin, L. A.; Zhou, X.; Burke, K. Restoring the Density-Gradient Expansion for Exchange in Solids and Surfaces. Phys. Rev. Lett. 2008, 100, 136406-136409.


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Figure 1 (a) Crystal structure evolution of FeP during the conversion reaction with lithium or sodium. (b) Schematic illustration of the synthesis process for H-FeP@C@GR nanocomposite. Figure 2 SEM images of (a-c) H-FeP nanospheres, (d-f) H-FeP@C nanospheres, and (g) H-FeP@C@GR nanocomposite. (h) Element mapping images and (i) EDS spectrum of H-FeP@C@GR nanocomposite. Figure 3 TEM and HRTEM images of (a-c) H-FeP nanospheres, (d-f) H-FeP@C nanospheres,

(g-h)

H-FeP@C@GR

nanocomposite.

(i)

SAED

patterns

of

H-FeP@C@GR nanocomposite. Figure 4 (a) XRD patterns and (b) TGA profiles of H-FeP nanospheres, H-FeP@C nanospheres, and H-FeP@C@GR nanocomposite. (c) Raman spectra of H-FeP@C and H-FeP@C@GR nanocomposite. Figure 5 XPS spectra of H-FeP@C@GR nanocomposite (a) survey and (b) Fe 2p, (c) P 2p, and (d) C 1s. Figure 6 Electrochemical performances of H-FeP@C@GR nanocomposite for LIBs; (a) Representative CV curves of H-FeP@C@GR nanocomposite at 0.1 mV s 1 within −

0.005 and 3.0 V. (b) Selected discharge-charge curves for the initial three cycles of the H-FeP@C@GR nanocomposite at a current density of 0.2 A g 1. (c) Rate −

performance and (d) Cycling performance of all the samples. (e) Long cycling performance of H-FeP@C@GR nanocomposite at a current density of 0. 5 A g 1. −

Figure 7 Nyquist plots of H-FeP nanospheres, H-FeP@C nanospheres, and H-FeP@C@GR nanocomposite (a) before and (b) after the rate tests. (c) The equivalent circuit used for fitting the experimental data. Figure 8 Electrochemical performances of H-FeP@C@GR nanocomposite for SIBs; (a) Representative CV curves of H-FeP@C@GR nanocomposite at 0.1 mV s 1 within −


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0.005 and 3.0 V. (b) Discharge-charge voltage profiles of H-FeP@C@GR nanocomposite at 0.1 A g 1. (c) Rate performance and (d) Long term cycling property −

of H-FeP@C@GR electrode at a current density of 0.1 A g 1. (e) The Nyquist plots of −

H-FeP@C@GR electrode in LIBs and SIBs. Figure 9 Ex-situ TEM/HRTEM images and corresponding SAED patterns of the H-FeP@C@GR at different discharge-charge stages in SIBs. (a-c)100th charged to 3 V, (d-e) initial discharged to 0.8 V, (f-g) initial discharged to 0.005 V. Figure 10 Ex-situ SEM and EDX spectra of H-FeP@C@GR electrode collected at various points in SIBs. (a) after first discharging to 1.5 V, (b) after first discharging to 0.8 V, (c) after first discharging to 0.005 V. Figure 11 Diffusion barrier energy (△E) of Li+ ions (a) and Na+ ions (b) along the optimized path in the bulk FeP.


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Figure 1 (a) FeP

p

Li/Na-insertion

Fe

LiFeP/NaFeP

Conversion

Li3P/Na3P+Fe

Li/Na

(b)

GO Carbon-coating

H-Fe3O4

Phosphatizing

Stirring

NaH2PO2

Calcination

H-Fe3O4@C

H-FeP@C


H-FeP@C@GR

ACS Nano

Figure 2 (a)

(b)

(c)

1µm

500 nm

100 nm (f)

(e)

(d)

1µm (g)

500 nm (h)

Fe 1µm

P

C

100 nm (i)

cps (eV)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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1µm


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Figure 3


ACS Nano

Figure 4 130

(b)

Weight Loss (%)

(a)

H-FeP@C@GR

Intensity (a. u.)

120

9.4 %

100

14.4 %

90 80

H-FeP@C

26.8 %

H-FeP H-FeP@C H-FeP@C@GR

110

0

200

400 600 o Temperature ( C)

800

(c) Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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H-FeP

I /I =1.07 D G

H-FeP@C@GR I /I =0.72 D G

JCPDS No:65-2595

20

30

40 50 2θ (degree)

60

G band

D band

H-FeP@C

70

1200

1400 1600 -1 Raman shift (cm )


1800

Page 41 of 47

Figure 5

(a)

(b)

Fe 2p

Intensity (a.u.)

Intensity (a.u.)

C 1s

O 1s Fe 2p

P 2p P 2s

0

200

400 600 Binding energy

(c)

Fe-P Fe-O

800 705

P-O

P 2p3/2 P 2p1/2

710

715 720 725 Binding energy

730

(d)

P 2p Intensity (a.u.)

Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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735 C1s

C=C/C-C

C-O C=O O-C=O

128


140 282

284



290

292

ACS Nano

Figure 6 0.4

3.0 Potential vs. Li/Li (V)

(a)

1st 2nd 3rd

-1.2 0.0

0.5 1.0 1.5 2.0 2.5 + Potential vs. Li/Li (V)

2.0 1.5 1.0

0.0

3.0

1st 2nd 3rd

0.5 0

1800

-1

(c)

H-FeP@C@GR H-FeP@C H-FeP

1500 0. 2 A g

1200

-1 -1

0. 2 A g

-1

0. 5 A g

900

-1

1Ag

2Ag

600

-1

4Ag

-1 -1

8Ag

300

-1 Specific capacity (mAh g )

10

20

30 40 50 Cycle number

60

1600

H-FeP@C@GR H-FeP@C H-FeP GR

1500 1200 900 600 300 0

0

(d)

-1

1800

400 800 1200 -1 Specific capacity (mAh g )

70

0

20

40 60 Cycle number

80

100

1200

100

1000 (e)

90 80

800

H-FeP@C@GR

70

600

60 50

400

Charge

Discharge Current density: 0.5 A g-1

200

40 30

0 0

50

100

150 200 Cycle number


250

20 300

Coulombic efficiency (%)

-0.8

0

(b)

2.5

+

-0.4

Specific capacity (mAh g )

Current (mA)

0.0

Specific capacity (mAh g )

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Figure 7

250

250

(a)

200 150 100 50 0

(b)

200 -Z? (ohm)

-Z? (ohm)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Nano


0

100

200

300 400 Z' (ohm)

(c)

500

600

Rf

150 100 50 0


0

100

200

Rct

Rs

CPE1

CPE2


300 400 Z' (ohm)

Zw

500

600

ACS Nano

Figure 8 0.4

3.0 Potential vs. Na/Na (V)

(a)

0.2

+

-0.2 -0.4 -0.6 -0.8 -1.0

1st 2nd 3rd

-1.2 0.0

0.5 1.0 1.5 2.0 2.5 + Potential vs. Na/Na (V)

1000 (c) 800

-1

0. 2 A g

0.4 A g

400

-1

0.8 A g

1.0 1st 2nd 3rd

0.5

200

0


10


(e)

100

200 0

200 400 600 800 1000 1200 -1 Specific capacity (mAh g )

-1

1.6 A g

0

1.5

300

-1 0. 1 A g

-1

2.0

400

0. 1 A g

600

2.5

0


-1

(b)

0.0

3.0

-Z? (ohm)

Current (mA)

0.0


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 44 of 47

50

60

H-FeP@C@GR in SIBs H-FeP@C@GR in LIBs

0

100

200 Z' (ohm)

300

400

1000 ( d) 800 600

H-FeP@C@GR

400 Charge 200

Discharge

Current density : 0.1 A g-1

0 0

25

50

75


175


200

225

250

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Figure 9


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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Figure 10

States of discharge Atomic ratio(Na:Fe:P)

Element

Na

Fe

Atomic %

9.38

10.83

P

Element

11.49

Atomic %

a 0.86:1:1.06

b

c

2.09:1:1.07

2.97:1:1.07

Na

Fe

P

11.12

5.31

5.71

Element

Na

Fe

P

Atomic %

3.48

1.17

1.25


Page 47 of 47

Figure 11 20

(a)

Li+

Relative Energy (eV)

15

Li P Fe

10

5

? E=11.133eV

0

0.0

30

0.2

0.4 0.6 Diffusion Path

0.8

(b)

1.0

Na+

25 Relative Energy (eV)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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20 15

Na

10

P Fe

? E=16.711eV

5 0 0.0

0.2

0.4 0.6 Diffusion Path

0.8

1.0


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