Molecular-level heterostructures assembled from layered black phosphorene and Ti3C2 MXene as superior anodes for high-performance sodium ion batteries

Nano Energy 65 (2019) 104037

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Molecular-level heterostructures assembled from layered black phosphorene and Ti3C2 MXene as superior anodes for high-performance sodium ion batteries

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Ruizheng Zhaoa, Zhao Qiana, Zhongyuan Liub, Danyang Zhaoa, Xiaobin Huia, Guanzhong Jianga, ⁎⁎ ⁎ Chengxiang Wanga, , Longwei Yina, a Key Laboratory for Liquid-Solid Structural Evolution and Processing of Materials, Ministry of Education, School of Materials Science and Engineering, Shandong University, Jinan, 250061, PR China b State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinghuangdao, PR China

A R T I C LE I N FO

A B S T R A C T

Keywords: Layered black phosphorene Ti3C2 Heterostructures Sodium-ion batteries Anode

Phosphorus, as one of the most promising anodes for sodium-ion batteries, its electrochemical performance improvement seriously suffers from insulated characteristics and poorly structural stability. Herein, we report an elaborately designed strategy to rationally synthesize molecular-level PDDA-BP/Ti3C2 nanosheet heterostructures taking advantages of high theoretical capacity of black phosphorene (BP) and high electronic conductivity, abundant functional groups of Ti3C2. Due to the face-to-face contact of both components, the parallel 2D interlayer spacing provides effective charge transfer and diffusion channels. DFT calculations show that strong interactions between BP and Ti3C2 could efficiently lower binding energy to facilitate the sodiation process. More importantly, surface functional groups of –F, –O and –OH in Ti3C2 play important roles to immobilize BP and act as more synergistic adsorption sites to accelerate sodiation. The PDDA-BP/Ti3C2 electrode displays extremely structure-stability confining monodispersed BP nanoparticle within Ti3C2 to buffer volume expansion and prevent aggregation of BP, exhibiting an ultrahigh reversible capacity of 1112 mA h g-1 at 500thcycle at 0.1 A g-1 and ultralong cycling stability of 658 mA h g-1 with only 0.05% degradation per cycle within 2000 cycles at 1.0 A g-1. The related soidation mechanism and effects of functional groups of Ti3C2 on sodiation/ desodiation redox reaction are deeply investigated.

1. Introduction Phosphorus (P) is one of the most promising candidates as anode materials for sodium-ion batteries (SIBs) owing to its high theoretical capacity of 2596 mA h g-1 by forming Na3P alloy and low electrode potential, which almost exceeds any other presently available SIBs anode materials [1–3]. However, P, as alloying-type SIBs anode materials, its electrochemical performance improvement seriously suffers from insulated characteristics and large volume expansion during electrochemical redox reaction process. Coupling P with conductive materials such as carbon is considered as an effective strategy to overcome the obstacle of low electrochemical redox kinetics and poor structural stability [4–7]. Since the pioneering works of Yang's group and Chou's group [4,5], P has been widely investigated as SIB anodes based on the strategy combining P with conductive matrix [6–8]. For

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example, MOF-derived carbon with 3D connected porous structure not only can effectively immobilize red P as SIB anodes, facilitate the rapid diffusion of organic electrolyte ions and improve the conductivity of the encapsulated red P, but also the porous matrix structure can buffer the volume change, exhibiting enhanced cycling stability with very low capacity fading of 0.02% per cycle, and a moderate capacity of 450 mA h g-1 at 1 A g-1 [8]. Recent reports show that chemical bonding between P and carbon matrix with a strengthened contact can greatly improve the electrochemical cycle stability and reduce the potential. For example, Song et al. prepared a long durable P-CNT hybrid by making P–O–C bonds to strengthen contact between P and CNTs, greatly improving the conductivity and helping to endure large volume expansion, maintaining 91% capacity after 100 cycles [6]. Although these routes are proved very effective to improve the energy storage performance, it is still challengeable to improve cycle stability and

Corresponding author. Corresponding author. E-mail addresses: [email protected] (C. Wang), [email protected] (L. Yin).

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https://doi.org/10.1016/j.nanoen.2019.104037 Received 9 July 2019; Received in revised form 12 August 2019; Accepted 17 August 2019 Available online 19 August 2019 2211-2855/ © 2019 Elsevier Ltd. All rights reserved.

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Fig. 1. Schematic illustration for the preparation process of PDDA-BP/Ti3C2 nanosheet heterostructures. The water molecules between the layers are not shown here for clarity.

SIBs. This robust nanosheet heterostructures are expected to constraint the serious volume changes and prevent aggregation and loss of BP resulting from phase transformations and structural collapse during the fast charge/discharge process, which is important to further enhance the durability of reversible capacity and long-term cycling stability [28]. In this paper, we report an elaborately electrostatic attraction strategy to rationally design molecular-level (PDDA-BP/Ti3C2) heterostructures based on layered black phosphorene with high theoretical capacity and Ti3C2 with high electronic conductivity, flexible plasticity, abundant functional groups. The modification of cationic polymer poly (diallyl dimethyl ammoniumchloride) (PDDA) on BP not only improves electronegativity of BP nanosheets, but also passivates them to prevent from oxidization in water, enhancing the stability and dispersity [22,23]. Due to the face-to-face contact of both components, the parallel 2D interlayer spacing provides suitable diffusion channels for ions, and thus greatly enhancing the electrochemical performance. Consequently, the PDDA-BP/Ti3C2 heterostructures electrode exhibits an ultrahigh reversible capacity of 1112 mA h g-1 at 0.1 A g-1 and ultralong cycling stability of 658 mA h g-1 with only 0.05% degradation per cycle within 2000 cycles at 1.0 A g-1. It reveals that PDDA-BP/Ti3C2 heterostructures electrode displays an extremely high structure stability that maintains initial monodispersed state of BP nanoparticles within Ti3C2 after long cycles. Furthermore, it demonstrates strong interactions between BP and Ti3C2, which could efficiently lower the binding energy and facilitate the sodiation process as certified by DFT calculations. The related mechanism at atomic scale and effects of functional groups of Ti3C2 on the sodiation/desodiation redox reaction are deeply investigated.

structural stability due to pulverization of P. MXene, as a rising star of 2D materials with a formula of Mn+1XnTx, M = Ti, Nb, Mo, V, etc., X = C, N, and T = O, F, OH surface functional groups, exhibits excellent conductivity and controllable flexible interlayer spacing, which can be conducive to enhance charge transport kinetics and electrochemical redox kinetics [9–15]. It is revealed that monolayer Ti3C2 nanosheet displays a high conductivity of ~ 6.76 × 105 S m-1 [9]. Meanwhile, the diffusion barriers of Na+ on –F, –O, and –OH terminated Ti3C2 are calculated to be about 0.19, 0.2 and 0.013 eV, respectively, demonstrating superior Na+ diffusion kinetics due to the presence of abundant functional groups on Ti3C2 surface [10]. It is shown that functional groups on Ti3C2 surface are electrochemically redoxable, superior to carbon materials as SIBs anodes [16–19]. Wu et al. reported that few-layer Ti3C2 nanosheets as SIBs anodes deliver a reversible capacity of 267 mA h g−1 due to facilitated fast electron transport and Na+ diffusion kinetics [17]. A layer-by-layer stacking of layered structure allows us to re-stack nanosheets with different functional characteristics into molecule-level heterostructures, which is proved definitely to make full use of both characteristics of components of the heterstructures for superior electrochemical performance compared with mechanically mixed layered nanosheets [20]. Layered black phosphorus (BP), as the most stable allotrope of P, can be easily exfoliated to phosphorene nanosheets with few-layer [21]. However, few-layer BP nanosheets are usually negatively charged in water, and would suffer from rapid degradation due to the reaction between phosphorus and oxygen. Surface medication and functionalization of BP is necessary to improve the structural stability and couple with other functional components [22,23]. On the other hand, Ti3C2 nanosheets display high conductivity, plastic characteristics and abundant functional groups, would provide a unique 2D confined environment for coupling with BP nanosheets, which is effective to buffer the volume expansion and prevent the aggregation and pulverization of BP during the charge/discharging process [23–27]. Taking consideration the high conductivity and plastic characteristics of Ti3C2, the high theoretical capacity of BP, also typically layered structure of both BP and Ti3C2, a unique molecular-level heterostructure coupling Ti3C2 with BP is strongly expected, which probably has higher specific capacity and longer lifespan, but still keeps unexplored. It is greatly challengeable and important to rationally design molecular-scale BP/Ti3C2 MXene heterostructures for high-performance

2. Results and discussion 2.1. Structure characterizations of PDDA-BP/Ti3C2 nanosheet heterostructures The PDDA-BP/Ti3C2 heterostructures are prepared through a flocculation process of exfoliated Ti3C2 (ex-Ti3C2) nanosheets and PDDAmodified exfoliated BP (PDDA-BP) nanosheets with oppositely charged states (Fig. 1). Owing to the surface functional groups, the ex-Ti3C2 nanosheet is negatively charged with a zeta potential of −49.6 mV. 2

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Fig. 2. Structure characterization of PDDA-BP/Ti3C2 nanosheet heterostructures. (a) Cross-sectional SEM image. Inset shows PDDA-BP/Ti3C2 flocculation. (b) TEM image. Inset shows photograph of PDDA-BP/Ti3C2 film is flexible and freestanding. (c) ED pattern of PDDA-BP/Ti3C2 flocculation. (d-e) AFM image and thickness profile of monolayer PDDA-BP/Ti3C2 nanosheet heterostructures. (f) TGA curves of ex-Ti3C2 and PDDA-BP/Ti3C2 nanosheet heterostructures from room temperature to 600 °C in N2 atmosphere. (g) XRD pattern of PDDA-BP/Ti3C2. (h) Raman spectra of PDDA-BP, ex-Ti3C2 and PDDA-BP/Ti3C2 nanosheet heterostructures. (i) HAADF-STEM image of the PDDA-BP/Ti3C2 nanosheet heterostructures and the corresponding elemental mapping images of Ti, C, O, F and P elements, respectively.

There are additional sharp peaks at 16.7°, 26.4°, 34.1°, 34.9° and 52.1°, corresponding to (020), (021), (040), (111) and (060) diffraction peaks of BP [31]. Raman spectra (Fig. 2h) are used to investigate the microstructures and interactions between PDDA-BP and ex-Ti3C2 nanosheets. For PDDA-BP sample, three peaks at 464, 433, and 352 cm−1 are assigned to Ag2, B2g and Ag1 vibrational modes of P–P bonds, respectively [26]. For ex-Ti3C2 sample, three peaks around 203, 400 and 622 cm-1 correspond to the vibrations from titanium carbide, which is consistent with previous report [11,13]. Besides, two obvious peaks at 1320 and 1578 cm-1 could be attributed to D band and G band of the layered structures, respectively. The ratio of intensity between ID and IG of exTi3C2 is 0.52, suggesting a high graphitization degree. For PDDA-BP/ Ti3C2 heterostructures, I and III peaks broaden and downshift, while peak II disappears, and Ag2, B2g and Ag1 modes blue-shift due to the breakup of partial P–P and Ti–C bonds for formation the P–O–Ti bonds at 1039 cm-1. Meanwhile, the G band of PDDA-BP/Ti3C2 heterostructures shifts to a lower wavenumber (1575 cm−1) because of the πp* conjugation compared to that of ex-Ti3C2, demonstrating the formation of P–O–Ti bonds [32]. Moreover, the ratio of 0.87 for intensity between ID and IG of PDDA-BP/Ti3C2 heterostructures, suggests a reduced graphitization degree, further confirming possible interactions within the heterostructures. A typical high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) image and the corresponding elemental mapping images of the PDDA-BP/Ti3C2 nanosheet heterostructures (Fig. 2i) show a uniform distribution of Ti, C, O, F and P elements, indicating that PDDA-BP and ex-Ti3C2 are uniformly restacked through the flocculation process. For comparison, mechanically mixed BP-Ti3C2 hybrid with 25.9 wt% BP content (Fig. S5) shows segregated diffraction peaks and restacked in a random

However, the zeta potential of PDDA-BP nanosheet is measured to be +44.3 mV, indicating an electrostatic interaction between PDDA-BP and ex-Ti3C2 nanosheets. Monodispersed nanosheet suspension turns to precipitated layer-structured products (Fig. 2a, inset), which could be easily made into a flexible film (Fig. 2b, inset) through vacuum filtration. Cross-section SEM observation (Fig. 2a) reveals a lamellar structure as expected layer-by-layer stacking of PDDA-BP and ex-Ti3C2 nanosheets. TEM image (Fig. 2b) presents that the ultrathin PDDA-BP nanosheets are uniformly distributed throughout the transparent exTi3C2 nanosheets with crumples. Electron diffraction (ED) pattern (Fig. 2c) shows the in-plane diffraction rings and discrete sharp diffraction spots are indexed to ex-Ti3C2 and PDDA-BP nanosheets, respectively, which can be further confirmed in Figs. S1–4 [26,29]. A typical stacking of both layered nanosheets is revealed by AFM images in Fig. 2d–e. PDDA-BP nanosheet with a thickness of 1.56 nm is superimposed on ex-Ti3C2 nanosheet (1.39 nm in thickness) face to face, and layer by layer due to electrostatic attraction force. In this regard, according to area matching hypothesis (1:1) (Table S1), we could optimize the mass ratio of ex-Ti3C2/PDDA-BP with 1.70 to get an ideal heterostructure [20]. The TGA profile of PDDA-BP/Ti3C2 composite (Fig. 2f) shows a 38.4 wt% mass reduction at 370 °C due to the sublimation of BP [30], assuming the almost same content of water and other adsorbed species in both PDDA-BP/Ti3C2 composite and ex-Ti3C2 materials. Therefore, the real mass ratio of ex-Ti3C2 to PDDA-BP is estimated to be 1.60, which is close to the calculated value. The lamellar structure of the heterostructures is also revealed by XRD pattern (Fig. 2g). XRD diffraction peaks at 4.0°, 7.9°, 20.2°, 24.3° and 28.1° could be ascribed to (001), (002), (005), (006) and (007) planes of a layered structure of Ti3C2, where (001) peak with the basal spacing of 2.2 nm is very close to basal spacing of the configuration (2.3 nm). 3

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Fig. 3. Electrochemical performance. (a) CV profiles at 0.1 mV s-1, (b) Initial five galvanostatic charge/discharge profiles at 0.1 A g-1, (c) Rate performance and (d) Long cycling performance at 1.0 A g-1 for 2000 cycles for PDDA-BP/Ti3C2 heterostructures electrode. (e) A red light-emitting diode (LED) could easily be powered by two SIBs based on PDDA-BP/Ti3C2 heterostructures electrode assembled in series. (f) The 5th charge/discharge profiles at 0.1 A g-1, (g) Cycling stability at 0.1 A g-1 and (h) Rate performance at 0.1, 0.2, 0.5, 1.0, 2.0 A g-1 for different electrodes, respectively.

electrolyte interface (SEI) film (Fig. S10). When U < 0.25 V, the cathodic current undergoes an intense growth, which is mainly due to the alloying process of P with Na+. These processes agree well with the plateaus of the first cycle in galvanostatic discharge profile (Fig. 3b). In the following anodic scan, a small peak at 0.67 V corresponds to the dealloying process of NaxP phases [30]. Besides, a weak reduction peak at 1.58 V, two weak oxidation peaks at 1.61 V and 2.17 V, are attributed to the Na+ adsorption/desorption between the heterostructure layers, respectively [17]. In the next cycles, the CV profiles almost overlap with each other except for a little decreasing cathodic current between 0.01 and 0.35 V, indicating the energy storage process is going to a stable state, i.e., a reversible alloying/de-alloying of Na+ with P. Besides, there is a broad cathodic peak at 0.3 V–1.0 V. According to Marbella et al.‘s calculation, it should be attributed to a gradually alloying process of Na3P11→Na3P7→NaP. When the potential of the electrode is discharged to < 0.3 V, a rapidly strengthen current suggests an active and high-capacity reaction happens, which probably corresponds to NaP→Na5P4 and finally Na5P4→Na3P [40].

manner because both kinds of nanosheets are negatively charged [30]. Notably, strong interactions of PDDA-BP/Ti3C2 nanosheet heterostructures film with superior wettability are further revealed by FTIR, XPS and contact angle measurements (Figs. S6–8) [11–13,30–39]. Furthermore, the PDDA-BP/Ti3C2 nanosheet heterostructures (Fig. S9 and Table S2) reveal mitigative restacking and increased specific surface area of 54.25 m2 g-1 with a pore size distribution around 7.44 nm, much higher than that of other materials, which could expose more accessible active sites to enhancing performance.

2.2. Electrochemical measurements and performances comparison Fig. 3a depicts cyclic voltammetry (CV) curves of PDDA-BP/Ti3C2 heterostructures electrode during the initial five cycles at a scan rate of 0.1 mV s−1. For the CV profiles of PDDA-BP/Ti3C2 heterostructures electrode, there are two obvious cathodic peaks for the first cycle at 1.0 and 0.45 V, which disappear in next cycles, meaning these two peaks correspond to irreversible process, i.e., the formation of solid 4

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percentage of BP = 351.8 × 61.6% + 2596 × 38.4% = 1213.5 mA h g−1). While PDDA-BP, ex-Ti3C2 and mechanically mixed BP-Ti3C2 electrodes only exhibit reversible capacities of 43, 185 and 465 mA h g−1, revealing excellent cycling stability, which is superior to those of other previously reported Ti3C2-based and BP-based anodes(Table S5). [16-19,26,30,37,43-47,49-53]. It evidently indicates that 2D PDDA-BP/ Ti3C2 heterostructures electrode intrinsically has a better ability to accommodate and survive large stresses and strains, and decrease energy barriers of ions movement during sodiation/desodiation process. Besides, the good electron conductivity and facile ion diffusion paths are simultaneously achieved in this heterostructures, thus offering the high reversible capacity and cycle stability. In addition, in order to further investigate the influence of BP content on the electrochemical performance, PDDA-BP/Ti3C2 heterostructures with three different mass ratios of Ti3C2/BP are prepared and measured under identical test conditions. Figs. S14a–b depicts SEM images PDDA-BP/Ti3C2-1, PDDA-BP/Ti3C2-2 samples. As shown in Fig. S14c, PDDA-BP/Ti3C2-2 heterostructures electrode (the mass ratio of 1.7 for Ti3C2/BP) shows much better electrochemical performance than PDDA-BP/Ti3C2-1 and PDDA-BP/Ti3C2-3 electrodes. The enhanced electrochemical performance of PDDA-BP/Ti3C2-2 heterostructures electrode could be attributed to the synergistic interactions between the high theoretical capacity of BP and the high conductivity of Ti3C2. Specifically, the initial discharge capacity increases with BP content, while the initial coulombic efficiency decreases contrarily with BP content increasing. For PDDA-BP/Ti3C2-1 electrode with higher BP content, although shows higher initial discharge capacity, it displays poorer cycle stability than PDDA-BP/Ti3C2-2 and PDDA-BP/Ti3C2-3 electrodes. The low cycle stability and electrochemical performance of PDDA-BP/Ti3C2-1 electrode should be attributed to poor charge transfer kinetics from the low electronic conductivity, the structural instability from BP aggregation and large volume changes from charge/ discharge process. For PDDA-BP/Ti3C2-3 electrode with relatively lower content of BP, it shows better cycle stability, although it shows lower capacity than PDDA-BP/Ti3C2-2 electrode. The rate performance of PDDA-BP/Ti3C2 heterostructures electrode (Fig. 3h) with high reversible capacities of 1767, 1059, 708 and 559 mA h g−1 is obtained at 0.1, 0.2, 0.5, and 1.0 A g−1, respectively. Even at 2.0 A g−1, a reversible capacity of 464 mA h g−1 can be maintained. Importantly, as the current density decreases to 0.1 A g−1, the reversible capacity is able to return to more than 1210 mA h g−1, showing an excellent rate capacity. In contrast, PDDA-BP, ex-Ti3C2 and mechanically mixed BP-Ti3C2 electrodes only show reversible capacities of 19, 80 and 264 mA h g−1 at 2.0 A g-1 after 50 cycles, obviously worse than that of PDDA-BP/Ti3C2 heterostructures electrode. It reveals that 2D molecular-level PDDA-BP/Ti3C2 heterostructures could provide multiple accessible active sites, further accelerate electrons/ions transport and facilitate access of electrolyte ions to the electrode, thus leading to higher specific capacity. This heterostructures are valid in achieving a synergistic effect and therefore very promising in developing next-generation high-performance sodium ion batteries.

Fig. 3b reveals the galvanostatic charge/discharge profiles of PDDABP/Ti3C2 heterostructures electrode for the first five cycles at a current density of 0.1 A g-1. It is noted that all the capacities in this work are calculated based on the mass of the total electrode (BP + Ti3C2). Initial discharge and charge capacities of 2588 and 1780 mA h g−1 are obtained, respectively, corresponding to a Coulombic efficiency (CE) of 68.7%. The irreversible capacity loss can be mainly attributed to the inevitable formation of SEI film, the undesirable irreversible reaction and electrolyte decomposition, which is also revealed in CV profiles [17,30]. As cycle number increases, the capacity gradually becomes stable, with CE up to 98.1% at the fifth cycle, much larger than that of the first cycle. Except for the first cycle, all the charge/discharge processes exhibit the similar profiles, corresponding to reversible sodiation/desodiation process, which can be roughly expressed by Ref. [40]: P + xNa++ x e- ↔ NaxP (0 ≤ x ≤ 3)

(1)

Ti3C2Tx + y Na + y e ↔ Ti3C2TxNay.

(2)

+

-

The long potential plateau blow 0.22 V (the fifth cycle) corresponds to the formation of crystalline Na3P. The voltage-profile of desodiation for the first five cycles exhibit monotonically increasing from the 0 V–3 V, which can be ascribed to the combination of multiple charge storage mechanisms with stepwise Na+ de-alloying from the fully charged Na3P phase to form the Na2P, NaP and NaP7 intermediates based on the Na/P phase diagram and calculation [40–42], as well as Na+ extraction out of the robust Ti3C2 matrix. These results agree well with the CV curves. Because BP nanosheets are sandwiched between Ti3C2 nanosheets, it is conducive to enhance the electron conductivity and ion transport kinetics (Fig. S11). The PDDA-BP/Ti3C2 heterostructures electrode (Fig. 3c) exhibits a high capacity of 1773 mA h g-1 at 0.1 A g-1. At higher current densities of 0.2, 0.5, 1.0, 2.0 A g-1, it can even retain capacities of 1062, 707, 560 and 461 mA h g-1, respectively, revealing much higher reversible capacity than other electrodes (Fig. S12) at the same current density. This sandwiched molecular heterostructure exhibits excellent cycle stability by confining BP between Ti3C2 “nanoreactors”. After 2000 cycles, it still retains a reversible capacity of 658 mA h g−1, with a capacity decay as low as 0.05% per cycle (Fig. 3d), much superior to those of other previously reported Ti3C2 MXene-based anodes (Table S3) [17,43–47]. Specially, a red lightemitting diode (LED) could easily be lighted by two PDDA-BP/Ti3C2 heterostructures electrode devices assembled in series shown in Fig. 3e, efficiently discharging for more than 6 h after charging for only 60 s to reach 3.0 V, confirming feasible and potential application of the device. Fig. 3f shows the fifth Galvano static charge/discharge profiles of different electrodes at 0.1 A g-1. The PDDA-BP, ex-Ti3C2, mechanically mixed BP-Ti3C2 and PDDA-BP/Ti3C2 heterostructures electrodes exhibit reversible capacities of 134, 228, 855 and 1635 mA h g−1 after five charge/discharge cycles. Since the mass ratio of ex-Ti3C2 to PDDA-BP is 1.70 in this heterostructures and the content of ex-BP is 25.9 wt% in mechanically mixed BP-Ti3C2 electrode, the capacity contributions of ex-Ti3C2 electrode to the heterostructures and mechanically mixed electrodes are estimated to be 144 and 169 mA h g−1, respectively. Therefore, the capacity of PDDA-BP electrode is 1491 mA h g−1 in this heterostructures, much higher than that of pure PDDA-BP (134 mA h g−1) and mechanically mixed BP-Ti3C2 electrodes (686 mA h g−1) (Table S4). It suggests that the heterostructures are more conducive to make full use of active BP materials than previously reported results. The cycling performance of different electrodes at 0.1 A g-1 as depicted in Fig. 3g undergo serious capacity degradation initially that are attributed to formation of SEI film, which could be confirmed by ex situ XPS (Fig. S13) [48]. The PDDA-BP/Ti3C2 heterostructures electrode still maintains an excellent reversible capacity of 1112 mA h g−1 with a CE of 99.3% after 500 cycles, which is close to the theoretical capacity contribution value (1213.5 mA h g−1) for the PDDA-BP/Ti3C2 heterostructures electrode with a BP content of 38.4 wt % (Ctheoretical = CTi3C2 × mass percentage of Ti3C2 + CBP × mass

2.3. Microstructure evolution of PDDA-BP/Ti3C2 heterostructures electrode Fig. 4a–d shows TEM images of the PDDA-BP/Ti3C2 heterostructures electrode at different stages. Fig. 4a gives the as-prepared one like that in Fig. 2b. As first discharged to 0.01 V (Fig. 4b), nanoparticles with several nanometers are formed uniformly on the matrix. The diffraction rings of the ED pattern (Fig. 4b, inset) correspond to (110) and (200) planes of Na3P, (100) and (210) planes of reflection Ti3C2. As charged to 3.0 V (Fig. 3c), the nanoparticles with a decreased particle size due to the alloying/de-alloying process do not disappear to return to the original structure as that in Fig. 4a, indicating an irreversible process. While the overall morphology is preserved by Ti3C2 as a matrix, as suggested by its persisting (100) and (210) reflection in ED pattern (Fig. 4c, inset). According to the reaction mechanism, an 5

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Fig. 4. Structure evolution of PDDA-BP/Ti3C2 heterostructures electrode. (a) TEM image of as prepared PDDA-BP/Ti3C2. (b) and (e) Ex situ TEM images of PDDABP/Ti3C2 first discharged to 0.01 V. (c) and (f) Ex situ TEM images of PDDA-BP/Ti3C2 first charged to 3.0 V. (d) and (g) Ex situ TEM images of PDDA-BP/Ti3C2 charged to 3.0 V after 500 cycles. In set in (b-d) shows ED patterns of PDDA-BP/Ti3C2: (b) first discharged to 0.01 V, (c) first charged to 3.0 V and (d) charged to 3.0 V after 500 cycles. T stands for Ti3C2, N is Na3P. (h) Ex situ Raman spectra and (i) Ex situ XPS spectra of PDDA-BP/Ti3C2 heterostructures electrode at different sodiation/desodiation states. The green region is related to the oxidized P with different valence state.

of the PDDA-BP/Ti3C2 heterostructures anode. Along with the cycle increases, the diameter decreases a little probably because the longtime pulverization. Furthermore, the microstructure evolution can also be well reflected by ex situ XRD (Fig. S15) [54]. We further carry out ex situ Raman and ex situ XPS analysis for the PDDA-BP/Ti3C2 heterostructures electrode to investigate structure changes in charge/discharge process. For as-prepared PDDA-BP/Ti3C2 heterostructures electrode (Fig. 4h), three peaks located at 464, 433, and 352 cm−1 are assigned to Ag2, B2g and Ag1 vibrational modes of P–P bonds, respectively [26]. Two obvious peaks at 1320 and 1578 cm-1 could be attributed to D band and G band of layered structured Ti3C2, respectively [11]. The ratio of intensity between ID and IG of ex-Ti3C2 is 0.79, suggesting a high graphitization degree, which is important for getting a good conductivity. When discharged to 0.01 V, the Ag2, B2g and Ag1 signatures disappear, corresponding to the fully breaking of P–P bonds after formation of Na3P phase. After completely charged to 3.0 V, Ag2, B2g and Ag1 of P–P bonds appear again but not as strong as that of as-prepared electrode, which is probably caused by different

evolution in structure and morphology for Ti3C2 would not take place, and nanoparticles should correspond to P or its compounds [26,30,40]. In discharge process, P–P bonds gradually break with the intercalation of Na+, and finally form crystalline Na3P. However, if Na+ is extracted in charge process, P species could not reorganize to form original BP crystal, leaving amorphous BP nanoparticles. Furthermore, the nanoparticles display crystal structure characteristics at 0.01 V (Fig. 4e) with the lattice spacing of 0.25 nm, corresponding to (110) plane of Na3P, and transform to amorphous ones (Fig. 4f) at 3.0 V in HRTEM images, which is well in agreement with the previous reports [26,30]. So far, there are seldom reports on monodispersed P nanoparticles on conductive matrix, especially for that with long-time stability. Herein, it is proved that BP nanosheet could convert into well-dispersed BP nanoparticles sandwiched between Ti3C2 nanosheet through first sodiation/ desodiation process. More importantly, these nanoparticles with a decreased particle size could keep the monodispersed state even after 500 cycles preserved by Ti3C2 with its in-plane (100) and (110) reflections as shown by Fig. 4d–g, further demonstrating excellent cycling stability

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More importantly, two different binding energies of -2.17 eV for BP-Na system and -2.83 eV for PDDA-BP-Na/Ti3C2 system are obtained (Fig. 5j, left panel). It means that Na prefers to combine with the BP nanoparticle anchored on Ti3C2 matrix, which evidently demonstrates the strong interactions between BP and Ti3C2. The difference charge densities after Na adsorption on BP nanoparticle or PDDA-BP/Ti3C2 heterostructure surfaces are shown in Fig. 5c–d. Obviously, the addition of Ti3C2 changes the charge density profiles around Na, P and Ti. Through the Bader charge calculations, it is found that the changes of charge states of most P atoms are more than -0.5, even up to -0.9 (for P2 and P6), except for P4 in Table S6. It suggests Ti3C2 nanosheet evidently changes the sodiation state of BP nanoparticles, makes it easier to get electrons from Na and facilitates the sodiation process. We further investigate the relaxation process when adding Na on top of surface functional groups, such as –F, –O and –OH. Their configurations after relaxation are shown in Fig. 5g (on top of –F) and Fig. 5h (on top of –O). In the former case of –F, Na bonds with two P atoms (2.87 and 2.81 Å) and one F atom (2.17 Å). In the latter configuration for –O, Na–O (2.29 Å), Na–F (2.28 Å) and Na–P (2.92 Å) bonds are formed. An interesting phenomenon here is that Na favors to combine with more than one atom to lower the binding energy as much as possible. The binding energy of the former case (-2.27 eV) is not as low as the latter (-2.45 eV) (Fig. 5j, right panel). While, when adding another Na on the top of –OH, the relaxation result indicates that Na could not be adsorbed on –OH surface functional group, but moves and combines with two P atoms (Na–P bonds of 3.19 and 3.08 Å, Fig. 5i), probably because there is a negative electron cloud which screens most of the attractive forces between the Na and the H atoms on the surface layer of Ti3C2. This case even lowers Na binding energy of -3.15 eV (Fig. 4j, right panel) than those in Fig. 5g–h, which means Na prefers to combine with P rather than to be adsorbed on the surface functional groups. The calculated differences in charge densities after another Na adsorption on the top of various surface functional groups are shown in Fig. S17. From the Bader charge analysis, it is calculated that there are 0.87, 0.86 and 0.73 electrons transferred from Na to the –F, –O and –OH surface functional groups of those heterostructures, leading to the decreased valence of F, O and P atoms and the increase of electron accumulation (Table S6). Based on the above analysis, the intercalated Na+ would combine with BP nanoparticles initially, then undergo a mixed adsorption on –F, –O and -P, in which case more bonds are favorable to lower the binding energy, further promoting fast ions diffusion and charge transfer kinetics.

state of P–P bonds between crystalline BP nanosheet and re-formed amorphous BP nanoparticles. Three small peaks around 226, 409 and 615 cm-1 correspond to the vibrations from titanium carbide, which is consistent with previous report [11]. Notably, some peaks at a range of 890-1200 cm-1 are possibly caused by the strong interactions between the BP nanoparticles and Ti3C2 nanosheets through P–O and P–O–Ti bonds during the electrochemical process, agreeing well with the results of ex situ XPS spectra [32,55–57]. Furthermore, the ratio of intensity between ID and IG related to disordered degree of carbon increases from 0.79 to 0.99 and then decreases to 0.96 for the asprepared, fully discharged and fully charged electrodes, respectively, which further confirms the structural sustainability and reversibility of the PDDA-BP/Ti3C2 electrode upon sodiation/desodiation process. For the typical XPS spectra, the as-prepared PDDA-BP/Ti3C2 heterostructures electrode (Fig. 4i) displays two bands at 129.3 and 130.0 eV, corresponding to 2p3/2 and 2p1/2 of BP nanosheet, respectively [35,36]. Furthermore, two peaks at 133.6 and 132.6 eV are ascribed to weak oxidized P–F and P–O (PxOy) bonds, suggesting that the surface of BP nanosheet is inevitably oxidized during the preparation process [58]. After discharged to 0.01 V, the peaks at 129.3 and 130.0 eV of P0 state disappear, meanwhile P–O and P–F bonds shift to higher binding energy, corresponding to the fully sodiation of BP. The former is in accordance with Raman spectra. The P–O and P–F bonds are helpful to immobilize the P species and prevent the aggregation and pulverization. An interesting point is that after charged to 3.0 V, the amorphous BP nanoparticles do not fully recover to P0 state with 129.3 and 130.0 eV, but still exhibit considerable P–F and P–O signals. It means strong interactions between the BP nanoparticles and Ti3C2 nanosheet take place through P–F and P–O bonds, rather than as isolated active materials. The interactions of BP nanoparticles with Ti3C2 nanosheet would be further investigated through DFT calculations in next section. Based on the discussions above, a schematic illustration (Fig. S16) of possible 2D-confined electrochemical reactions is proposed. The interlayer distance changes when Na+ inserts/extracts between robust Ti3C2 layers, while BP nanosheet transforms to Na3P nanoparticles and further converts into amorphous BP nanoparticles during the first sodiation/desodiation process, which are homogeneously dispersed and confined between the Ti3C2 gallery networks through strong interactions of P–F and P–O bonds. Thus, the electrochemical reaction processes are highly reversible from the second cycle, which is advantageous to the excellent cycling stability and rate capacity. 2.4. Sodiation process revealed by DFT calculations

2.5. Electrochemical measurements to reveal the dynamic difference To figure out the possible mechanism of enhanced sodium storage performance for the PDDA-BP/Ti3C2 heterostructures electrode, we further perform DFT calculations to investigate the interactions between Na and BP nanoparticle, and the interactions between BP and Ti3C2 with functional groups. The BP used for calculation is mimicked by a BP nanoparticle of six atoms. For a simple Na-BP system, as a Na atom is absorbed in this system (Fig. 5a), an equilibrium configuration after relaxation process will finally reach a relatively stable state, as shown in Fig. 5b. Five kinds of Na–P bonds with length of 2.78, 2.80, 2.98, 3.28, 3.29 Å are formed, respectively. For the system of PDDA-BP/ Ti3C2 without considering the functional groups on Ti3C2, the initial sixatom BP nanoparticle are absorbed on Ti3C2 matrix (Fig. 5e), an equilibrium configuration after relaxation process can be depicted in Fig. 5f. There form four types of Na–P bonds with lengths of 2.85, 2.95, 2.91, 2.33 Å, respectively. In both cases, the Na–P bonds are very close to those in Na3P with Na–P bond lengths of 2.87, 2.92, 3.24 Å, respectively. This indicates the present DFT calculation results are reasonable. It is noted that (Fig. 5b–f) a Ti3C2 matrix will lead to some significant differences. The BP nanoparticle on Ti3C2 are in favor to move at the PDDA-BP/Ti3C2 interface rather than in a free state. These nanoparticles are separated into two groups a bit far from each other, which may be corresponding to the volume expansion in sodiation process.

Electrochemical impedance spectroscopy (EIS) is conducted to investigate the charge transfer kinetics at the interface and Na+ diffusion ability of different electrodes [59]. As shown in Fig. 6a, the Nyquist plot mainly consists of a depressed semicircle in the high frequency region associated with the charge transfer resistance (Rct) and a straight line in the low-frequency region related to the Na+ diffusion ability defined as Warburg impedance (W0). Rct values for the PDDA-BP, ex-Ti3C2, mechanically mixed BP-Ti3C2 and PDDA-BP/Ti3C2 heterostructures electrodes are 257.8, 66.5, 222.2 and 127.6 Ω, respectively, suggesting that PDDA-BP/Ti3C2 heterostructures electrode exhibits the smaller charge transfer resistance than PDDA-BP and mechanically mixed BP-Ti3C2 electrodes, which reveals that this heterostructures electrode shows improved electronic conductivity and charge transfer ability. According to relevant reports [59,60], the Na+ diffusion ability is related to the phase angle in the low frequency region (below 1 Hz) of the Bode plots, and the smaller the phase angle, the faster the Na+ diffusion ability. As can be seen in Fig. 6b, the PDDA-BP/Ti3C2 heterostructures electrode exhibits the smallest phase angles among these electrodes besides exTi3C2 electrode in the low-frequency region, demonstrating faster ion diffusion kinetics than others in electrolyte. The Na+ diffusion coefficient in the electrode could be obtained by the plots in the low7

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Fig. 5. Sodiation process revealed by DFT calculations. (a), (e) The initial models. (b), (f) The most stable adsorption configuration between Na and BP nanoparticle or PDDA-BP/Ti3C2 heterostructure. (c), (d) The difference charge densities for one Na adsorption on BP nanoparticle and PDDA-BP/Ti3C2 heterostructure. Cyan and yellow regions show the depletion and accumulation of electrons, respectively. The isosurface values are chosen as 0.0005eV/Å3. (g-i) The most stable adsorption configurations for another Na adsorption on top of functional groups in the heterosturctures, such as –F, –O and –OH. (j) The calculated binding energies for Na adsorption on the surface of BP nanoparticle or PDDA-BP/Ti3C2 heterostructures with different surface functional groups. The PDDA and water molecules between the layers are not shown here for clarity.

Where a is a coefficient, b = 1 stands for surface-controlled process completely, while b = 0.5 for diffusion-controlled one completely. Fig. 6d shows the CV profiles of PDDA-BP/Ti3C2 heterostructures electrode. By plotting log i vs log v, b is the corresponding slope as shown in Fig. 6e. For cathodic and anodic current, the b-value at peak potential is 0.82 and 0.87, respectively, indicating that the surface faradaic redox reactions are predominantly capacitive behaviors. The same kinetics analysis based on the CV curves for the PDDA-BP (Fig. S18) and mechanically mixed BP-Ti3C2 electrodes (Fig. S20), where much smaller b-values are obtained beside ex-Ti3C2 (Fig. S19), which further suggests that the ex-Ti3C2 and PDDA-BP/Ti3C2 heterostructures electrode shows improved ion diffusion and charge transfer kinetics than others. We further check b-values for anodic currents at 0.5, 0.75, 1.0, 1.25 and 1.5 V, respectively (Fig. 6f). The b-values does not change much for each electrode in the order ex-Ti3C2 > PDDA-BP/Ti3C2 > BPTi3C2 > PDDA-BP with average b-values of 0.95, 0.82, 0.76, 0.52, respectively. It means mechanically mixed BP-Ti3C2 and PDDA-BP electrodes are much influenced by ion diffusion, while ex-Ti3C2 and PDDABP/Ti3C2 heterostructures electrodes are not. That's the reason why

frequency region based on the following equation [60,61]: |Z | = Rct + Re + σω 2 2

−1/2

(3)

CNa2σ2)

2 4 4

DNa = R T /(2A n F

(4)

+

Where D stands for the Na diffusion coefficient, R represents the gas constant, T is the absolute temperature, A is the surface area, n is the number of electrons per molecule oxidized, F is Faraday's constant, C is the concentration, σ is the Warburg factor, and ω is the frequency. Obviously, Fig. 6c demonstrates that the PDDA-BP/Ti3C2 heterostructures electrode displays the smallest σ among these electrodes besides ex-Ti3C2 electrode, revealing the largest Na+ diffusion coefficient among the electrodes. For CV measurements, the current of all the electrodes consist of diffusion-controlled and surface-controlled (capacitive) parts. It can be expressed as [51]: i = avb, (0.5 ≤ b ≤ 1)

(5)

log i = b log v + log a

(6)

8

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Fig. 6. Electrochemical mesurements to reveal the dynamic difference. (a) Nyquist plots. (b) Bode plots before cycling in the frequency range from 100 KHz to 10 m Hz for different electrodes, respectively. Inset exhibits the equivalent circuit. (c) Extracted |Z| vsω1/2 plot in the low frequency region for different electrodes, respectively. (d) CV curves at various scan rates and (e) Relationship between log (i) vs log (v). (f) Extracted b-values for different electrodes at different potentials. (g) Capacitive-controlled and diffusion-controlled contributions at 1.0 mVs-1. (h) Normalized contribution ratio of capacitive-controlled capacities for PDDA-BP/ Ti3C2 and mechanically mixed BP-Ti3C2 electrodes at different scan rates. (i) Charge transfer mechanism of PDDA-BP/Ti3C2 heterostructures electrode.

efficiently take advantage of both components and more efficient ions diffusion from Ti3C2 surface to BP; (ii) The flexible Ti3C2 with high electron and Na+ conductivity provides effective electron transfer paths and then penetrates to BP through the face-to-face contact, leading to strong interactions, efficiently decreasing the binding energy of Na and BP to facilitate the sodiation process and inducing intrinsic more capacitive behaviors; (iii) The monodispersed BP nanoparticles with small diameter are well thoroughly and chemically bonding stabilized and confined within and between 2D Ti3C2 layers after a longtime cycle, providing one freedom to buffer the volume expansion, further preventing aggregation and loss of BP.

PDDA-BP/Ti3C2 heterostructures electrode displays much larger capacity than mechanically mixed BP-Ti3C2 and PDDA-BP electrodes at the same current density. Actually, the current can be further written as [17,43]: i(v) = k1v + k2v1/2 i(v)/v

1/2

= k1v

1/2

+ k2

(7) (8)

Where k1 and k2 are adjustable values. The first item is capacitivecontrolled contribution (k1v) and the second item is diffusion-controlled contribution (k2v1/2). It enables to resolve the diffusive and capacitive component quantitatively by plotting i/v1/2 vs v1/2. As shown in Fig. 6g, the capacitive capacity reaches 60.7% of the total capacity when PDDA-BP/Ti3C2 electrode is scanned at a rate of 1.0 mV s-1. Similarly, capacities of PDDA-BP/Ti3C2 and mechanically mixed BP-Ti3C2 electrode at different scan rates are resolved (Fig. 6h). The PDDA-BP/ Ti3C2 heterostructures electrode always shows larger capacitive component than mechanically mixed BP-Ti3C2 electrode as the b-values indicates. It is believed that sufficient contact of Ti3C2 and BP in PDDABP/Ti3C2 heterostructures electrode leads to a more efficient ion diffusion from Ti3C2 surface to BP. Apparently, the PDDA-BP/Ti3C2 heterostructures electrode exhibits superior electrochemical performance compared to other electrodes, which is mainly due to the following benefits (Fig. 6i): (i) The molecular-level heterostructure with orderly stacked structure ensures the intimate contact between Ti3C2 and BP to

3. Conclusions We rationally design a novel molecular-level PDDA-BP/Ti3C2 nanosheet heterostructrues. Consequently, the PDDA-BP/Ti3C2 heterostructures electrode exhibits a reversible capacity of 1112 mA h g-1 at 0.1 A g-1, and shows a superior cycling stability with 0.05% degradation per cycle within 2000 cycles at 1.0 A g-1, which is attributed to the unique heterostructures. The molecular-level heterostructure with orderly stacked structure ensures the intimate contact between Ti3C2 and BP to efficiently take advantage of both components. The flexible Ti3C2 with high electron and Na+ conductivity provides effective charge transfer and ion diffusion paths, and then penetrates to BP through the 9

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face-to-face contact. The monodispersed BP nanoparticles with small diameter are well thoroughly and chemically bonding stabilized and confined within and between 2D Ti3C2 layers after a long-time cycle, providing one freedom to buffer the volume expansion, further preventing aggregation and loss of BP. Strong interactions between BP and Ti3C2 could efficiently lower binding energy and enhance interfacial charge transfer to facilitate sodiation process. The intercalated Na+ would combine with BP nanoparticles initially, then undergo a mixed adsorption on –F, –O and –OH, in which case more bonds are favorable to lower the binding energy, further promoting fast ion diffusion and charge transfer kinetics. The related atomic-scale sodiation/desodiation mechanism of molecular-level PDDA-BP/Ti3C2 nanosheet heterostructrues are deeply investigated. The artificial molecular-level heterostructure design and the underlying mechanism should be inspirable to the application of other materials.

heterostructures. The prepared PDDA-BP/Ti3C2 heterostructures with three different mass ratios were denoted as PDDA-BP/Ti3C2-1, PDDABP/Ti3C2-2 and PDDA-BP/Ti3C2-3 by controlling ex-Ti3C2 and PDDABP suspensions with the quality feeding rates of 1.2:1, 1.7:1 and 2.2:1. For comparison, mechanically mixed BP-Ti3C2 hybrid was also prepared by directly mixing ex-Ti3C2 and ex-BP suspensions under the same mass ratio. 4.4. Materials characterization The crystal structures of obtained materials were carried out on a Rigaku D/Max-KA X-ray diffractometer equipped with a Cu Kα source (λ = 1.5406 Å) at a step size of 0.02° with 0.5 s dwelling time. The chemical content was estimated by thermogravimetric analysis (TGA) under a flow of nitrogen atmosphere with a heating rate of 10 °C min-1 from RT to 600 °C. Fourier transform infrared (FTIR) characterization was recorded on Bruker spectrometer (TENSOR 27). Raman spectra were performed on a JY HR800 micro Raman spectrometer, using a 633 nm laser as the excitation source. The exfoliated nanosheets were characterized by atomic force microscope (AFM, FM-Nanoview 1000). The morphologies, microstructures and elemental components were investigated by field emission scanning electron microscopy (FESEM, SU-70), and high-resolution transmission electron microscopy (HRTEM, JEM-2100) equipped with energy dispersive spectrometry (EDS) at an acceleration voltage of 200 kV, respectively. Chemical composition of the samples was further analyzed by high-resolution X-ray photoelectron spectroscopy (XPS) recorded with an ESCALAB 250 instrument equipped with a 150 W Al Kα probe beam. Zeta potential measurements were performed with a Zetasizer Nano ZS apparatus from Malvern Instruments. Contact angles were performed on an OCA20 machine (Data-Physics, Germany).

4. Experimental section 4.1. Preparation of exfoliated Ti3C2 nanosheets Ti3AlC2 powders (> 98 wt % purity) were bought from 11 Technology Co., Ltd. Exfoliated Ti3C2Tx nanosheets were synthesized as described in previous reports with some modifications [13]. Typically, 2.0 g LiF (Alfa Aesar, 98.5%) was dissolved in 9 M HCl (Beijing Chemical Factory, 40 mL) solution. Then 1.0 g Ti3AlC2 powders were slowly added into the mixture solution and maintained at 35 °C for 24 h under stirring. Afterward, the resultant was repeatedly washed with deionized water, centrifuged and decanted until the pH of the supernatant was approximately 6, and the clay-like sediment was obtained. Next, the sediment was dispersed in 200 mL of deionized water and bath sonicated for 60 min under Ar flow, followed by centrifuging for 1 h at 3500 rpm. Finally, the dark green supernatant was filtered through a Jin Teng membrane, whose concentration was calculated by filtering a known volume of the suspension and measuring the weight of the film after dried in air. Herein, the clay-like sediment was exfoliated with H2O labeled as ex-Ti3C2.

4.5. Electrochemical measurements To test electrochemical performance, the active materials, polyvinylidene fluoride binder and acetylene black (Super-P) were firstly mixed with the weight ratio of 70:15:15 and then dissolved in N-methyl-2-pyrrolidinone to form homogenous slurry. After the homogeneous slurry was coated onto Cu foil substrate and dried at 80 °C for 12 h, the electrode was cut into disks with diameters of 12 mm. Therefore, the average mass loading of the electrode is ~1.0-1.2 mg or 0.69-0.83 mg cm−2. Next, the working electrodes were assembled into 2025 coin-cell by using sodium metal as the reference electrode, glass fiber as the separator and 1 M NaClO4 in EC:DMC:EMC (1:1:1 wt%) with 5 wt% FEC additives as the electrolyte. Galvanostatic charging/ discharging (GCD) tests were performed on a LAND CT2001A instrument (Wuhan, China) with a potential range of 0.01-3.0 V at room temperature. Cyclic voltammetry (CV) curves between 0.01 and 3.0 V at different scan rates and electrochemical impedance spectroscopy (EIS) with the frequency range between 100 KHz and 10.0 mHz were tested on an electrochemical work station (PARSTAT2273).

4.2. Preparation of PDDA modification exfoliated BP nanosheets In detail, 2.149 g PDDA 35% solution, 0.363 g tris, 0.173 g NaCl and150 mL H2O were added into a vial and sonicated for 5 min, forming a mixture aqueous. Then, 100 mg bulk BP crystal was added into 10 mL the above mixture aqueous in an agate mortar and then grounded for 10 min. Then, this mixture was transferred to a vial with 140 mL, the above mixture aqueous was bubbled with argon for 10 min to eliminate the dissolved oxygen molecules to prevent bulk BP crystal from oxidation. Next, the mixture solution was carefully sealed and sonicated in an ice-bath for 10 h at a power of 200 W. After settling down for 12 h, the resultant suspension was centrifuged at 12000 rpm for 30 min to remove the residual unexfoliated and multilayered BP nanosheets. Control exfoliated BP nanosheets without PDDA functionalized were also prepared. Herein, the bulk BP crystals exfoliated with and without PDDA functionalized in H2O were labeled as PDDA-BP and ex-BP, respectively.

4.6. Calculations The structural relaxation and binding energy calculations were conducted by using Vienna Ab-initio Simulation Package (VASP) code according to Density functional theory (DFT) [62–65]. The projector augmented wave (PAW) method was performed considering spin-polarized calculations [66]. The scheme of Perdew-Burke-Ernzerhof (PBE) was employed as the exchange correlation functional in generalized gradient approximation (GGA) and a plane wave basis set with an energy cutoff of 520 eV was used to describe the electronic wave function [67]. The K-points mesh of 9 × 3 × 7 was used in pure BP adsorption system. In other adsorption systems, the K-points were tested and used in 3 × 3 × 3. The geometry optimization was performed by minimizing the forces on atoms with a conjugate gradient algorithm and

4.3. Fabrication of flexible PDDA-BP/Ti3C2 nanosheet heterostructures The PDDA-BP/Ti3C2 nanosheet heterostructures were synthesized via an electrostatic attraction self-assembly process. Typically, ex-Ti3C2 and PDDA-BP suspensions with a quality feeding rate of 1.7:1 were controllably dropped into a beaker simultaneously, respectively. After both kinds of solutions were completely added, the mixed solution was sonicated for 60 min and then stirred for another 12 h, filtered using a Jin Teng membrane with copious deionized water, and then freezedrying for 24 h to obtain the freestanding heterostructures film. Herein, the PDDA-BP/ex-Ti3C2 film was labeled as PDDA-BP/Ti3C2 10

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simultaneously minimizing stresses on unit cells without any symmetry constraint. An ionic relaxation convergence criterion of 0.05 eV/Å was tested and employed in the work. The Bader charge calculations were analyzed the charge transfer between Na and BP nanoparticle or PDDABP/Ti3C2 heterostructures with different surface functional groups [68]. The structures and charge density visualization were generated using VESTA. To evaluate interactions between Na and BP nanoparticle or PDDABP/Ti3C2 heterostructures with different surface functional groups, the binding energies had been calculated by the following equation:

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Eads = Ehs+Na - (Ehs + ENa) where Eads means the binding energy, Ehs+Na means the total energy of relaxed adsorption structure, Ehs means the total energy of heterostructures structure with various surface functional groups, and ENa means the total energy of the Na atom. The difference charge density was calculated by the following equation: Δρ(r) = ρhs+Na(r) –ρhs(r) - ρNa(r) Where ρhs+Na(r) means the charge density of adsorption structure, ρhs(r) is the charge density of heterostructures structure, and ρNa(r) represents the charge density of Na atom. Acknowledgements We acknowledge support from the project supported by the Sate Key Program of National Natural Science of China (No.: 51532005), National Natural Science Foundation of China (No.: 51472148, 51272137, 51602181), the Tai Shan Scholar Foundation of Shandong Province, General Financial Grant from the China Postdoctoral Science Foundation (No: 2015M582088), the Fundamental Research Fund of Shandong University and the Young Scholars Program of Shandong University (YSPSDU). Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.nanoen.2019.104037. References [1] J. Ni, L. Li, J. Lu, ACS Energy Lett 3 (2018) 1137–1144. [2] L.D. Ellis, T.D. Hatchard, M.N. Obrovac, J. Electrochem. Soc. 159 (2012) A1801–A1805. [3] P.K. Allan, J.M. Griffin, A. Darwiche, O.J. Borkiewicz, K.M. Wiaderek, K.W. Chapman, A.J. Morris, P.J. Chupas, L. Monconduit, C.P. Grey, J. Am. Chem. Soc. 138 (2016) 2352–2365. [4] J. Qian, X. Wu, Y. Cao, X. Ai, H. Yang, Angew. Chem. Int. Ed. 52 (2013) 4633–4636. [5] W.J. Li, S.L. Chou, J.Z. Wang, H.K. Liu, S.X. Dou, Nano Lett. 13 (2013) 5480–5484. [6] J. Song, Z. Yu, M.L. Gordin, X. Li, H. Peng, D. Wang, ACS Nano 9 (2017) 11933–11941. [7] Y. Hu, B. Li, X. Jiao, C. Zhang, X. Dai, J. Song, Adv. Funct. Mater. 28 (2018) 1801010. [8] W. Li, S. Hu, X. Luo, Z. Li, X. Sun, M. Li, F. Liu, Y. Yu, Adv. Mater. 29 (2017) 1605820. [9] X. Sang, Y. Xie, M.W. Lin, M. Alhabeb, K.L. Van Aken, Y. Gogotsi, P.R. Kent, K. Xiao, R.R. Unocic, ACS Nano 10 (2016) 9193–9200. [10] Y.-X. Yu, J. Phys. Chem. C 120 (2016) 5288–5296. [11] R. Zhao, M. Wang, D. Zhao, H. Li, C. Wang, L. Yin, ACS Energy Lett 3 (2017) 132–140. [12] J. Li, X. Yuan, C. Lin, Y. Yang, L. Xu, X. Du, J. Xie, J. Lin, J. Sun, Adv. Energy Mater. 7 (2017) 1602725. [13] M. Boota, B. Anasori, C. Voigt, M.Q. Zhao, M.W. Barsoum, Y. Gogotsi, Adv. Mater. 28 (2016) 1517. [14] D. Zhao, R. Zhao, S. Dong, X. Miao, Z. Zhang, C. Wang, L. Yin, Energy Environ. Sci. 12 (2019) 2422–2432. [15] X. Hui, R. Zhao, P. Zhang, C. Li, C. Wang, L. Yin, Adv. Energy Mater. (2019) 1901065, https://doi.org/10.1002/aenm.201901065. [16] X. Yu, D.A. Yohan, N. Michael, G. Yury, M.W. Barsoum, H.L. Zhuang, P.R.C. Kent, ACS Nano 8 (2014) 9606–9615. [17] Y. Wu, P. Nie, J. Wang, H. Dou, X. Zhang, ACS Appl. Mater. Interfaces 9 (2017)

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