Interfaces Modification for Vacancies Enhancing Lithium

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Surfaces/Interfaces Modification for Vacancies Enhancing Lithium Storage Capability of Cu2O Ultra-Small Nanocrystals Huawei Song, Yue Gong, jian su, Yinwei Li, Yan Li, Lin Gu, and Chengxin Wang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b11592 • Publication Date (Web): 20 Sep 2018 Downloaded from http://pubs.acs.org on September 20, 2018

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ACS Applied Materials & Interfaces

Surfaces/Interfaces Modification for Vacancies Enhancing Lithium Storage Capability of Cu2O Ultra-Small Nanocrystals Huawei Song a#, Yue Gong b#, Jian Su a, Yinwei Li c, Yan Li a*, Lin Gu b*, Chengxin Wang a* a

State Key Laboratory of Optoelectronic Materials and Technologies, School of Materials Science and

Engineering, Sun Yat-Sen (Zhongshan) University, Guangzhou 510275, People’s Republic of China b

Laboratory for Advanced Materials & Electron Microscopy, Beijing National Laboratory for Condensed

Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing 100190, People’s Republic of China c

School of Physics and Electronic Engineering, Jiangsu Normal University, Xuzhou 221116, People’s Republic

of China

KEYWORDS: Surface/Interface modification, Vacancy, intercalation pesudocapacitance, Ultra-small nanocrystals

# These authors contributed equally to this work. * Correspondence and requests for materials should be addressed to: E-mail: [email protected] (C. X. Wang), Tel & Fax: +86-20-84113901; [email protected] (Y. Li); [email protected] (L. Gu). 1

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ABSTRACT

Theoretically, Cu2O delivers a poor Li storage capacity ~373.9 mAhg-1 based on a so-called conversion reaction (Cu2O+2Li→2Cu+Li2O). Herein, we broke through the bottleneck, and acquired an impressive lithium storage capability (1122 mAhg-1) tripled more than the theoretical one by an in-situ surfaces/interfaces engineering process for the first time. The surfaces/interfaces modification enabled us to fabricate ultra-small nanocrystals of Cu2O with Cu vacancies (VCu) of high concentration, somewhat like monovalent-anion doping. Except for the conversion-reaction-typed capacity, VCu enhancing intercalation pesudocapacitance in Cu2O and its reduction product-Cu also contributed a lot to the Li-storage capability. First-principles calculation substantiated that intercalation energy of Li was severely lowered for both Cu-vacancy-riched Cu2O and Cu comparing with their stoichiometric counterparts. Another important factor for the enhancement was the surfaces/interfaces organic species themselves which could reversibly store Li by redox reactions. The surfaces/interfaces modification for vacancies, vacancy inheritance from metal oxide to single metal, and vacancy-enhancing Li-storage capability in metal oxide and single metal all will inspire us a lot in fabricating new-generation advanced electrodes for rechargeable batteries.

2


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INTRODUCTION Generally, the strategies towards electrochemical energy storage could be faradic and non-faradic. The former worked through redox reactions which could be diffusion-controlled as in rechargeable batteries or capacitive in pseudocapacitors, depending on the kinetics factors, while the latter stored charges through ions electrostatic adsorption on the surfaces/interfaces such as materials with high specific surface in double layer capacitors (DLCs)

1-2

. In rechargeable batteries, fabricating nanostructures capable of storing ions through

multi-electrons redox reaction afforded to elevate the energy storage capability. Specifically, Sn-, Si- based materials, and various transition metal oxides (TMOs), nitrides, sulfides, and fluorides had been selected to design advanced electrodes for lithium ion batteries 3-9. The initial purpose mainly lied in improving utilization of the active materials, i.e. raising the diffusion-controlled charge storage performance, which benefited from the impressive advantages of nanomaterials such as high majority surface atoms, shortened ion diffusion depth, and enhancing reactivity, in spite of resulting in a good redox pesudocapacitance due to the enlarged surface area. In addition, designing materials with layered or tunneled crystal structures, large layer intervals or tunnels or gaps, and abundant active intercalating sites was also popular to realize efficient diffusion-controlled energy storage, particularly those with large ion radius (Na+, Mg2+, and Ca2+). The typical prototypes were nanostructures of metal-organic framework which are usually characteristic of relative fast kinetics due to the large intervals and gaps

10-15

. As to the improvement in capacitive energy storage, there were fewer reports on

intercalating pesudocapacitance, mainly about redox pesudocapacitance contributed by surface redox reaction in electrochemical supercapacitors 1-2, 16-19. The intercalation pesudocapacitance could be defined as faradic charge storage by ions intercalation that happened in the same timescale with that for capacitive charge storage. Typical examples were two dimensional (2D) carbides and sulfides which intrinsically possessed large interlayer spacings for fast ions intercalation kinetics

7, 20-21

. Vacancies in TMOs nanocrystals, affording abundant sites for ion intercalation and shortened 3



ion-diffusion depth, also were verified to enhance the intercalation pesudocapacitance

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. Consequently, it is

now very meaningful to fabricate TMOs nanocrystals with abundant vacancies. In the meantime, vacancies in electrochemical inactive nanocrystals such as Fe and Cu for further storing ions (Li+) have been never reported, which also may play important roles to further elevate the energy storage capability of various advanced electrodes. Herein, we reported surfaces/interfaces modification of ultra-small Cu2O nanocrystals (~3 nm) to bring about abundant Cu vacancies for enhancing the Li storage capability. After electrochemical reduction, the formed Cu nanocrystals well inherited the parent vacancies and exhibited excellent Li-ion intercalation charge storage capability never reported before. RESULTS AND DISCUSSION The surfaces/interfaces of Cu2O modified with organic monovalent anions, as illustrated in Scheme 1a, worked similarly with those of F- or OH- doped TiO2, readily generating abundant cations vacancies

22

. To

realize the supposition, micro/nano hollow structures of Cu (II)-MOF were firstly achieved through PVPdirected solvothermal treatment of copper acetate (Figure S1). The Cu (II)-MOF precursor rinsed with ethanol was then mildly annealed at vacuum atmosphere, resulting in the formation of Cu2O ultra-small nanocrystals and their in-situ encapsulation of organic layers full of organic monovalent anions such as -COO-, -CO-, -NHand etc. Herein, ethylene glycol served as one of the ligands in the Cu-MOF, which probably co-influenced with PVP and acetate radicals on the surface/interface decorating groups. Those chemically bonded organic monovalent anions might serve as valence-mismatching anions in Cu2O, somewhat taking after surface/interface doping. Suppose the surfaces/interfaces of Cu2O nanocrystals appeared with the most closepacked planes of (110), in which terminal O atoms were all modified with organic monovalent anions in the organic coating layer. The concentration of VCu could be predicted when the size of Cu2O crystals changed (VCu (at. %) =0.452/d*100%, herein, d (nm) represents the size of the crystals). The theoretical prediction verified the VCu concentration exhibited a strong dimensional dependency (Scheme 1b). When the dimension of Cu2O 4


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crystals was reduced to 3 nm, the concentration would rise up to a high level of 15%. In contrast, stoichiometric Cu2O nanocrystals with non-modified interfaces were also fabricated through a modified solvothermal methods. (See the Experimental Parts in the Supporting Information, Figure S2). The diffraction peaks of powder XRD patterns in Figure 1a well consistent with that of literature cubic Cu2O indicated both samples were phase-pure Cu2O crystals (Space group: Pn-3m, a=0.426 nm). The wide diffraction peaks verified those crystals were in nanosize. The wider were the diffraction peaks, the smaller dimensions did the crystals have. Figure S3 presented the XPS, EDX, and FT-IR spectra of the as-prepared Cu2O nanohybrids in comparison with those of the precursor Cu-MOF. The markedly satellite peaks around 942 eV and 962 eV for the Cu 2p core XPS in Cu-MOF and their disappearance for that of Cu2O nanohybrids, in combination with the higher binding energy of Cu 2p3/2 peak verified Cu element in Cu-MOF was cupric, while that in Cu2O nanohybrids was cuprous

23-24

(Figure S3a). Weakening intensity of C=O and its red shift from

Cu-MOF to Cu2O nanohybrids in the deconvoluted C1s XPS spectra, in addition to the appearing π-π* satellite peak, indicated -COO- functional groups were largely reduced in the annealing process due to pyrolysis (Figure S3b). EDX spectra confirmed the existence of C, O, and Cu in both materials, but the content of C and O in Cu2O nanohybrids was much lower (Figure S3c). N element non-detected in the EDX indicated its content was very little. On the whole, it was substantiated that only few of the organic species in the Cu-MOF were maintained in the Cu2O nanohybrids. Those included abundant monovalent organic functional groups such as COO-, -CO-, -NH-, etc as shown in the FT-IR spectra (Figure S3d). In addition to the organic functional groups, the Cu2O nanohybrid also exhibited a relative high BET surface area of 38.9 m2 g-1 in accompany with a type IV N2 adsorption-desorption isotherm and a wide range of BJH desorption pores (Figure S4a)25. A thermogravimetric analysis indicated the nanohybrid contained Cu2O of 68.8% (in weight) and various bonding organic functional groups in the surfaces/interfaces of 26.5%, except for some physically absorbed solvent molecules about 4.7% (Figure S4b). In comparison with stoichiometric Cu2O nanocrystals, the common absorbance band around 624 cm-1 ascribed to Cu-O stretching vibration mode 5



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for Cu2O in the FT-IR spectra indicated the core of both was mainly cuprous oxide, while the difference in the absorbance nearly 439 and 497 cm-1 reflected the interface-modified and interface-non-modified Cu-O vibrations of Cu2O (Figure 1b) 26-28. The shift and markedly varying intensity indicated the different vibration absorbance of surfaces/interfaces Cu-O due to transformation from chemically bonded to physically absorbed organic functional groups and the quite lessening amount as reflected by the varying absorbance intensity of CO (1000~1100, 1400~1480 cm-1), -COO- (1300-1400, 1500-1630 cm-1), - NH- (700~940), and etc 14, 29. The as-prepared Cu2O nanohybrids resembled the parent-Cu (II)-MOF in microstructures, exhibiting a multilayered porous global skeleton resulted from the Ostwald ripening mechanism 4 (Figure S5a and 5b). C, O, and Cu were uniformly distributed in the porous skeleton, which consisted of Cu2O nanocrystals with a dimension about 3~5 nm in good accordance with the powder XRD diffraction patterns and the ring-like polycrystalline SAED image ascribed to cuprites (Figure 2a-2c and Figure S5c-5e). Atomic resolution TEM image of a Cu2O nanocrystal allowed us to visually achieve the distribution of VCu according to the intensity variation and dark contrast which reflected the occupation change of Cu atom (Figure S5f and Figure S6). As displayed in the select parallelogram zone of 10×10 atom columns, the ratio of vacancies was up to 20% comparable to that achieved by theoretical calculations (15%) in consideration of another factor for vacancy formation due to O insufficient in vacuum atmosphere (Figure 2d). Figure 2e and 2f more clearly showed us two Cu vacancies indicated by the varying intensity and dark contrast in the corresponding colored image with a profile showing intensity variation of the marked rectangular zone. As mentioned above, those abundant vacancies would play a significant role in enhancing the related intercalation pesudocapacitance.

To substantiate the supposition, the Li-ion storage capabilities of both interface-modified and interfacenon-modified Cu2O samples were tested by fabricating lithium batteries with them as the cathodes and lithium plates as the anode, and 20 µL 1M LiPF6 in ethylene carbonate (EC) and diethyl carbonate (DEC) with a weight ratio of 1:1 as the electrolyte. All the capacities involved were calculated based on the weight of the 6


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composite. The typical charge-discharge curves, cyclic voltammograms, cycling performance, and rate performance tests (Figure 3a-3d, and Figure S7) clearly demonstrated that the VCu-riched interface-modified Cu2O nanohybrids exhibited a superior ion storage capability comparing with interface-non-modified Cu2O nanocrystals. When galvanostatically tested at a rate of 0.1 A g-1, the former delivered a superior reversible capacity ~1122 mAh g-1 three times more than that of the latter (~367 mAh g-1 approximate the theoretical value, in Figure 3a and 3b). Except for the multiplied charge-storage capability, the vacancies-riched nanohybrids also delivered excellent cycling performance with high capacity of 1163 mAh g-1 at 0.1 A g-1 for 100 cycles and 1067 mAh g-1 at 0.5 A g-1 for approximately 600 cycles (Figure 3b and 3d), outperforming all the reported Cu2O based electrodes ever reported (See Table S1, which listed most of the representative Cu2Obased electrodes ever reported for Li-storage)30. Moreover, an impressive rate performance (Figure 3c) was also achieved in the nonstoichiometric nanohybrids, specifically, 1143 mAh g-1 at 50 mA g-1, 1122 mAh g-1 at 100 mA g-1, 1089 mAh g-1 at 200 mA g-1, 1002 mAh g-1 at 500 mA g-1, 909 mAh g-1 at 1 A g-1, 755 mAh g-1 at 2 A g-1, 684 mAh g-1 at 5 A g-1, and 553 mAhg-1 at 10 A g-1respectively (versus 427 mAh g-1 at 50 mA g-1, 367 mAh g-1 at 100 mA g-1, 323 mAh g-1 at 200 mA g-1, 290 mAh g-1 at 500 mA g-1, 230 mAh g-1 at 1 A g-1, 192 mAh g-1 at 2 A g-1, 120 mAh g-1 at 5 A g-1, and 44 mAhg-1 at 10 A g-1 of interface-non-modified Cu2O nanocrystals). When the rate returned to 50 mA g-1, the storage capability quickly recovered nearby.

Interestingly, the slight increase in the capacity shown in the cycling tests was somewhat like that reported for inorganic-organic hybrids, which was probably attributed to the evolution of the organic functional groups in the surface/interface of the nonstoichiometric Cu2O nanocrystals

14, 31

(not the case common in various

nanostructures of transition metal oxide, which is to be explained in the following). As to the improved rate performance, it should be not only contributed by the porous microstructures and nonstoichiometric related vacancies enhancing fast charge carriers kinetics and diffusion processes (including the charge transfer of ions, electrons, and electron holes), but also highly dependent on the conductive organic matrix derived in the 7



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battery operation process. The improved conductivity could be reflected by the lowered charge-transfer resistances of interface-modified Cu2O versus interface-non-modified ones represented by depressed semicircles at the mid-high frequency region in the Nyqusit plots (Inset view of Figure S7, ~120 ohm vs. ~220 ohm), and the non-severely varied charge-transfer resistances of interface-modified Cu2O at different state-ofcharges (SOCs, 120 ohm for the pristine electrode, ~90 ohm for the lithiated electrode, and 140 ohm for the de-lithiated electrode) in the EIS spectra (Figure S8), while organic interface-modification derived superior microstructure could be verified by post-cycled electrodes characterization with Cu nanocrystals uniformly encapsulated by organic-Li species in the TEM images (Figure S9)32. The almost unchanged electrolyte resistances reflected by the same intercept value (~8 ohm) on the Z axis indicated the process of electrolyte decomposition was not marked after electrochemical lithiation and de-lithiation, unlike that of the progressively formation-decomposition of a gel-like film in various nanostructures transition metal oxides which would contribute to a marked successive rise in the capacity and a final electrode failure. It was worth mentioning that the Li-ion storage capability for the Cu2O nanohybrids was also superior to their precursor CuMOF and the fully anneal products of CuO (Figure S10). The Cu-MOF initially delivered a high reversible capacity ~800 mAhg-1 which probably also benefited from the surface/interface organic species like that for Cu2O nanohybrids, but decreased to ~ 584 mAhg-1 largely due to the unstable intermediates evolutions in-situ derived from the organic species

14

, while CuO exhibited an ordinary capacity of ~503 mAhg-1 and slightly

rose up to ~ 660 mAhg-1, approaching the theoretical capacity of CuO (673 mAh g-1, CuO+2Li→Cu+Li2O) 33.

Sweep voltammetry of interface-modified Cu2O nanohybrids was performed at different scan rates in comparison with the interface-non-modified ones (Figure S11) to analyze the kinetics. According to the equation (i=k1v+k2v0.5), the response current (i) is linearly dependent on sweep rate (v) in capacitive chargestorage processes and square root of sweep rate (v0.5) in diffusion-controlled charge-storage processes respectively 1, 16. Consequently, the values of k1 and k2 could be defined with two different voltages in the CVs. 8


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According to the calculated k1 and k2, the capacity contribution of diffusion-controlled and non-diffusioncontrolled parts (mainly capacitive-controlled) in the CVs could be achieved. As shown in Figure 3e and 3f, the capacitive charge-storage capability of Cu2O had been significantly enhanced through interfacemodification engineering Cu vacancies similar with those reported in MoO3-x and F-doped TiO2 2, 22. Moreover, the diffusion-controlled charge-storage capability was also revitalized due to smaller nanocrystals, as more large peaks and larger area are presented in the cyclic voltammograms. Those agreed well with our abovementioned supposition of vacancies-enhancing intercalation pesudocapacitance.

To fully understand the vacancies-enhancing intercalation capacitance, we further analyzed the Li-ion storage mechanism in the nonstoichiometric Cu2O nanohybrids. As shown in Figure 4a, the discharging profile comprised of three successive sloping plateaus around 2.2 V, 1.4 V, and 0.3 V respectively. Among them, the plateau at 1.4 V was usually ascribed to the generally accepted conversion reaction process of Cu2O reduction to Cu well consistent with those in various copper oxide nanostructures and composites previous reported

33-36

. The first plateau at ~2.2 V was firstly observed in CuO/Cu2O nanoparticles after initial cycle,

and hence had not attracted any attention

37

. Later reappearance in CuO, it was attributed to the reduction of

CuO accompanying with the formation of mixed valence copper oxides and Cu2O33, 38. However, based on our ex-situ XRD results in Figure 4b, the diffraction peaks around 36.4o, 42.3o (2theta) ascribed to Cu2O lasted until 1.75 V with only marked shift to high 2theta range probably resulted from severe structural deformation due to introduction of large amount Li. The interesting phenomenon indicated the plateau-2.2 could not be just simply explicated as the reduction of Cu (I). Due to the introduction of Cu-vacancy, a more sophisticated lithiation mechanism had happened. As to the last slope below 1 V, Cu2O had been already converted into Cu nanocrystals (Figure 4b), a Cu (I)-reduction process could not be prolonged to the stage, which also should be focused. As confirmed by ex-situ FT-IR spectra in Figure 4c, organic species as –NH (780~840 cm-1), -COO(1300~1600 cm-1), and -C=C (~1670 cm-1) had also suffered severe bonding transformation leading to a 9



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reversible redox reaction of C-O, C=O or C=C functional groups similar with previous reported organic counterparts14, 39-40.

Except for XRD and FT-IR characterization, the HR-TEM and atomic resolution TEM images of the performed electrodes at various stages of charge and discharges (SOCs, SODs) showed us the more direct results being caused by Li-ion intercalation. The severely deformed lattices fringes for the electrodes preformed at SOD of 1.75 V indicated series amount of Li ions had intercalated into Cu2O before conversion reaction (Figure S12a and 12b). The Li intercalation in this stage accounting for a discharge capacity of ~200 mAhg-1 largely relied on the Cu-vacancy enhancing intercalation pesudocapacitance similar with those in TiO2 and MoO3

1-2, 15

. As to the 1V stage, The atomic resolution images of newly formed Cu nanocrystals around

SOD of 1 V exhibited a large number of vacancies with a ratio up to 43%, which probably confirmed previously intercalated Li-ion had occupied those lattice points of Cu somewhat like solutions and superstructures, although we had not observed them limited by the resolution of TEM (Figure 5(a) and (b), Figure S12(c) and Figure S13). Moreover, the gradually widening lattice fringes ascribed to deformed Cu nanocrystals in the atomic resolution TEM image (Figure 5c), the concentrated Li distribution in Cu nanocrystals confirmed by element mapping images for fully lithiated (SOD of 0.005V) (Figure 5d), and residual Li only in the center of Cu nanocrystals for the partially de-lithiated Cu (SOC of 2.1 V) nanocrystals (Figure 5e and 5f) all substantiated many Li-ions have intercalated into the vacancies-riched newly formed Cu nanocrystals in the stage.

Consequently, it’s not hard to see that Li storage in the Cu2O nanohybrids included not only traditional conversion-reaction part between Cu2O and Cu, but also the abnormal energy storage phenomena of Cu2O and Cu themselves. The phase evolutions of Cu2O in the lithiation-delithiation were presented in combination with the typical cyclic voltammogram and discharge-charge curves (Figure 4a and Figure 6a). Considering the vacancy-enhancing capacitance energy storage processes, it’s not surprising that Li-ion storage capability in 10


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the interface-modified Cu2O would be significantly enhanced due to the considerable amount of vacancies. Theoretical calculations based on the density functional theory (DFT) confirmed the energy essential for Liions intercalating into non-stoichiometric Cu2O with various vacancies had been severely lowered in comparison with those of stoichiometric counterparts (Figure 6(b) and Figure S14). However, Cu vacancies are not enough to accommodate such a large number of Li-ions accounting for the discharge capacity before 1.75 V (~200 mAh g-1). Therefore, interstices would be the sites for further Li ions intercalation as the Cu vacancies were fully occupied by Li. Our DFT simulations demonstrated the intercalation energy of one additional Li+ into non-stoichiometric Cu2O with 20% Cu vacancies occupied by Li atoms was -0.56 eV indicating further Li+ intercalation into interstices was energetically favored (Figure S15(b)).

For the vacancy enhancing Li-storage in Cu nanocrystals, our DFT-calculations also revealed the energy for Li-ion intercalation into Cu with substituted Li of 43% could be lowered for 1.23 eV comparing with perfect counterpart (interstices in Li43Cu57 versus those in perfect Cu) (Figure 6 (c)). As to the formation of vacancies in newly formed Cu nanocrystals, we speculated that Li ions intercalation into Cu2O pre-conversion reaction had occupied some of the sites for Cu in the formation of vacancy-riched Cu nanocrystals, while subsequently intercalated Li ions participated into the Cu2O reduction reaction. In comparison with the experimental XRD pattern of Cu2O nanohybrids, there were no change for the simulated XRD pattern of Cu2O with Cu vacancy of 20% and its Li-occupied-vacancy counterparts (Figure S16). Only with another Cu interstices of 23% occupied with Li ions, a distinct change of the relative intensity for peaks ascribed to (111) and (200) crystal planes had happened, but still similar with the peaks for the experimental XRD at SOD of 1.75 V (Figure S16). Those results all well supported the above-mentioned speculation. In consideration of charge conservation law, it seemed that the vacancy enhanced charge storage which were not accompanied with the reduction of Cu species in the lithiation process, should result into reduction reaction of another surrounding species such as various organic functional groups, being indicated by the reversible conversion of 11



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νN-H, νC-O, and νC=O stretching vibrations (791 cm-1, 1080 cm-1, and 1584 cm-1). Actually, those monovalent organic groups acted as O-, which could be further reduced with Li storage. In this sense, surfaces/interfaces modifications were essential for those abnormal Cu-vacancy enhanced Li storage.

CONCLUSIONS

In summary, we presented a facile strategy to significantly improve the energy storage capability of Cu2O. Due to surface/interface modification engineering Cu vacancies, the charge-storage capability of ultra-small nanocrystals of Cu2O has been elevated up to three times higher than that of traditional counterpart. Based on the nanohybrids, we substantiated there was Li-ion storage process based on intercalation-like reaction in Cu2O and Cu themselves, except for the common conversion reaction based on redox. And consequently, nanocrystals of Cu2O and Cu both exhibited impressive vacancies-enhanced intercalation pesudocapacitance. Moreover, the success in Cu2O will probably be also applicable to other various advanced electrode materials.

ASSOCIATED CONTENT

Supporting Information The Supporting Information is available free of charge on the ACS Publications website Supporting Information Available: **[ Experimental part, Theoretical calculation results, additional XRD, XPS, EDX, FT-IR, TG, BET, SEM, and TEM characterization and electrochemical test, structure refinement results ]**

AUTHOR

INFORMATION

Corresponding Author

12


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E-mail: [email protected] (C. X. Wang), Tel & Fax: +86-20-84113901; [email protected] (Y. Li); [email protected] (L. Gu). ORCID Huawei Song: 0000-0001-9091-2297 Yue Gong: 0000-0002-5764-3117 Jian Su: 0000-0003-0239-1351 Yinwei Li: 0000-0002-9974-504X Yan Li: 0000-0002-7099-6830 Lin Gu: 0000-0002-7504-031X Chengxin Wang: 0000-0001-8355-6431 Notes The authors declare no competing financial interest.

ACKNOWLEDGEMENTS

This work was supported by National Natural Science Foundation of China (grants Nos. 51602355, 51872339, 51522212, 11722433 and U1401241).

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Schematics

Scheme 1 (a) Interface modifications with organic monovalent anions for Cu vacancy (VCu) formation in Cu2O nanohybrids; (b) theoretical calculations depict dimensional dependence of VCu concentration in interfacemodified Cu2O crystals.

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Figures

Figure 1 (a) Powder X-ray diffraction pattern and (b) FT-IR spectrum of the interface-modified Cu2O nanohybrids versus that without interfaces modification.

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Figure 2 Cu2O nanohybrids with ultra-small crystals and abundant interface-modifications enhanced vacancies (a) STEM image, (b) SAED image, (c) HR-TEM image, (d), (e) and (f) atomic resolution TEM images of a Cu2O nanocrystal, and its colored view with a profile showing intensity variation of the marked rectangular zone indicating the existence of abundant vacancies, confirmed as those of Cu by the corresponding atom stacking profiles in the top-right view (Cu in green balls and O in red balls), high ratio of vacancies up to 20% was implied in the marked parallelogram zone with vacancies highlighted by dashed lined small circles (see Figure S5f and S6).

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Figure 3 Vacancies enhanced Li-ion charge storage capability of Cu2O with interface-modified versus interface-non-modified nanocrystals (a) typical charge-discharge profiles at 0.1 A g-1, (b) galvanostatic cycling performance at 0.1 A g-1 and (c) rate performance at varying currents from 50 mA g-1 to 10 A g-1, (d) cycling stability of interface-modified Cu2O nanohybrids at 0.5 A g-1 in the long term test, and diffusion-controlled versus capacitive capacity contribution analysis for (e) interface-modified and (f) non-interface-modified Cu2O 22


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nanocrystals in the cyclic voltammogram at 0.3 mV s-1 based on calculated k1 and k2 (i=k1v+k2v0.5, See Figure S11).

Figure 4 Ex-situ XRD and FT-IR analysis of the Li storage mechanism in the as-prepared Cu2O nanohybrids electrodes (a) typical discharge-charge curves with marked state-of-discharges (SODs) and charges (SOCs), and corresponding Cu-containing species confirmed by ex-situ XRD patterns, (b) ex-situ XRD patterns for the electrodes at selected SOCs and SODs, (c) ex-situ FT-IR spectra for the electrodes at selected SOCs and SODs.

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Figure 5 Atomic-resolution TEM analysis of the vacancies-enhanced Li-ion storage mechanism in Cu nanocrystals generated from VCu-riched Cu2O nanocrystals after conversion reaction, the voltages for electrodes at selected states of charge and discharge (SOC and SOD) are marked for each image (a) a Cu nanocrystal with abundant vacancies confirmed by varying profiles intensity of different lines shown in Figure S12c and Figure S13, and high ratio of vacancies up to 43% highlighted by dash-lined small circles was implied in the marked parallelogram zone, (b) the corresponding colored image of (a) and the equivalent theoretical atom packing pattern of Cu is shown in the top-right view, (c) Cu nanocrystals after full lithiation with gradually widening interplanar crystal spacing resulted from Li+ intercalating and (d) their element mapping images showing Li uniformly distributing in the Cu nanocrystals, (e) Li+ de-intercalating out of Cu nanocrystals, and (f) the corresponding element mapping images indicating residual Li+ in the core of Cu nanocrystals.

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Figure 6 Phase evolutions of interface-modified Cu2O ultra-small nanocrystals in Li-ion charge storage and theoretical analysis (a) CV curve marked with corresponding Cu-containing species confirmed by ex-situ XRD patterns at different stages of charge or discharge (SOCs or SODs), (b) DFT-calculated intercalation energy variation for Li+ intercalating into Cu2O of stoichiometric versus non-stoichiometric at various circumstances before the conversion reaction (Cu in blue balls, O in red balls, and Li in cyan, see Figure S14, Table S2 and S4), (c) DFT-calculated energy for Li+ intercalating into interstices of Cu with/without substituted Li of 43% (interstices in Li43Cu57 versus those in Cu, see Figure S15).

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TOC

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Interfaces Modification for Vacancies Enhancing Lithium

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