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Metal (M ¼ Ru, Pd and Co) embedded in C2N with enhanced lithiumstorage properties
Chunmao Huang a, e, 1, Javeed Mahmood b, 1, Zengxi Wei c, Dan Wang d, Shenghong Liu a, e,Yanming Zhao e, Hyuk-Jun Noh b, Jianmin Ma c, Jiantie Xu a, **, Jong-Beom Baek b, *
a Guangdong Provincial Key Laboratory of Atmospheric Environment and Pollution Control, National Engineering Laboratory for VOCs Pollution ControlTechnology and Equipment, School of Environment and Energy, South China University of Technology, Guangzhou, 510640, Chinab School of Energy and Chemical Engineering/Center for Dimension-Controllable Organic Frameworks, Ulsan National Institute of Science and Technology(UNIST), 50 UNIST, Ulsan, 44919, South Koreac School of Physics and Electronics, Hunan University, Changsha, 410082, Chinad State Key Laboratory of Organic-Inorganic Composites, Beijing University of Chemical Technology, Beijing, 100029, Chinae School of Physics, South China University of Technology, Guangzhou, 510640, China
a r t i c l e i n f o
Article history:Received 7 August 2019Received in revised form9 October 2019Accepted 11 October 2019Available online 7 November 2019
Keywords:C2NHoley nanocarbonMetal incorporationAnodesLithium ion batteriesNitrogenated structures
The improvement of the high-efficiency electrode materials for lithium ion batteries is one of theresearch priorities. The research on the carbon with M � N active sites (M-N-C) as anodes for lithium ionbatteries is still quite rare. Herein, we report a series of highly crystalline M@C2N hybrids in which themetal (M ¼ Ru, Pd and Co) are uniformly embedded in the two-dimensional C2N networks. Together withthe unique structural features (e.g., uniform two-dimensional nanohole structure, high surface area andenlarged distance between C2N nanosheets), rich N-M coordination homogenously distributing throughthe M@C2N structure are favorable for the brisk infiltration of electrolyte, swift transportation of Liþ/electrons and more storage of Liþ, thereby leading to the superior lithium storage properties.
Owing to their environmental friendliness, long cycle life, highenergy density and falling battery costs, rechargeable lithium ionbatteries (LIBs) have been treated as one of leading technologiesapplied inelectric vehicles (EVs) andenergystorage systems (ESSs). Inorder to successfully implement the large-scaleapplicationof EVsandESSs, the improvement of electrode materials for LIBs with sufficientenergy density is insistently required [1]. The electrodematerials canbesimplydividedintoorganicandinorganicmaterials[2,3].Comparedto the inorganic electrodesdominating the currentmarket of LIBs, theorganicelectrodeshavebeenalso intensively recognizedaspromisingelectrode candidates. This ismainly due to their outstanding features,including the light molecular weights, eco-efficient sustainability,
multiple electron transfer capability and chemical diversity [4e9]. Todate, the organic electrodes can be divided into three major types,namely, carbonyl compounds, organosulfur compounds and freeradical compounds. The carbonyl compounds are common organicmaterials containing a functional group of C]O and thus have theabilityofoxidation[6].Asaresult,avastnumberofcarbonylcontainingmaterialsareconsideredascathodecandidatesforLIBs.Thisisbecausethe considerable capacities of themaremainly securedat>2V vs. Liþ/Li, such as dichloroisocyanuric acid [10], chloranil [11], nonylbenzo-hexaquinone (NBHQ) [12], polymeric quinone [13], polyimides [14]and polyaryltriazine [15]. Nevertheless, a few organic compoundsdeliver low average discharge potentials of ~0.8 V vs. Liþ/Li, such asconjugated carboxylates [16] and PPy-colbalt-oxygen complex [17].Despitetheachievedprogress, theorganicelectrodesoftensuffer frompoor crystallinity, low intrinsic electrical conductivity and high solu-bility innon-aqueous electrolytes [7,18]. These drawbacksmake themdifficultto fullyreleasetheircapacity.Furthercarbonizationof themtocarbon with well-defined structures and improved conductivity is acommonbut effective route [19]. After the carbonization, the resultedcarbon materials tend to show the higher capacity and better rate
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proficiency in comparison to graphite as the most commonly usedanode for commerical LIBs [20].
In order to further boost the capacityand rate capabilityof carbonbased materials mentioned above, the introduction of heteroatoms(e.g., N, F, S, and P) doping and metal incorporation to carbonframeworks are commonly empolyed. These strategy not only im-proves the electric conductivity of the carbon but also generatesadditional active sites (e.g., vacancy and enlarged accessible surfacearea) for Liþ storage. Inparticular, a large number of N-doped carbonhave been intensively applied as anodes for high-performance LIBs[21], along with their most popular applications in the field of oxy-gen reduction reaction (ORR) [21e23]. Recently, many studies haverevealed that the doped nitrogen atoms coordinated bymetals (e.g.,Co, Fe and/or Ni) enabled nitrogenemetal (e.g., Co, Fe and/or Ni)-carbon (N-M-C) catalysts with superior electrocatalytic activity forORR [24,25], as well as hydrogen evolution reaction (HER) [25e27]and oxygen evolution reaction (OER) [26,28]. Compared to manyprogress achieved in thefields of electrocatalysis, direct studyon theN-M-CwithoutM relatedmass (e.g., oxdies, sulfides and carbides) asanodes for LIBs is still quite rare. As expected, the coordination be-tween N and M in the N-M-C as anode for LIBs could also provideadditional and efficient active sites for Liþ storage and trans-portation, as well as those using N-M-C as catalysts [29e38].Moreover, the existence of nanopores (i.e., nanoholes) within thecarbon frameworks could facilitate the brisk infiltration of electro-lyte, quick diffusion of ions and creation of copious edges as activesites for ion storages, as well as easy introduction of heteroatomswith high doping level [39e42]. However, it is difficult to experi-mentally control the carbon frameworks with the desired holes,porous structures and their uniform distribution. The preparation ofholey carbon often involves the use of catalyst or strong acid, whichcould destroy the pristine structure in different degree. Thus, it ishighly preferable to explore novel organic materials (especiallycontaining N and M) with abundant open channels, high electronicconductivity and numerous active sites, allowing more Liþ rapiddiffusion and efficient storage.
Recently, a novel carbon nitride, nitrogenated 2D layeredstructure (C2N) was synthesized by wet-chemical reaction (Fig. S1)[43]. The holes and N/C atoms are uniformly distributed throughthe whole C2N structure. Based on TGA anlaysis (Fig. S2), the C2Nwas further annealed at 450 �C for 4 h to enhance the layer to layerinteraction. As displayed in Fig. S3, the texture of the resultingmaterial (C2N) almost maintains its original morphology. When itwas measured as anode for LIBs, however, its stable cycling dura-bility and rate capability are yet unsatisfied [44]. The plausiblereason could be the limited electronic conductivity associated witha optical band gap of 1.70 eV (vs. zero bandgap for graphene), aswell as its bulk pattern, thereby leading to low utilization of activesites and holes. Inspired by the progress on the N-M-C with theenhanced catalytic activity, herein, we report a series of highlycrystalline M@C2N hybrids. The hybrids are composed by the M/MOx (M¼ Ru, Pd and Co) embedded in the two-dimension (2D) C2Nnetworks. As shown in Scheme 1, the hybrids display rich N-Mcoordination disrtributing through the M@C2N along with uniform2D nanohole structure, high surface area, and enlarged distancebetween C2N nanosheets. As expected, these structural features aremore favorable for the rapid diffusion and more storage of Liþ.
2. Experiments
2.1. Synthesis of C2N and M@C2N (M ¼ Ru, Pd and Co)
First of all, metal salt (4.82 mmol RuCl3, 5.64 mmol PdCl2 or8.50 mmol CoCl2) was dissolved in 50 mL of deoxygenated N-Methyl-2-pyrrolidone (NMP) in a 3-necked round bottom flask and
placed on the ice bath under argon atmosphere. Secondly, 2.81 g ofhexaketocyclohexane octahydrate (HKH) and 2.50 g of hex-aaminobenzene trihydrochloride (HAB) were added in flask and letslowly warm up ambient condition for 2 h. Then, the ice bath wasexchanged with oil bath and further warmed to 175 �C for 8 h. Aftercooling down the reaction mixture to 80 �C, slowly 40 ml of sodiumborohydride (NaBH4) solution (10 wt% in NMP) was added. Theresulting solution was again refluxed at 175 �C for 3 h. The reactionsolution was then cooled down to room temperature and pouredinto water (in case of Ru@C2N and Pd@C2N, pour into acetone). Theblack solid products were obtained on polytetrafluoroethylene(PTFE) (0.5 mm) membrane by suction filtration. These productswere further washed by Soxhlet extraction with methanol andwater respectively, and freezed dried under reduced pressure(0.05 mmHg) at �120 �C. In case of the Ru@C2N and Pd@C2N, theprecipates were purified by Soxhlet extraction using acetone andmethanol for three days and one day, respectively. Finally, thefreeze dried products were further heat treated at various tem-perature under inert atmosphere for 2 h. It should be noted thatafter heat treatment the final products were treated with HCl (4 M)for 4 h to get rid of unreacted free standing metallic impurities, ifany. The C2N sample was prepared by the same way without theuse of metal salt and NaBH4 (10% NMP solution).
2.2. Material characterization
The materials structures were identified by powder X-raydiffraction (XRD) using a X-ray diffractometer, Bruker D8 advance,with CueKa radiation (l ¼ 1.5406 Å). The morphologies were char-acterized by scanning electron microscopy (SEM, JEOL JSM-6360LV)and transmission electron microscopy (TEM, HitachiH-9500). Theelemental mappings were collected by the EDX coupled SEM. Ther-mogravimetric analyses (TGA) were carried out with ramping rate of10 �C min�1 using a Thermogravimetric Analyzer Q200 (TA Instru-ment, USA). The nitrogen (N2) gas adsorption-desorption isothermswere obtained using the BET method on BELSORP-max (BEL Japan,Inc., Japan). X-rayphoto-electron spectroscopy (XPS)wereperformedon a Kratos Axis Ultra spectrometer using C 1s (B.E. ¼ 284.6 eV) as areference. Raman spectra were recorded on a Raman spectrometer(Horiba) using an excitation wavelength of 633 nm.
2.3. Electrochemical measurements
The electrochemical characterization of the samples was eval-uated using coin cells. The working electrodes were fabricated bymixing the active materials (M@C2N/C2N, 80 wt%), acetylene black(AB, 10 wt%) and polyvinylidene fluoride (PVDF, 10 wt%) binder inthe solvent of NMP. The obtained slurries were uniformly coated onCu foil and then heated at 90 �C overnight. Then, the electrode waspunched into a round disk with a controlled mass loading of1.5 mg cm�2. The coin-type cells were assembled in an argon-filledglove box (Mikrouna, H2O, O2 < 0.1 ppm) using the as-preparedelectrode as the working electrode, lithium pellet as the counterand reference electrode, a porous polypropylene as the separatorand 1 M LiPF6 in a mixture of EC/DEC (ethylene carbonate/diethylcarbonate, 1:2 (v/v)) as the electrolyte. The discharge-chargemeasurements were carried out by an automatic battery testersystem (Land®, China). Cyclic voltammetry (CV) measurementswere performed at a scan rate of 0.1 mV s�1 from 0.02 to 3.0 V (vs.Liþ/Li) using an AUTOLAB PGSTAT100 type electrochemical work-station (Metrohm AG company). Electrochemical impedancespectroscopy (EIS) was measured on an AUTOLAB PGSTAT302N(Metrohm, Netherlands) with an Ac signal of 5 mV in the frequencyrange of 0.1 Hze1 MHz. All the measurements were carried out atroom temperature.
Scheme 1. Scheme of the synthesis and lithiathion-delithiation mechanism of M@C2N (M ¼ Ru, Pd and Co) and C2N.
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2.4. Computational details
Density functional theory (DFT) calculations were accomplishedby using the Vienna Ab Initio Simulation Package (VASP) code[45,46]. The projector augmented wave (PAW) approach wasadopted to describe the ionic cores and the electronic exchange-
Fig. 1. Characterizaiton of the materials. (a) XRD patterns, (b) Raman spectra, (c) BET me(M ¼ Ru, Pd and Co) and C2N in air condition at ramping rate of 10� min�1.
correlation energy was defined by the Perdew-Burke-Ernzerhof(PBE) functional within the generalized gradient approximation(GGA) [47,48]. The cut-off energy was set as 520 eV. A Monkhorst-Pack 3 � 2 � 1 k-point grid was used as the Brillouin zone. Theconvergence criterion for the electronic structure iteration was setto 10�6 eV, and 0.02 eV/Å for the force. A 20 Å vacuum in the z
asurements and inset corresponding pore size distribution (d) TGA curves of M@C2N
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direction was obatined in the optimized geometry. The DFT-D3scheme was considered in the van der Waals (vdW) interaction.
3. Results and discussion
3.1. Structural and Compositional analysis of M@C2N (M ¼ Ru, Pdand Co) and C2N
Fig. 1a shows the XRD patterns of M@C2N (M ¼ Ru, Pd and Co)and C2N. Apart from the broad diffraction peak at ~25.5� indexedto C2N, the characteristic peaks of cubic hexagonal Ru (ICSD#43710), cubic Pd (ICSD #41517) and Co (ICSD #52935) are clearlyobserved in the M@C2N (M ¼ Ru, Pd and Co), respectively. Itshould be noted that minor diffraction peaks correspond to hex-agonal Co (ICSD #52934). Compared to Ru@C2N and Co@C2N, thePd@C2N has higher crystallinity M (Pd). This is evidenced by thepeaks intensity ratio of the M/C2N (M ¼ Ru, Pd and Co) in theorder of Pd > Ru > Co, as well as the well-define and strong
Fig. 2. Study of bonding nature by XPS. (a) XPS of M@C2N (M ¼ Ru, Pd and Co) and C2N. High, i and p) O 1s peaks of (b-d) C2N, (e-h) Ru@C2N, (i-l) Pd@C2N and (m-p) Co@C2N.
characteristic peaks of Pd. Fig. 1b displays the Raman Spectra ofM@C2N (M ¼ Ru, Pd and Co) and C2N. As can be seen, all thehybrids show two typical D and G bands located at 1350 and1595 cm�1, associated with defects/disorder and graphitic latticevibration mode with E2g symmetry, respectively [49]. Compared toC2N, the M@C2N (M ¼ Ru, Pd and Co) have larger amounts ofdefects than C2N (<1), which is reflected by their higher intensityratio of ID/IG (>1). The specific surface areas of the M@C2N(M ¼ Ru, Pd and Co) and C2N (Fig. 1c) were measured using theBrunauer-Emmett-Teller (BET) method. The specific surface area ofM@C2N (M ¼ Ru, Pd and Co) and C2N are 312.9, 452.2, 402.9 and351.1 m2 g�1, respectively. The smaller surface area and total porevolume (Table S1) of Ru@C2N is most likely due to its largeramounts of Ru metals occupying at holes. This is further confirmedby the thermogravimetric analysis (TGA) results (Fig. 1d) with ahigher estimated value of ~33.7 wt% weight residuals for Ru@C2N,compared to 26.1 wt% and 18.7% for Pd@C2N and Co@C2N,respectively.
h resolution XPS spectra of (e, i andm) M, (b, f, j and n) C 1s, (c, g, k and o) N 1s and (d,
Fig. 3. SEM and TEM characterization. (a, h and o) SEM images and (b-d, i-k and p-r) elemental mappings, (e-f, l-m and s-t) TEM images and (g, n and u) SAED for (a-g) Ru@C2N,(h-n) CoC2N and (o-u) Pd@C2N.
C. Huang et al. / Materials Today Energy 14 (2019) 100359 5
The elemental compositions and their bonding states ofM@C2N (M ¼ Ru, Pd and Co) and C2N, X-ray photoelectron spec-troscopy (XPS) was performed. As shown in Fig. 2a, all the samplesshow the peaks of C 1s at ~285 eV, N 1s at ~399 eV and O 1s at~532 eV. The characteristic M (Ru, Pd and Co) peaks of M@C2N areobserved at ~ 283, ~340 and ~781 eV, respectively. The presence ofoxygen in the samples is cheifly due to trapped or physicallyadsorbed oxygen and moisture as well as diketonic groups at theedges of C2N. The decovoluted XPS spectra of M, C 1s, N 1s and O1s for M@C2N (M ¼ Ru, Pd and Co) and C2N are shown in Fig. 2b-p.As can be seen, the C 1s spectra (Fig. 2b, f, j and n) are assigned toCeC at 284.6 eV, CeN at 286.0 eV and C-heteroatom at 288.8 eVwhile N 1s spectra (Fig. 2c, g, k and o) can be resolved into N-pyrazine like at 398.8 eV and “pyrazine like” N coordinated withmetal at 400.4 eV [50]. The O 1s peaks (Fig. 2d, h, l and p) are well
fitted to O]O, C]O and HeOeH at 531.1, 532.0 and 533.1 eV,respectively [43]. Fig. 2e, i and m show high resolution M XPSspectra of M@C2N. For the Ru 3d spectrum of Ru@C2N (Fig. 2e),there are two typical Ru 3d3/2 at 281.2 eV and Ru 3d1/2 at 285.1 eV,corresponding to the metallic Ru [51,52]. Similarly, the character-istic peaks of metallic Pd and Co can be also observed in the Pd 3d(Fig. 2i) and Co 2p spectra (Fig. 2m) [50,53].
The morphologies of the M@C2N (M ¼ Ru, Pd and Co) and C2Nwere characteriezed by scanning electron microscopy (SEM) andtransmission electron microscopy (TEM). As shown in Fig. 3aed, h-k and o-r, elemental mapping results from EDS reveal that the C, Mand N are homogenously distributed through the M@C2N struture.The M (M ¼ Ru, Pd and Co) particles uniformly distributed throughthe C2N structure can be further verified by TEM (Fig. 3e, l and s), inparticular M¼ Ru. High resoultion TEM (HRTEM) images of M@C2N
Fig. 4. Electrochemical characterization. (a) Rate capability measurements at various C-rates, (b) initial discharge/charge profiles of M@C2N and C2N at 0.1C and (c) cyclingperformance of M@C2N (M ¼ Ru, Pd and Co).
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(Fig. 3f, m and t) reveal that the M (M ¼ Ru, Pd and Co) particleswith a nanoparticle size of ~10 nm are well encapsulated in C2N.The selected area electron diffractions (SAED, Fig. 3g, n and u), onceagain, confirm M@C2N particles with the different crystallinity inthe order of Pd@C2N > Co@C2N > Ru@C2N, which is almost con-sitent with XRD results (Fig. 1a).
3.2. Electrochemical measurements of M@C2N (M ¼ Ru, Pd and Co)and C2N
Fig. 4a describes the rate capability of M@C2N (M ¼ Ru, Pd andCo) and C2Nmeasured at different C-rates from 0.1 to 10 C between0.02 and 3.0 V. Fig. 4b and S4 show the typical electrochemicalbehaviors of M@C2N and C2N, which agree well with previous re-sults of carbonaceous materials [54]. As shown in Fig. 4b, theM@C2N (M ¼ Ru, Pd and Co) and C2N deliver initial capacities of1104.2, 1168.3, 789.7 and 1231.6 mA h g�1 at 0.1C, with initialCoulombic efficiencies (CEs) of 66.9%, 52.7%, 66.7% and 46.9%,respectively. The massive irrevocable capacity in the first cycle (i.e.,low initial CE) are mainly appears by the formation of solid stateinterphases (SEI) [44]. This is consistent with the cyclic voltam-metry (CV) measurements of M@C2N at a scan rate of 0.1 mV s�1
(Fig. S5). As can be seen, there are typical reduction peaks at ~0.7 Vin the first cathodic scans, related to the formation of SEI. In thefollowing rounds, the disappear of reduction peaks at ~0.7 V sug-gests the evolution of a strong and stable SEI in the beginning cycle.Moreover, the highly overlapped CV curves in the following 4 cyclesconfirm the stable electrode structures (including SEI film) duringthe cycles. It is noteworthy that a minor reduction peak occur at~1.7 V for Ru@C2N (Fig. S5d) in the first cycle, associated to anobvious plateau at ~1.8 V (Fig. 4b), which is attributed to theinteraction between Liþ and partial RuOx [55]. This process may bean irrversible process because of its disappearance in the following4 cycles (Figs. S5f, S5d). When the C-rate is expanded from 0.2 to 10
C, the Ru@C2N and Pd@C2N electrodes, followed by Co@C2N, exhibithigher average specific capacities of 569.8 (527.8), 462.0 (415.3),413.9 (340.6), 361.7 (259.6), 290.6 (199.3) and 231.7 (135.1) mA hg�1 than C2N at 0.2, 0.5,1, 2, 5, and 10 C respectively, indicating theirexcellent rate capability (inset of Fig. 4a). The higher specific ca-pacities of M@C2N (M¼ Ru, Pd and Co) than that of C2N at a high C-rate indicating the structures with bigger nanohole are morefavorable for the rapid diffusion and more storage of Liþ.
For the sake of comparison the relative long-term cycling per-formance, the cells based on the M@C2N (M ¼ Ru, Pd and Co)electrodes were measured for 200 cycles (Fig. 4c). At 1 C, theRu@C2N provides an initial capacity of 483.6 mA h g�1, higher thanPd@C2N (396.5 mA h g�1) and Co@C2N (390.4 mA h g�1). After 200cycles, the Pd@C2N displays the cycling superiority with a higherinitial capacity retention of 46.2% than Ru@C2N (45.0%) andCo@C2N (41.8%). At a higher C-rate of 5 C, the Pd@C2N still shows itscycling superiority after 200 cycles with an initial capacity reten-tion of 47.0%, compared to 44.4% for Co@C2N and 38.4% for Ru@C2N.
Despite the capacities and cycling stability in differences, all theelectrodes show high and close average CE (~99.5%) over 200 cycles(Fig. 4c). To further investigate the electrode kinetics, the electro-chemical impedance spectra (EIS) of the M@C2N electrodes evalu-ated at 5 C before cycling and after 200 cycles were compared. Asillustrated in Fig. S6, EIS spectra of all the electrodes after cyclingdisplay a noticeable double semicircle (first and second) atmedium-high frequencies, and a flat line at small frequency, relatedto the internal resistance (including SEI film and electrolyte) andcharge-transfer kinetics-controlled resistance (Rct), and the mass-transfer-controlled Warburg resistance (W), respectively [56]. Thecorresponding model is formulated in inset of Fig. S6a. Conse-quently, Rct, a basic indicator of electrode kinetics, of Ru@C2N,Co@C2N and Pd@C2N are apparently reduced from 95.8, 131.7 and137.1 U to 13.6, 34.3 and 45.3 U, respectively. The drcreased Rct isprimarily because of the electrode-electrolyte activation. The order
Fig. 5. Work functions of M@C2N (M ¼ Ru, Pd and Co) and C2N and their corresponding schematic diagram of M � C2N interfaces.
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of Rct values almost agree well with their associated electro-chemical performance. The nitrogen doping in the carbon frame-works has been demonstrated to be a powerful concept for thefabrication of carbon materials with improved lithium storageproperties [57e59]. Mainly due to the nitrogen atom in the struc-ture which modulates the physical and chemical characteristics ofthe host structure, leading to improved electrical conductivity,additional defects as active sites and better electrode/electrolytewettability of the carbon electrode. Nevertheless, a large number ofN atoms in the carbon frameworks and/or their non-uniform dis-tribution may also result in structural instability of host structure[60]. It is expected that C2N with rich but uniform CeN distributioncould be an import platform for Liþ absorption and diffusion. Thisexplains well that the C2N has comparable or even higher initialcapacity but poor rate capability and cycling stability (Fig. 4a) thanM@C2N. Moreover, the poor cycling stability of C2N is also probablydue to its uptake of large amounts of Liþ during the lithiationprocess, which make it easy agglomeration and even formation ofirreversible Li metal at the hole edges. Unlike C2N, the N atomslocalized at the hole edges interacting with adequate M atoms canprovide M@C2N a more stable and attractive platform for Liþ stor-age and release, thus bring about excellent rate capability andcycling stability. Furthermore, C2N structure is a narrow band-gapsemiconductor with a direct band gap of 1.70 eV. M atoms canbridge the two adjacent C2N layers to decrease band gap and extendp- p conjugated systems, which could be also favorable for theenhancement of the electronic conductivity, and formation ofadditional accessible active sites for Liþ storage as well as enlargedchannels for Li diffusion, respectively.
Recent studies have revealed that conductive angents interact-ing with C2N could effectively reduce the work functions of C2Nwith the improved electronic conductivity [61e64]. For example, aseries of metal (M ¼ Al, Ag, Au, Pd, Pt, and Sc) contacts withmonolayer and bilayer C2N could significantly enhance the stronginteractions at the M � C2N interfaces, leading to the improvedelectron injection efficiency at the interfaces [61]. To further un-derstand the Liþ interaction at the interfaces of M � C2N (M ¼ Ru,Pd and Co) and C2N, the work functions were evaluated by first-principle density functional theory (DFT) calculations. As shownin Fig. 5, the highwork function of C2N (5.81 eV)make it difficult forlithium storage. Compared to C2N, theM� C2N (M¼ Ru, Pd and Co)
display lower work functions of 4.62 eV (Co@C2N), 4.88 eV(Ru@C2N) and 5.00 eV (Co@C2N). As expected, the reduced re-sistances of M� C2N are favorable for the improved lithium stroageproperties.
4. Conclusions
In summary, we have developed an effective bottom-up wet-chemical route to the synthesis of C2N and M@C2N hybrids as an-odes for LIBs. For theM@C2N hybrids, themetal (M¼ Ru, Pd and Co)are homogenously embedded in the 2D C2N networks. Benefitingfrom their unique structural features (e.g., uniform 2D nanoholestructure, high surface area and enlarged distance between C2Nnanosheets), the M@C2N electrodes display outstanding lithiumstorage properties, including high specific capacity (1104.2, 1168.3and 789.7 mA h g�1 for Ru@C2N, Pd@C2N and Co@C2N at 0.1C,respectively), excellent rate capability (413.9, 340.6 and333.4 mA h g�1 for Ru@C2N, Pd@C2N and Co@C2N at 1 C, 290.6,199.3 and 239.0 mA h g�1 for Ru@C2N, Pd@C2N and Co@C2N at 5 C,respectively) and stable cyclability (45.0%, 46.2%, and 41.8% initialcapacity retention after 200 cycles for Ru@C2N, Pd@C2N andCo@C2N, respectively). As expected, this work could not only offersa general approach for the rational design of holey structure withrich and even distribution of M-N-C as active sites for energystorage and conversion applications, but also provides an impor-tant guideline for the development of other nanocomposites withM-N-C active sites as anodes toward high-performance LIBs.
Acknowledgements
The authors are grateful for financial support from the “YoungTalent Fellowship” program and “the Fundamental Research Fundsfor the Central Universities”(2018JQ06) through South China Uni-versity of Technology, Grant (No. 51672086) through NSFC Com-mittee of China, Grant (No. 2017B030308005) through the Scienceand Technology Bureau of Guangdong Government, and the Crea-tive Research Initiative (CRI, 2014R1A3A2069102), the BK21 Plus(10Z20130011057), Science Research Centre (SRC,2016R1A5A1009405) and Young Researcher (2019R1C1C1006650)programs through the National Research Foundation (NRF) ofKorea.
C. Huang et al. / Materials Today Energy 14 (2019) 1003598
Appendix A. Supplementary data
Supplementary data to this article can be found online athttps://doi.org/10.1016/j.mtener.2019.100359.
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