Copper Phosphate as a Cathode Material for Rechargeable LiBatteries and Its Electrochemical Reaction MechanismGuiming Zhong,† Jingyu Bai,†,‡ Paul N. Duchesne,§ Matthew J. McDonald,† Qi Li,† Xu Hou,†

Joel A. Tang,∥ Yu Wang,Δ Wengao Zhao,¶ Zhengliang Gong,¶ Peng Zhang,§ Riqiang Fu,∥

and Yong Yang*,†,¶

†State Key Lab of Physical Chemistry of Solid Surfaces, Collaborative Innovation Center of Chemistry for Energy Materials andDepartment of Chemistry and ¶School of Energy Research, Xiamen University, Xiamen, Fujian 361005, China‡Shanghai Institute of Space Power-Sources, No. 2965 Dongchuan Road, Shanghai 200245, China§Department of Chemistry, Dalhousie University, Halifax, Nova Scotia B3H4R2, Canada∥National High Magnetic Field Laboratory, 1800 East Paul Dirac Drive, Tallahassee, Florida 32310, United StatesΔShanghai Synchrotron Radiation Facility, Shanghai Institute of Applied Physics, Chinese Academy of Science, Shanghai 201204,China

*S Supporting Information

ABSTRACT: In the search for new cathode materials for rechargeable lithium batteries, conversion-type materials have greatpotential because of their ability to achieve high specific capacities via the full utilization of transition metal oxidation states. Here,we report for the first time that copper phosphate can be used as a novel high-capacity cathode for rechargeable Li batteries,capable of delivering a reversible capacity of 360 mAh/g with two discharge plateaus of 2.7 and 2.1 V at 400 mA/g. Theunderlying reaction involves the formation as well as the oxidation of metallic Cu. The solid-state NMR, in situ XAFS, HR-TEM,and XRD results clearly indicate that Cu can react with Li3PO4 to form copper phosphate and LixCuyPO4 during the chargingprocess, largely determining the reversibility of Cu3(PO4)2. This new reaction scheme provides a new venue to explorepolyanion-type compounds as high-capacity cathode materials with conversion reaction processes.

■ INTRODUCTION

Lithium ion batteries, a promising power system for hybrid andpure electric vehicles, have been widely used in various types ofportable devices. Those important applications have driven theglobal search for novel electrode materials with higher specificcapacities and energy densities.1−5 Certain promising anodematerials, such as Si and Sn, have shown a large capacity of upto 3000 mAh/g.6−9 On the other hand, commercializedintercalation-type cathode materials exhibit much lowercapacities, for example LiCoO2 (theoretical capacity: ∼ 274mAh/g) and LiFePO4 (theoretical capacity: ∼ 170 mAh/g).Thus, cathode materials have become a technical “bottleneck”in the further development of Li-ion batteries with a highenergy density.

Recently, many researchers have investigated severalmaterials, such as Li2MSiO4 and Cu2.33V4O11, which undergomultiple-electron reactions, leading to higher capacities forrechargeable Li batteries.10−16 However, the capacities of thesematerials have not yet been entirely realized. Poizot et al. foundthat Li2O can react with transition metals, M, in the chargingprocess, thus enabling nanosized transition metal oxide (MO)particles (e.g., cobalt oxides) to exhibit high capacities byutilizing the full valence states of the metal.17 Unfortunately,among the MxNy family of conversion materials (M = Fe, Co,Ni, Cu, etc.; N = N, O, F, P, S, etc.), only the MFx family has

Received: June 16, 2015Revised: July 22, 2015Published: July 24, 2015

Article

pubs.acs.org/cm

© 2015 American Chemical Society 5736 DOI: 10.1021/acs.chemmater.5b02290Chem. Mater. 2015, 27, 5736−5744

shown sufficiently high conversion voltages to make thempossible candidates for use as rechargeable Li batterycathodes.5,18−27 However, it is still important to explore newtypes of conversion-type cathode materials with differentchemistries.Many works have demonstrated that CuO can be used as an

electrode material for lithium ion batteries, showing a relativelylarge capacity and good cycling performance.17,28 Replacementof oxide ions by phosphate may increase the conversionvoltage, due to the inductive effect. Therefore, in this work, werevisit copper phosphate, a well-known cathode for primarybatteries.29,30 To the best of our knowledge, this material hasnever been reported in the literature as a rechargeable cathodematerial. We observed for the first time that, by utilizing theCu2+/Cu0 redox reaction, a carbon-coated amorphous copperphosphate cathode was capable of delivering a reversiblecapacity larger than 360 mAh/g at two discharge plateaus of 2.7and 2.1 V with a high current density of 400 mA/g within 10cycles. Additionally, both amorphous and crystalline materialswere studied, showing that the material with an amorphousstate delivered better electrochemical performance. This maybe due to a more uniform distribution of reduction products.Furthermore, we employed several techniques to understand

the detailed reversible electrochemical reaction mechanism forthis new type of rechargeable material. X-ray diffraction (XRD)and high-resolution transmission electron microscopy (HR-TEM) were used to observe the morphology and fine structureof the material; high-resolution solid-state 7Li and 31P magicangle spinning (MAS) nuclear magnetic resonance (NMR)spectroscopy was used to track the local structure and phasechanges in the material;31−33 while X-ray absorption finestructure (XAFS) spectroscopy34,35 was implemented toinvestigate the oxidation state changes of Cu.

■ EXPERIMENTAL SECTIONMaterials. Crystalline Cu3(PO4)2 powders were prepared by the

solid-state reaction of stoichiometric amounts of (NH4)2HPO4 andCuO at 900 °C. The amorphous Cu3(PO4)2 were prepared throughprecipitation of (NH4)2HPO4 and Cu(NO3)2 aqueous solutions,followed by calcining at 500 °C. High-energy ball milling ofCu3(PO4)2 and acetylene black was utilized to produce the carbon-coated material Cu3(PO4)2/C (90 wt % Cu3(PO4)2). The electrodeswere then fabricated using 80 wt % Cu3(PO4)2/C, 10 wt % acetyleneblack, and 10 wt % PVDF, with NMP as a solvent. The loading of theactive material was around 2 mg/cm2. The electrochemical perform-ances of the cathode materials were examined using CR2025 coin-typecells. These cells were assembled in an argon-filled glovebox andincluded the prepared electrode as the cathode, lithium metal as theanode, 1 M LiPF6 in ethylene (EC)/dimethyl carbonate (DMC) as theelectrolyte, and Celgard 2300 as the separator. After the cells hadfinished cycling, they were disassembled, and the electrodes werewashed using DMC in an argon-filled glovebox. Finally, the electrodeswere dried at room temperature and used for the ex situ study.XRD and TEM. The XRD patterns of samples were collected using

a Panalytical X-pert diffractometer (PANalytical, Netherlands)equipped with a Cu Kα radiation source and operating at 40 kV and30 mA. In order to prevent oxidation of the samples, electrodes werecovered with Kapton brand tape in the glovebox prior to analysis. TheKapton tape generated a broad but weak peak between 10° and 35°,but there was no effect on signal collection from the electrodes. TEMand high resolution TEM imagery was collected with a Tecnai F20 STTEM operating at 200 kV. The samples were loaded on copper gridsin an argon-filled glovebox, sealed in an argon-filled transfer holder,and then transferred into the TEM column.XAFS. A standard CR2025 coin cell with small holes in each side

was used to perform the in situ XAFS measurements. Once the in situ

cell had been discharged (or charged) to a certain state, charging washalted, and XAFS measurements were immediately performed. TheXAFS data at the Cu K-edge were collected in transmission mode atroom temperature, using ion chamber detectors at beamline BL14W1of the Shanghai Synchrotron Radiation Facility (SSRF) and a Si(111)double-crystal monochromator. The focal spot size at the position ofthe sample was ∼0.25 mm. The monochromator was calibrated toreject higher harmonics of the selected wavelength (harmonic content<10−4), and data was collected over a range of energies, from 200 eVbelow to 800 eV above the Cu K-edge (8979 eV). The incidentphoton energy was calibrated using a standard Cu metal foil just priorto data collection in all in situ XAFS measurements. Processing andfitting of the XAFS data were performed using WinXAS software.36

Solid-State NMR. All NMR experiments were performed in a 9.4T magnetic field with a Bruker Avance III spectrometer. The 7Li NMRspectra were recorded using 1.3 mm probe-heads at a magic-anglespinning speed of 60 kHz with a Hahn-echo pulse sequence (90° pulse− τ − 180° pulse − τ). A recycle delay of 10 s and a 90° pulse lengthof 1.3 μs were used for 7Li. A stepwise spin−echo pulse sequence witha recycle delay of 0.01 s and a 90° pulse length pulse width of 1.6 μswas applied to acquire 31P NMR spectra. The irradiation frequencywas varied with a step size of 0.15 MHz over the range of 1.8 MHz.The spin−echo mapping 31P NMR spectra were obtained by summingthe individual spectra at different irradiation frequencies. The chemicalshifts of 7Li and 31P were calibrated using LiCl powder (0 ppm) andadenosine diphosphate (1 ppm), respectively.

■ RESULTS

Electrochemical Performance of the Cu3(PO4)2/CCathode. The rechargeability of transition metal phosphateas a conversion cathode material was demonstrated first. Figure1a shows the discharge/charge profiles of the crystallineCu3(PO4)2/C material operated between 1.5 and 4.8 V at aspecific current of 10 mA/g. A capacity of up to 425 mAh/g(theoretical capacity: 423 mAh/g) was observed in the first

Figure 1. Electrochemical profiles of carbon coated crystalline (a) andamorphous (b) Cu3(PO4)2 materials under 10 and 400 mA/g,respectively. The inset figures are the cycling performance of thematerials under 400 mA/g.

Chemistry of Materials Article

DOI: 10.1021/acs.chemmater.5b02290Chem. Mater. 2015, 27, 5736−5744

5737

discharge, suggesting that the transfer of two electrons per Cuatom takes place and Cu2+ is completely reduced to Cu0 uponcompletion of the discharge. In the second cycle, the crystallineCu3(PO4)2 material displayed a reversible capacity of only 290mAh/g (with 4.1 electrons per formula unit) in the secondcycle and decayed quickly (Figure S1). Nevertheless, as shownin Figure 1b, the amorphous copper phosphate materialexhibited much better cycling performance, with a reversiblecapacity larger than 360 mAh/g at 400 mA/g between 1.5 and4.8 V (Figure 1b and Figure S1) within 10 cycles, meaning that5.1 electrons per Cu3(PO4)2 unit, or 1.7 electrons per Cu atom,are reversibly transferred. The capacity decayed to 220 mAh/gafter 30 cycles, which may be due to the contact problem ofreduction products and dissolution of Cu.37,38The bettercycling performance of the amorphous sample may be due tothe smaller particle sizes of discharge products, copper andlithium phosphate in the amorphous sample, providing betterphysical contact between Cu and Li3PO4, thus facilitating therecharge reaction. Therefore, the cycling performance of thematerials may be further improved through optimizing particlesize and morphology.28,39 It should be noted that there weretwo clearer discharge plateaus, at 2.7 and 2.1 V, in the seconddischarge process, which may be due to the morphologychanges. From the crystal structure of Cu3(PO4)2 (Figure 2), it

can be seen that there is no vacant site in the Cu3(PO4)2structure into which Li may insert, suggesting that a conversionreaction occurs even in the initial discharge process. Thus, thetwo discharge plateaus in the second discharge process may bedue to different conversion reactions. As a new type of materialfor rechargeable Li batteries, the reaction mechanism may giveus some important information. In order to understand thereactions of this new type of cathode, especially the twodischarge plateaus and the recharging process, the crystallinematerial was examined in more detail, because of its highcrystallinity and purity.Crystal Phase Evolution Detected Using XRD. XRD is a

very useful tool for monitoring crystalline phase changes duringelectrochemical reactions. Figure 3a depicts the XRD patternsof crystalline Cu3(PO4)2 and Cu3(PO4)2/C. A broader halfpeak width appeared in the XRD pattern of Cu3(PO4)2/C,indicating that smaller particles were formed after ball milling,which brought about the rechargeable property. Figure 3b andFigure S2 are the XRD patterns of the material during the firsttwo discharge/charge cycles. It can be clearly seen that theintensities of the XRD peaks corresponding to Cu3(PO4)2decreased, and a new and broad Bragg peak at 43.3° (attributedto Cu (111)) was detected when the x value in LixCu3(PO4)2reached 2 during the first discharge, indicative of thedecomposition of Cu3(PO4)2 and the formation of metallicCu during discharge. In the recharge process, the intensity ofthe 43.4° peak decreased, indicating that the copper was being

reoxidized during the charge process. The discharge curve ofthe electrode (Figure 1) supports the claim of the completeconversion of Cu3(PO4)2 to Li3PO4 and Cu metal when thedischarge process is finished. However, no other new Braggpeaks were visible during cycling other than that at 43.3°(Cu(111)), indicating that Li3PO4 and any other productsformed on recharging are amorphous and making it verydifficult to obtain further structural information regarding anypossible products and/or reaction intermediates from thepowder XRD patterns. Nevertheless, the ex situ XRD studyconfirmed that Cu3(PO4)2 was converted into copper duringthe discharge process, while the copper was reoxidized duringthe charge process.

Fine Structure and Differences between Amorphousand Crystalline Materials Observed with TEM. Figure 4shows TEM images of both amorphous and crystallinematerials at different stages (i.e., before discharge, after firstdischarge, and after first recharge). Figure 4a shows that thepristine amorphous (main image) material exhibits a largedegree of disorder, and the crystalline (inset) Cu3(PO4)2sample possesses high crystallinity. Figure 4b demonstratesthe submicron level structure of the discharged amorphousmaterial. This structure consisted of often-interconnectedparticles, each several hundred nm in diameter and with acarbon coating of varied thickness (typically 5−20 nm). Theselarger particles were in turn formed from fairly uniform Cunanoparticles, approximately 5 nm in diameter, and surround-ing crystallites of Li3PO4, which can be seen in the image asdark “dots” and lighter regions, respectively. The dischargedcrystalline samples shared this general structure, but theconstituent Cu nanoparticles exhibited a greater size as well

Figure 2. Crystal structure of Cu3(PO4)2 with the P-1 space group.

Figure 3. (a) XRD patterns of crystalline Cu3(PO4)2 and Cu3(PO4)2/C. (b) Ex situ XRD patterns of LixCu3(PO4)2 in different lithiationstates.

Chemistry of Materials Article

DOI: 10.1021/acs.chemmater.5b02290Chem. Mater. 2015, 27, 5736−5744

5738

as a wider size distribution, which are likely due to the largercontiguous Cu3(PO4)2 regions in the pristine material. Figure4c shows this structural model, in which the Li3PO4 and Curegions are labeled in the amorphous sample that underwent asingle discharge. The dark, circular areas correspond to Cu andthe lighter surroundings, Li3PO4. The Li3PO4 was in the formof small crystallites for both amorphous and crystalline samples,although it was difficult to determine if the average Li3PO4

crystallite size differed between these two. Finally, Figure 4dshows the crystalline sample after the first recharge. Areas of Cuand Cu3(PO4)2 are seen, demonstrating the incompletereaction of Cu with surrounding Li3PO4 and the presence ofsubstantially larger Cu nanoparticles compared to amorphoussamples, which is in accordance with the XAFS results (Figure8a). However, the TEM imagery of amorphous samples afterthe first recharge did not display any crystal planes ofCu3(PO4)2, indicating that the recharged material was highlyamorphous. This TEM data is consistent with the previousobservation: the degree of disorder of Cu3(PO4)2 crystals in thepristine material had an impact on the charge capacity of thesample. The material with higher crystallinity formed larger Cunanoparticles upon discharge and larger regions of Cu3(PO4)2upon recharge, presumably interfering with Li ion transportduring the reaction process and leaving “islands” of unreactedmaterial dispersed throughout the structure.

Local Structure Changes during Cycling Detected bySolid-State NMR. High-resolution solid-state NMR is apowerful method for understanding the local structure ofmaterials, especially those in amorphous states. Figure 5 showsthe mass-normalized 7Li MAS NMR spectra of crystallineCu3(PO4)2/C during the first discharge and second chargeprocesses. The inset is the enlarged view of selected states.Upon discharge, two peaks at −0.3 and −13 ppm wereobserved, as shown in Figure 5a (x = 0.3, 2.7 V). The peak at−13 ppm can be assigned to the Li close to Cu2+ (3d9), as theunpaired electrons from paramagnetic Cu2+ causes a large shiftdue to Fermi-contact interactions.31 The peak at −0.3 ppm wasdue to the Li located in a diamagnetic environment. In addition,it can be clearly seen that the line-width of the −0.3 ppm peakhas a trend of being reduced during discharge. Since thematerial was completely converted to Cu metal and Li3PO4, thepeak at −0.3 ppm in Figure 5a (x = 6.0) could be confidentlyassigned to Li3PO4. The −0.3 ppm peak with a relativelybroader line-width for other discharge states (x = 0.3, 1.0, 2.0)may indicate the formation of another diamagnetic environ-ment during the discharge process, such as Cu+ (3d10), withstrong evidence for this provided by XAFS results (Figure 8a).Figure 5b presents the 7Li NMR spectra during the charge

process in the second cycle. The decrease in the intensity of thepeak at −0.3 ppm indicated that the Li in Li3PO4 deintercalatedduring the charge process. As shown in the inset of Figure 5b,

Figure 4. (a) TEM image of pristine amorphous material and pristine crystalline material (inset). (b) Submicron structure of amorphous materialupon first discharge. (c) and (d) are nanoscale structures of amorphous and crystalline material upon first discharge and first recharge, respectively.

Chemistry of Materials Article

DOI: 10.1021/acs.chemmater.5b02290Chem. Mater. 2015, 27, 5736−5744

5739

when the material was charged to 2.7 V only one peak at −0.3ppm with decreased intensity and a slightly broader line-widthwas observed compared to the discharged sample, which maybe due to the formation of Cu+. Another peak at −13 ppmcould be seen when the voltage reached 3.3 V. As mentionedbefore, the 7Li NMR peak at −13 ppm was due to the Lilocated around the Cu2+. Thus, we can conclude that Cu wasoxidized to Cu+ when the voltage was below 2.7 V, while Cu2+

was formed afterward. However, the intensity of the peak at−13 ppm did not increase in the following charge process,indicating that another compound without Li formed. Thedischarge in the second cycle shows reaction processes similarto the first cycle, as shown in Figure S3. The intensity of thepeak at −13 ppm increased when the material discharged to 2.7V and decreased when the voltage was 2.0 V, indicating thatCu2+ and Li+ contained intermediates that primarily formedwhen the discharge voltage was larger than 2.7 V and convertedinto Li3PO4 when the discharge voltage was lower than 2.5 V.Figure 6 shows the mass-normalized 31P NMR spectra of

crystalline materials at different discharge states, giving moredetailed information regarding the reaction products. The 31Ppeaks between 1500 and 2000 ppm were assigned to crystallineCu3(PO4)2. The

31P resonance of paramagnetic Cu3(PO4)2 willshift under different temperatures, and although the roomtemperatures when testing different samples were the same, thereal temperatures of the samples in the high-spinning rotorsmay have been different, due to the unavoidable presence of Alpieces scraped from the electrodes. Differing amounts of Alcause different temperature rises under a spinning frequency of60 kHz, resulting in different chemical shifts observed forCu3(PO4)2. However, the same 31P shifts were obtained undera low spinning frequency of 25 kHz, as shown in Figure S4.Figure 6a depicts the 31P NMR spectra of the crystallinematerial when the x value in LixCu3(PO4)2 was less than 2. Inaddition to the peak assignable to crystalline Cu3(PO4)2, a peakat around 9 ppm, whose line-width at half height decreasedwith discharge, can be assigned to Li3PO4 and diamagneticspecies, which was in accordance with the 7Li NMR results.Therefore, we can conclude that the diamagnetic intermediateformed in the discharge process was LixCu(I)yPO4.

40 Thereduction in the line-width may have been due to the

decomposition of LixCu(I)yPO4 and the formation of Li3PO4.Additionally, a broad 31P peak at around 400 ppm wasobserved, which was due to the P close to Cu2+. The intensityof the broad peak at around 400 ppm increased with dischargewhen the x value was less than 2 and then decreased in thefollowing discharge process. When the x value reached 5, thepeak almost vanished, as shown in Figure 6b.Figure 7a shows the 31P NMR spectra of the material during

the charge process in the second cycle. It can be seen that the

broad peak at around 400 ppm was obtained again when thecharge voltage reached 3.3 V. Considering that almost the sametrend of the 7Li peak at −13 ppm and the broad31P peak ataround 400 ppm can be found during the first and secondcycles, it is reasonable to believe that the Li−Cu(II)−P−Ocompound formed during the discharge/charge processes. Thegeneral material formula LixCu(II)yPO4

41,42 will be used here.This LixCuy(II)PO4 formed mainly when the charging voltagewas lower than 3.3 V. However, another compound formed inthe following charge process. As shown in Figure 7a, a very

Figure 5. Mass-normalized 7Li solid-state NMR spectra of crystallineCu3(PO4)2/C during the first discharge (a) and second charge (b)processes. The values on the left axis refer to the x value inLixCu3(PO4)2. The inset figures are the enlarged views.

Figure 6. Mass-normalized 31P MAS NMR spectra of crystallineCu3(PO4)2/C during the first discharge process. The inset figure is anenlarged view.

Figure 7. Mass-normalized 31P MAS NMR spectra of crystallineCu3(PO4)2/C during the charge (a) and discharge (b) processes in thesecond cycle.

Chemistry of Materials Article

DOI: 10.1021/acs.chemmater.5b02290Chem. Mater. 2015, 27, 5736−5744

5740

broad 31P peak between 500 and 3000 ppm was observed whenthe material was charged to 4.0 V. The ex situ 7Li NMR spectraduring the charge process demonstrates that a compoundwithout Li was possibly formed when the material was chargedto 4.0 V. Additionally, the 31P shift of amorphous copperphosphate was observed to be between −500 and 4000 ppm(Figure S5). Therefore, the broad peak between 500 and 3000ppm can be assigned to the amorphous copper phosphate. Inother words, amorphous copper phosphate formed when thecharge voltage reached 4.0 V. Figure 7b is the 31P NMR spectraof carbon coated crystalline materials during the seconddischarge process. It can be seen clearly that the intensity ofthe signal from the amorphous copper phosphate decreasedwith discharge, and the peak from LixCu(II)yPO4 was reducedwhen the voltage was up to 2.0 V, which was consistent withthe 7Li NMR results. It should be noted that a recycle delay of0.01 s was applied for the acquisition of 31P NMR spectra,causing a suppression of the signal of the peak at 9 ppm fromLi3PO4 and LixCu(I)yPO4, which showed a very long spin−lattice relaxation time (T1), larger than 10 s. However, therelaxation times of the 31P peaks corresponding to copperphosphate and LixCu(II)yPO4 were much less than 1 ms,allowing a full relaxation when using a recycle delay of 0.01 s.The 31P NMR spectra at different states using a recycle delay of60 s are shown in Figure S6.The high-resolution solid-state 7Li, 31P MAS NMR spectra

thus demonstrate that two intermediate compounds, LixCu-(I)yPO4 and LixCu(II)yPO4, were formed during the firstdischarge process. When x was larger than 2, the conversionreaction of Cu3(PO4)2 and LixCu(I, II)yPO4 into Li3PO4 wasmainly occurring. The charge process of the materialunderwent three reaction processes. First, LixCu(I)yPO4formed, while the voltage was less than 2.7 V. Second, whenthe voltage was between 2.7 and 3.3 V, the formation of

LixCu(II)yPO4 dominated. Finally, when the voltage was above3.3 V, amorphous copper phosphate was the primary reactionproduct. In addition, the second discharge process looks likethe reverse of the charge process. LixCu(I)yPO4 and LixCu-(II)yPO4 were formed when the discharge voltage was largerthan 2.7 V, and LixCu(II)yPO4 started to convert into Li3PO4when the voltage was lower than 2.5 V.

Tracking Cu Oxidation States and Local StructureChanges Using in Situ XAFS. The XAFS study gives a directview of the oxidation of copper. Figure 8a shows first-derivativecurves of in situ X-ray absorption near-edge structure (XANES)spectra (black) and linear combination fitting curves (red),obtained using the first-derivative XANES curves of thecrystalline Cu3(PO4)2 electrode and fully dischargedLi6Cu3(PO4)2 as standards. It can be clearly seen that inaddition to the peaks located at 8979 eV (Cu) and 8984 eV(Cu2+), a signal was also observed at 8981.4 eV in thedischarge/charge spectra, which can be attributed to the Cu+

ion.43,44 Due to the fact that the valence state of copper inCu3(PO4)2 is Cu

2+, we can deduce that the Cu+ is present inthe intermediate product, LixCu(I)yPO4. These results agreewell with evidence from NMR testing. The first-derivativespectra and XANES spectra (Figure S7) reveal that theformation of Cu+ in the first discharge process occurred mainlybefore the x value in LixCu3(PO4)2 reached 0.5 and decreasedslowly until disappearing when x increased to 5.5. Combiningthese results with the observation of the continuouslyincreasing intensity of the 7Li peak at −0.3 ppm (Figure 5a)and 31P peak at 9 ppm (Figure 6a), we can tell that Li3PO4formed continuously during discharge and was the maindischarge product during the whole discharge process. Duringcharging, the amount of Cu+ was found to gradually increase asx decreased, demonstrating the continuous formation of a smallamount of LixCu(I)yPO4 during the charge process. Of course,

Figure 8. (a) The first derivative curves of in situ XANES spectra (black) and the linear combination fitting curves (red), obtained using the firstderivative XANES curves of Cu3(PO4)2 (x = 0) and Li6Cu3(PO4)2 as standards. The vertical axis (x) indicates the corresponding value inLixCu3(PO4)2. (b) Discharge/charge profiles and fitting results (coordination numbers and bond lengths) for Cu−O and Cu−Cu from the in situEXAFS data.

Chemistry of Materials Article

DOI: 10.1021/acs.chemmater.5b02290Chem. Mater. 2015, 27, 5736−5744

5741

the main reaction was the oxidation of Cu to Cu2+, which isshown clearly in Figures S7 and S8. Additionally, the peak at8979 eV was observed in the XANES spectrum ofLi0.5Cu3(PO4)2, indicating that Cu metal formed as soon asthe battery began discharging; thus, a complex reactionmechanism involving Cu metal formation could be anticipated.Additionally, the XAFS data shows that the local environmentof Cu changes during the discharge/charge process. Figure 8bgraphs the bond lengths of Cu−O and Cu−Cu, as well as theircoordination number (CN), during cycling. Both Cu−O andCu−Cu bond lengths remained relatively unchanged through-out this process; however, the CN of Cu−O decreased withdischarging, while that of Cu−Cu increased. On charging, thesetrends were reversed. The increasing CN of Cu−Odemonstrated that Cu was oxidized to a Cu−O containingcompound, in accordance with the NMR results. Furthermore,the decreasing CN of the Cu−Cu unit indicated that theparticle size of Cu decreased with charging, which was due tothe reaction of Li3PO4 and Cu to form LixCuyPO4. Of note isthe fact that the Cu−Cu CN for the discharged material wasabout 9, which was less than that of bulk Cu (12), indicatingthat Cu metal nanoparticles were formed. Furthermore, someCu−Cu bonding was retained even after the battery was fullycharged, indicating that some unreacted Cu was still presenteven at 4 V, which is consistent with the TEM results.The XAFS study confirmed the existence and change of Cu+

in the intermediates LixCu(I)yPO4 during the cycling ofcrystalline copper phosphate. Additionally, it revealed that aCu mass was formed upon discharge, and then was reoxidizedto Cu+ and Cu2+ during the charging process. The analysisresults from XAFS are highly consistent with those from NMR.

■ CONCLUSION

The Cu3(PO4)2 material has superior characteristics that makeit suitable, with some additional process tuning, as a cathodematerial for lithium ion batteries. Due to the inductive effect ofthe phosphate group, the Cu3(PO4)2 shows a higher workingvoltage of 2.7 and 2.1 V compared to CuO (1.5 V).Furthermore, it is possible to control/select the dischargevoltage by changing the transition metal and anion groupmaterials. Presently, the material can deliver a reversible high

capacity of larger than 300 mAh/g within 20 cycles at highcurrent density. The reversibility of the material is very sensitiveto the particle size; the as-prepared crystalline sample showsalmost no cycle performance. However, the ball-milled carboncoated sample, which has smaller particles, displays reversibility.In addition, the amorphous material has a better rechargeability.Considering the different morphologies of the dischargedcrystalline and amorphous samples as given by TEM results, wecan deduce that the rechargeability of Cu3(PO4)2 dependsstrongly on the particle sizes of the discharge products but notthat of the pristine samples (Figures S10 and S11). It isspeculated that the smaller particle size of the Cu and Li3PO4phases formed during discharge leads to a more uniformreaction product distribution and provides better contactbetween Cu and Li3PO4, thus facilitating the recharge reaction(oxidation of Cu). In addition, the high diffusity of copper ionsin the polyanion framework could be another factor that affectsthe reversibility of the conversion reaction in Cu3(PO4)2.

45

Although the capacity of the material currently decays to 220mAh/g within 30 cycles, it is likely to be improved afterintroducing morphology and size control.28,46

Several local structure characterization techniques such asXAFS, NMR, and XRD have been carried out to probe thecrystal and local structure changes in the nanostructuredCu3(PO4)2 during the discharge/charge processes and thus tounderstand these complex conversion reactions, as summarizedin Figure 9. In the first discharge process, small amounts ofLixCu(I)yPO4 were formed when the x value in LixCu3(PO4)2was less than 0.5, material that was gradually reduced to Cu andLi3PO4 in the following discharge process. The result wasstrongly supported by in situ XAFS study. However, this is onlya side reaction. Li3PO4, LixCu(II)yPO4, and Cu are the threemain products when the x value in LixCu3(PO4)2 is less than 2.Moreover, the reduction of both Cu3(PO4)2 and LixCu(II)yPO4

to Li3PO4 and Cu primarily occurred during subsequentdischarge processes. In the charge process, Li3PO4 and Cu wereconverted to LixCu(I)yPO4 when the voltage was less than 2.7V. Subsequently, LixCu(II)yPO4 was primarily formed until thevoltage reached 3.3 V, where amorphous copper phosphatestarted to generate. The second discharge curve showed twoclear plateaus at 2.7 and 2.1 V. However, the reactions in the

Figure 9. Reaction scheme of a crystalline Cu3(PO4)2/Li battery during the first two cycles. CP, LC(II)P, LC(I)P, and LP refer to copper phosphate,LixCu(II)yPO4, LixCu(I)PO4, and Li3PO4, respectively.

Chemistry of Materials Article

DOI: 10.1021/acs.chemmater.5b02290Chem. Mater. 2015, 27, 5736−5744

5742

second discharge process were similar to the first discharge.Amorphous copper phosphate converted to LixCu(II)yPO4,Li3PO4, and Cu before 2.7 V, and then LixCu(II)yPO4 andamorphous copper phosphate were reduced to Cu and Li3PO4.The discharge plateau at 2.7 V in the second cycle had largercapacity than that in the first cycle, which may be due tomorphology change after the first discharge. A surface withhigher specific area may form, facilitating the reaction. It is well-known that the transition metal oxides, MO, can be used asconversion electrodes for lithium ion batteries. However, thelow conversion voltages of MO have thus far been insufficientfor use as cathode materials. By contrast, polyanion compoundsdisplay a higher conversion voltage because of the strongerinductive effect of the polyanion groups. Our results indicatethat transition metals, M, such as Cu, and Li3PO4 can beconverted into LixMPO4, which holds important implicationsfor the further optimization of these materials and the design ofnew polyanion-type cathode materials with large specificcapacities.

■ ASSOCIATED CONTENT*S Supporting InformationThe Supporting Information is available free of charge on theACS Publications website at DOI: 10.1021/acs.chemma-ter.5b02290.

Detailed information for the whole set of XRD patterns,NMR spectra, TEM figures, XANE spectra, and fittingcurves of extended k3-weighted EXAFS spectra duringcycling of the material (PDF)

■ AUTHOR INFORMATIONCorresponding Author*E-mail: yyang@xmu.edu.cn.Author ContributionsG.Z. and J.B. contributed equally. G.Z., J.B., and Q.L. carriedout the materials preparation and electrochemical test. G.Z. andX.H. acquired the NMR data, G.Z., R.F., and J.A.T. analyzedthe NMR data. G.Z., W.Y., and Z.G. carried out the XAFSexperiments. G.Z., P.N.D, and P.Z. processed and analyzed theXAFS data. M.M. and W.Z. recorded and analyzed the TEMresults. Q.L. acquired and processed the XRD data. Y.Y. andG.Z. proposed the research, and Y.Y. attained the main financialsupport for the research.NotesThe authors declare no competing financial interest.

■ ACKNOWLEDGMENTSThe authors acknowledge the financial support of their researchfrom the National Natural Science Foundation of China(Grants No. 21233004, 21474138, 21428303, and 21321062)and National Basic Research Program of China (973 program,Grant No. 2011CB935903). In particular, R. Fu is indebted as aPCOSS fellow supported by the State Key Lab of PhysicalChemistry of Solid Surfaces, Xiamen University, China. P.Zhang thanks NSERC Canada for supporting the XAFSanalysis program. J. A. Tang and R. Fu are grateful for financialsupport from the User Collaboration Grants Program (UCGP)at the National High Magnetic Field Laboratory (NHMFL)which is supported by USA National Science Foundation(NSF) Cooperative Agreement No. DMR-1157490, the Stateof Florida, and the U.S. Department of Energy. Special thanks

are given for insightful comments made by Prof. Clare P. Greyabout the NMR results.

■ REFERENCES(1) Whittingham, M. S. Lithium batteries and cathode materials.Chem. Rev. 2004, 104, 4271−4302.(2) Armand, M.; Tarascon, J.-M. Building better batteries. Nature2008, 451, 652−657.(3) Bruce, P. G.; Scrosati, B.; Tarascon, J.-M. Nanomaterials forRechargeable Lithium Batteries. Angew. Chem., Int. Ed. 2008, 47,2930−2946.(4) Goodenough, J. B.; Park, K.-S. The Li-Ion Rechargeable Battery:A Perspective. J. Am. Chem. Soc. 2013, 135, 1167−1176.(5) Cabana, J.; Monconduit, L.; Larcher, D.; Palacin, M. R. Beyondintercalation-based Li-ion batteries: the state of the art and challengesof electrode materials reacting through conversion reactions. Adv.Mater. 2010, 22, E170−92.(6) Li, H.; Huang, X.; Chen, L.; Wu, Z.; Liang, Y. A high capacitynano Si composite anode material for lithium rechargeable batteries.Electrochem. Solid-State Lett. 1999, 2, 547−549.(7) Chen, H.; Xiao, Y.; Wang, L.; Yang, Y. Silicon nanowires coatedwith copper layer as anode materials for lithium-ion batteries. J. PowerSources 2011, 196, 6657−6662.(8) Chan, C. K.; Peng, H.; Liu, G.; McIlwrath, K.; Zhang, X. F.;Huggins, R. A.; Cui, Y. High-performance lithium battery anodes usingsilicon nanowires. Nat. Nanotechnol. 2008, 3, 31−35.(9) Ng, S. H.; Wang, J.; Wexler, D.; Konstantinov, K.; Guo, Z. P.; Liu,H. K. Highly Reversible Lithium Storage in Spheroidal Carbon-CoatedSilicon Nanocomposites as Anodes for Lithium-Ion Batteries. Angew.Chem., Int. Ed. 2006, 45, 6896−6899.(10) Melot, B. C.; Tarascon, J.-M. Design and Preparation ofMaterials for Advanced Electrochemical Storage. Acc. Chem. Res. 2013,46, 1226−1238.(11) Rangappa, D.; Murukanahally, K. D.; Tomai, T.; Unemoto, A.;Honma, I. Ultrathin nanosheets of Li2MSiO4(M= Fe, Mn) as high-capacity Li-ion battery electrode. Nano Lett. 2012, 12, 1146−1151.(12) Lv, D.; Bai, J.; Zhang, P.; Wu, S.; Li, Y.; Wen, W.; Jiang, Z.; Mi,J.; Zhu, Z.; Yang, Y. Understanding the High Capacity of Li2FeSiO4: Insitu XRD/XANES study combined with first-principles calculations.Chem. Mater. 2013, 25, 2014−2020.(13) Cheng, F.; Chen, J. Transition metal vanadium oxides andvanadate materials for lithium batteries. J. Mater. Chem. 2011, 21,9841−9848.(14) Morcrette, M.; Rozier, P.; Dupont, L.; Mugnier, E.; Sannier, L.;Galy, J.; Tarascon, J.-M. A reversible copper extrusion−insertionelectrode for rechargeable Li batteries. Nat. Mater. 2003, 2, 755−761.(15) Whittingham, M. S.; Song, Y.; Lutta, S.; Zavalij, P. Y.; Chernova,N. A. Some transition metal (oxy) phosphates and vanadium oxidesfor lithium batteries. J. Mater. Chem. 2005, 15, 3362−3379.(16) Gong, Z.; Yang, Y. Recent advances in the research ofpolyanion-type cathode materials for Li-ion batteries. Energy Environ.Sci. 2011, 4, 3223−3242.(17) Poizot, P.; Laruelle, S.; Grugeon, S.; Dupont, L.; Tarascon, J.Nano-sized transition-metal oxides as negative-electrode materials forlithium-ion batteries. Nature 2000, 407, 496−499.(18) Ma, D.-l.; Cao, Z.-y.; Wang, H.-g.; Huang, X.-l.; Wang, L.-m.;Zhang, X.-b. Three-dimensionally ordered macroporous FeF3 and itsin situ homogenous polymerization coating for high energy and powerdensity lithium ion batteries. Energy Environ. Sci. 2012, 5, 8538−8542.(19) Kim, S.-W.; Seo, D.-H.; Gwon, H.; Kim, J.; Kang, K. Fabricationof FeF3 Nanoflowers on CNT Branches and Their Application to HighPower Lithium Rechargeable Batteries. Adv. Mater. 2010, 22, 5260−5264.(20) Badway, F.; Mansour, A. N.; Pereira, N.; Al-Sharab, J. F.;Cosandey, F.; Plitz, I.; Amatucci, G. G. Structure and Electrochemistryof Copper Fluoride Nanocomposites Utilizing Mixed ConductingMatrices. Chem. Mater. 2007, 19, 4129−4141.

Chemistry of Materials Article

DOI: 10.1021/acs.chemmater.5b02290Chem. Mater. 2015, 27, 5736−5744

5743

(21) Badway, F.; Pereira, N.; Cosandey, F.; Amatucci, G. Carbon-Metal Fluoride Nanocomposites Structure and Electrochemistry ofFeF3:C. J. Electrochem. Soc. 2003, 150, A1209−A1218.(22) Yamakawa, N.; Jiang, M.; Grey, C. P. Investigation of theConversion Reaction Mechanisms for Binary Copper(II) Compoundsby Solid-State NMR Spectroscopy and X-ray Diffraction. Chem. Mater.2009, 21, 3162−3176.(23) Yamakawa, N.; Jiang, M.; Key, B.; Grey, C. P. Identifying theLocal Structures Formed during Lithiation of the Conversion Material,Iron Fluoride, in a Li Ion Battery: A Solid-State NMR, X-rayDiffraction, and Pair Distribution Function Analysis Study. J. Am.Chem. Soc. 2009, 131, 10525−10536.(24) Li, T.; Chen, Z. X.; Cao, Y. L.; Ai, X. P.; Yang, H. X. Transition-metal chlorides as conversion cathode materials for Li-ion batteries.Electrochim. Acta 2012, 68, 202−205.(25) Liu, J.-l.; Cui, W.-j.; Wang, C.-x.; Xia, Y.-y. Electrochemicalreaction of lithium with CoCl2 in nonaqueous electrolyte. Electrochem.Commun. 2011, 13, 269−271.(26) Amatucci, G. G.; Pereira, N. Fluoride based electrode materialsfor advanced energy storage devices. J. Fluorine Chem. 2007, 128, 243−262.(27) Zhang, W.; Ma, L.; Yue, H.; Yang, Y. Synthesis andcharacterization of in situ Fe2O3-coated FeF3 cathode materials forrechargeable lithium batteries. J. Mater. Chem. 2012, 22, 24769−24775.(28) Zheng, S.-F.; Hu, J.-S.; Zhong, L.-S.; Song, W.-G.; Wan, L.-J.;Guo, Y.-G. Introducing dual functional CNT networks into CuOnanomicrospheres toward superior electrode materials for lithium-ionbatteries. Chem. Mater. 2008, 20, 3617−3622.(29) Dey, A. N.; Needham; Holmes, R. W.; Mass, D. Metalphosphate and metal arsenate organic electrolyte cells. U.S. Patent3736184 A, 1973.(30) Caillaud, T.; Lecerf, A. Positive active material for an electric cellwith a non-aqueous electrolyte. U.S. Patent 4260668 A, 1981.(31) Grey, C. P.; Dupre,́ N. NMR studies of cathode materials forlithium-ion rechargeable batteries. Chem. Rev. 2004, 104, 4493−4512.(32) Blanc, F. d. r.; Leskes, M.; Grey, C. P. In Situ Solid-State NMRSpectroscopy of Electrochemical Cells: Batteries, Supercapacitors, andFuel Cells. Acc. Chem. Res. 2013, 46, 1952−1963.(33) Hu, Y.-Y.; Liu, Z.; Nam, K.-W.; Borkiewicz, O. J.; Cheng, J.;Hua, X.; Dunstan, M. T.; Yu, X.; Wiaderek, K. M.; Du, L.-S. Origin ofadditional capacities in metal oxide lithium-ion battery electrodes. Nat.Mater. 2013, 12, 1130−1136.(34) Zhang, W.; Duchesne, P. N.; Gong, Z.-L.; Wu, S.-Q.; Ma, L.;Jiang, Z.; Zhang, S.; Zhang, P.; Mi, J.-X.; Yang, Y. In SituElectrochemical XAFS Studies on an Iron Fluoride High-CapacityCathode Material for Rechargeable Lithium Batteries. J. Phys. Chem. C2013, 117, 11498−11505.(35) Song, M.-K.; Park, S.; Alamgir, F. M.; Cho, J.; Liu, M.Nanostructured electrodes for lithium-ion and lithium-air batteries: thelatest developments, challenges, and perspectives. Mater. Sci. Eng., R2011, 72, 203−252.(36) Ressler, T. WinXAS: a program for X-ray absorptionspectroscopy data analysis under MS-Windows. J. Synchrotron Radiat.1998, 5, 118−122.(37) Whitehead, A. H.; Schreiber, M. Current Collectors for PositiveElectrodes of Lithium-Based Batteries. J. Electrochem. Soc. 2005, 152,A2105−A2113.(38) Hua, X.; Robert, R.; Du, L.-S.; Wiaderek, K. M.; Leskes, M.;Chapman, K. W.; Chupas, P. J.; Grey, C. P. Comprehensive Study ofthe CuF2 Conversion Reaction Mechanism in a Lithium Ion Battery. J.Phys. Chem. C 2014, 118, 15169−15184.(39) Kang, W.; Shen, Q. The shape-controlled synthesis and novellithium storage mechanism of as-prepared CuC2O4·xH2O nanostruc-tures. J. Power Sources 2013, 238, 203−209.(40) Snyder, K.; Raguz, B.; Hoffbauer, W.; Glaum, R.; Ehrenberg, H.;Herklotz, M. Lithium Copper(I) Orthophosphates Li3‑xCuxPO4:Synthesis, Crystal Structures, and Electrochemical Properties. Z.Anorg. Allg. Chem. 2014, 640, 944−951.

(41) Upreti, S.; Chernova, N. A.; Xiao, J.; Miller, J. K.; Yakubovich,O. V.; Cabana, J.; Grey, C. P.; Chevrier, V. L.; Ceder, G.; Musfeldt, J.L.; Whittingham, M. S. Crystal Structure, Physical Properties, andElectrochemistry of Copper Substituted LiFePO4 Single Crystals.Chem. Mater. 2012, 24, 166−173.(42) Amine, K. Active material for lithium batteries. U.S. Patent6319632 B1, Nov. 20, 2001, 1999.(43) Lamberti, C.; Bordiga, S.; Bonino, F.; Prestipino, C.; Berlier, G.;Capello, L.; D’Acapito, F.; i Xamena, F. L.; Zecchina, A. Determinationof the oxidation and coordination state of copper on different Cu-based catalysts by XANES spectroscopy in situ or in operandoconditions. Phys. Chem. Chem. Phys. 2003, 5, 4502−4509.(44) Cui, Y.-H.; Xue, M.-Z.; Zhou, Y.-N.; Peng, S.-M.; Wang, X.-L.;Fu, Z.-W. The investigation on electrochemical reaction mechanism ofCuF2 thin film with lithium. Electrochim. Acta 2011, 56, 2328−2335.(45) Yu, H. C.; Ling, C.; Bhattacharya, J.; Thomas, J. C.; Thornton,K.; Van der Ven, A. Designing the next generation high capacitybattery electrodes. Energy Environ. Sci. 2014, 7, 1760−1768.(46) Gao, X. P.; Bao, J. L.; Pan, G. L.; Zhu, H. Y.; Huang, P. X.; Wu,F.; Song, D. Y. Preparation and Electrochemical Performance ofPolycrystalline and Single Crystalline CuO Nanorods as AnodeMaterials for Li Ion Battery. J. Phys. Chem. B 2004, 108, 5547−5551.

Chemistry of Materials Article

DOI: 10.1021/acs.chemmater.5b02290Chem. Mater. 2015, 27, 5736−5744

5744