Exploring Lithium Deficiency in Layered Oxide Cathode for Li‐Ion...

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Exploring Lithium Deficiency in Layered Oxide Cathode for Li-Ion Battery

Sung-Jin Cho,* Md-Jamal Uddin, Pankaj K. Alaboina, Sang Sub Han, Manjula I. Nandasiri, Yong Seok Choi, Enyuan Hu, Kyung-Wan Nam, Ashleigh M. Schwarz, Satish K. Nune, Jong Soo Cho, Kyu Hwan Oh, and Daiwon Choi

DOI: 10.1002/adsu.201700026

to develop high energy density LIBs. Efforts have been made to develop high energy anodes such as silicon-based mate-rials. However, high-energy cathode mate-rials need to be developed first since the LIBs are still cathode limited. For the inter-calation chemistry, the cathode capacity of LIBs can be improved by supplying more Li-ion; and it can be done in two ways—either put more Li-ion in the cathode or increase the voltage to extract more Li-ion. The first one requires new chemistry such as xLi2MnO3·(1−x)Li(Nix’MnyCoz)O2 (LMR-NMC). Unfortunately, the commer-cialization of such material is far from the reality because of the unsolved voltage fading issues.[1] The second approach is also challenging because of the significant amount of phase transition involved at high voltage.[2–4] For that reason, practical capacity of the existing layered cathode materials is still limited. For example, lay-ered NMC (LiNi1/3Mn1/3Co1/3O2) (space group R-3m) materials show a significant

amount of phase transformation and oxygen evolution when cycled either at high voltage (>4.3 V) or at high temperature (60 °C).[5–7] In such state, migration of the transition metal into the lithium layer takes place and NiO like phase or spinel phase

The ever-growing demand for high capacity cathode materials is on the rise since the futuristic applications are knocking on the door. Conventional approach to developing such cathode relies on the lithium-excess materials to operate the cathode at high voltage and extract more lithium-ion. Yet, they fail to satiate the needs because of their unresolved issues upon cycling such as, for lithium manganese-rich layered oxides—their voltage fading, and for as nickel-based layered oxides—the structural transition. Here, in contrast, lithium-deficient ratio is demonstrated as a new approach to attain high capacity at high voltage for layered oxide cathodes. Rapid and cost effective lithiation of a porous hydroxide precursor with lithium deficient ratio is acted as a driving force to partially convert the layered material to spinel phase yielding in a multiphase structure (MPS) cathode material. Upon cycling, MPS reveals structural stability at high voltage and high temperature and results in fast lithium-ion diffusion by providing a distinctive solid electrolyte interface (SEI) chemistry—MPS displays minimum lithium loss in SEI and forms a thinner SEI. MPS thus offers high energy and high power applica-tions and provides a new perspective compared to the conventional layered cathode materials denying the focus for lithium excess material.

Li-Ion Batteries

1. Introduction

To enable the advanced application of Li-ion batteries (LIBs), such as electric vehicle and ultraportable devices, it is mandatory

Prof. S.-J. Cho, M.-J. Uddin, P. K. Alaboina, Dr. J. S. ChoJoint School of Nanoscience and NanoengineeringNorth Carolina Agricultural and Technical State UniversityGreensboro, NC 27401, USAE-mail: [email protected]. S. Han, Dr. Y. S. Choi, Prof. K. H. OhDepartment of Materials Science EngineeringSeoul National UniversitySeoul 08826, Republic of KoreaDr. M. I. NandasiriImaging and Chemical Analysis LaboratoryDepartment of PhysicsMontana State UniversityBozeman, MT 59718, USADr. E. HuChemistry DivisionBrookhaven National LaboratoryUpton, NY 11973, USA

Prof. K.-W. NamDepartment of Energy Materials EngineeringDongguk UniversitySeoul 04620, Republic of KoreaA. M. SchwarzEnvironmental and Molecular Sciences LaboratoryPacific Northwest National LaboratoryRichland, WA 99352, USADr. S. K. Nune, Dr. D. ChoiEnergy and Environmental DivisionPacific Northwest National LaboratoryRichland, WA 99352, USA

The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/adsu.201700026.

Adv. Sustainable Syst. 2017, 1, 1700026

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© 2017 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim1700026 (2 of 10)


forms which are associated with a volume change due to the migration of dissimilar size of ions.[8–10] The repeated volume changes result in the pulverization of the active material and metal dissolution in the electrolyte. These dissolved metal ions participate in side reactions with the electrolyte and eventu-ally increase the solid electrolyte interface (SEI) layers on both cathode and anode sides. Overall, the cell impedance rises, and capacity fading occurs.[10–12] At the same time, transition metal (TM) cation migrations lead to the breaking of TMO bonds and eventually evolves significant amount of oxygen which proves a major risk in battery safety.

To prevent from such structural failure and to improve the structural stability of layered NMC material in such condition (up to 4.5 V, <60 °C), much work has been done. The most common approach is to use a thin, protective layer of inac-tive materials such as Al2O3, ZrO2, etc., that can suppress the side reactions with electrolyte.[13–19] Alternatively, doping with similar size of the cation is also obtained to stabilize the struc-ture.[20–24] Eventually, core–shell structures with different struc-tures of active material are also implemented to stabilize the surface which is particularly popular for Li-rich NMC mate-rial.[25–32] Later on, full concentration gradient cathodes are introduced.[33,34] Indeed, the degradation of the layered NMC (R-3m) material is strongly related to their structural transfor-mation. Using a structurally stable cathode material (such as spinel) will definitely help maintain the intrinsic structure. Nevertheless, most of the previous studies have focused on materials with a separately created artificial layer of a spinel (or other stable) phase on its outmost surface.[25,30,31,33–35] While they stabilize the structure up to a certain extent, under tre-mendous stress—such as a combination of high voltage and high temperature, or mechanical stress—these materials tend to fail, and new interfaces are generated. Thus, they lose their advantages provided by a stable coating and result in cathode degradation.

Herein we report a new strategy to develop an NMC cathode material using Li-less ratio to reach closer to the theoretical capacity by using them at high voltage. We have synthesized a multiphase structure cathode material (MPS) where both the layered structure and the spinel structure is present. The structural stability of the defect-spinel structure (Fd-3m) NMC mate-rial is significantly outperforming because of the pillaring structure in lithium layer. As such, spinel NMC can sustain up to a very high voltage of 4.5 V and even higher than 4.5 V. Also; the spinel structure contains 3D channels for the Li-ion to transfer during lithiation/delithiation,[36–39] which eventually improves the rate capability. Spinel NMC thus offers improved cycling stability as well as high rate capability. Yet, comparing to the layered structure, the spinel structure has less number of active sites for Li-ion since transition metal ions stay and act as pillars in the lithium layer. Hence, the specific capacity of Spinel NMC is lower than that of the lay-ered NMC. In this new concept, we have

synthesized a multiphase structure cathode material to exploit the benefit of both the spinel and layered structure, i.e., to gain high structural stability with improved rate performance while keeping the capacity high enough, we have used similar manu-facturing process as of current commercial production of NMC layered cathode materials, with the added advantage of signifi-cantly shorter calcination time that can save electrical cost by 37%. The prepared MPS cathode material showed high struc-tural, thermal, and electrochemical stability with high cycle life and improved rate performance.

2. Results and Discussion

Concept diagram of the multiphase structure is depicted in Figure 1. Using a lithium deficient ratio in the solid state lithi-ation of the spray-dried metal hydroxide precursor, and using a controlled sintering rate to avoid the oxidation of Ni2+ carefully, we have produced a multiphase structure of cathode material. The porous metal hydroxide precursor was synthesized using a simple spray-drying method (Figure S1, Supporting Infor-mation); the multiphase cathode material was synthesized by simply calcination from the precursor. During the solid state lithiation, some of the Ni2+ ions may have migrated to the lithium layer from the transitional metal ions layer in the well-known octahedral (TM)-tetrahedral-octahedral (Li) pathway.[40] The transition metal migration is partially imposed here by using a small deficiency of lithium (0.90–0.99). Also, the use of porous metal hydroxide precursor has helped in faster lithia-tion which significantly reduces the sintering time (we refer to the Experimental Section for details).

To compare the MPS cathode material, a commercial single-phase layered NMC (hereafter denoted as SPS) with the regular, over-stoichiometric lithium content was used as the control. A coprecipitated cathode material is used as SPS, instead of a spray dried one, since the spray dried cathode materials with regular lithium content usually do not show promising


Figure 1. Concept diagram of multiphase structure cathode materials: Li-deficiency resulted in the formation of additional spinel phase apart from the major layered phase.




electrochemical performance.[41–43] Hence, one of the best per-forming commercial cathode material, Umicore MX7 was used as the control.

The morphology of the commercial SPS and MPS, and their respective precursors after scanning electron microscopy (SEM) are shown in Figure 2. The coprecipitated nonporous SPS precursors have flake like primary particles and spherical sec-ondary particle morphology, shown in Figure 2a. The secondary particle sizes are roughly translated to the SPS cathode material after high-temperature sintering, Figure 2b. Spray-dried porous MPS precursors also have spherical morphology as depicted in Figure 2c and are retained after solid state sintering, as seen in Figure 2d. However, the porous nature of the MPS precursor resulted in a faster sintering compared to SPS. Calcination of the metal hydroxide precursor took only 467 min to obtain the final MPS cathode material.

Comparing to SPS and MPS, particle size distribution is quite different. SPS has a wide particle size distribution with an average particle size of 7 µm. For MPS, the average sec-ondary particles size is bigger (12 µm) than that of SPS since the largest secondary particles are of a similar size for both SPS and MPS. This is also in agreement with the Brunauer–Emmett–Teller (BET) analysis results where the MPS parti-cles showed low specific surface area than the SPS particles (Table S1, Supporting Information). BET results also show the pore volume and average pore size, both of which are reduced for MPS indicating that the solid state lithiation during sin-tering was enough to fill the nanoporous precursors.

Primary particles of both the SPS and MPS are mostly is in nanoscale (100–200 nm). The shape of the primary particles is, however, quite different in SPS and MPS. SPS primary par-ticles are mostly of spherical to an octahedral shape which is common for layered structure material. For MPS, in addition to the SPS-like primary particles, tetrahedron shaped primary par-ticles are also present which are most possibly from a separate

phase. The presence of these particles indi-cates that there may be multiple phases involved in the cathode material.

The energy dispersive X-ray spectros-copy (EDS) mapping result of both SPS and MPS, shown in Figure S2 (Supporting Information), reveals that the distribution of transition metals is homogeneous. The compositional information obtained from the EDS are shown in Table S2 (Supporting Information). Although the composition data from EDS are localized, they still reveal the Ni:Mn:Co ratio of both SPS and MPS, which is close to the theoretical ratio of 1:1:1. Induc-tively coupled plasma optical emission spec-troscopy (ICP-OES) results also confirm the Ni:Mn:Co ratio to be close to the theoretical ratio of 1:1:1 for both SPS and MPS as well as for the metal hydroxide precursor (Table S3, Supporting Information). In addition to that, the ICP-OES results also show the lithium to the metal ratio for both SPS and MPS. SPS has a Li:M ratio of 1.089 which is fairly common for commercial NMC materials.

The Li:M ratio of MPS is 0.966, which reveals that lithium defi-cient nature of MPS cathode material.

To explore the structure of the MPS, synchrotron X-ray dif-fraction (SR-XRD) of SPS and MPS powder materials were car-ried out. Figure 3a,b shows the SR-XRD results of both these powders and their fitted results. For the SPS, as shown in Figure 3a, all the peaks could be indexed by the R-3m hexagonal phase, indicating the formation of a single phase. Thus, Riet-veld refinement was carried out on this sample to estimate the lattice parameter information. On the other hand, for MPS, Le Bail fitting was employed as no certain information is known about each phase’s (main phase and others) composition. By using Le Bail fitting, we can get accurate lattice parameter information without the knowledge of detailed structure. All peaks were indexed for MPS sample by using a hexagonal phase with R-3m space group and a secondary spinel phase with Fd-3m space group, as shown in Figure 3b. However, the actual amount of spinel phase could not be determined because of the nature of Le Bail fitting.

For the R-3m hexagonal phase in MPS, a = 2.8683 (1) and c = 14.2739(4), while for SPS, a = 2.8641 (1) Å and c = 14.2592 (1) Å. The hexagonal phase in MPS has larger lattice parameter than that of SPS. This correlates with the fact that MPS has a secondary spinel phase with a relatively small lattice parameter of 8.1430 Å. This small spinel lattice parameter is very likely due to the presence of cations with small radii such as Co3+ or Ni2+, the loss of which can increase the lattice para-meter as in the case of hexagonal phase of MPS. It is possible that some Ni or Co ion migrate from the main hexagonal phase in MPS to form the spinel phase, leading to the multi-phase structure in cathode material.[44,45]

Raman spectra can give useful and very accurate informa-tion from a localized region. Thus, Raman spectra of SPS and MPS powders were analyzed and reported in Figure 3c,d. For SPS, the two significant Raman peaks appeared around


Figure 2. SEM images of the a) SPS precursor at 5000×, b) SPS cathode material at 5000×, c) MPS precursor at 4000×, and d) MPS cathode material at 5000×.




480.2 and 582.0 cm−1 in the spectra as shown in Figure 3c, pos-sibly belongs to the bending Eg and stretching A1g modes of the oxygen vibration. In MPS, Figure 3d, in addition to the same peaks as of SPS, there is also a small peak around 634.6 cm−1 which is the characteristic peak for the spinel phase, associated with MnO bonding.

To further confirm the phases present in the cathode mate-rials, transmission electron microscopy (TEM) with selected area electron diffraction (SAED) was employed. Line scan-ning on the high angle annular dark field (HAADF)—scan-ning transmission electron microscopy (STEM) image of SPS showed almost no variation of the transition metals (Figure S3, Supporting Information). A quick line scanning on the HAADF-STEM image of MPS revealed fluctuation in cobalt content, as appeared in differential contrast in Figure S4a (Sup-porting Information). Analyzing further, the fluctuation of the cobalt content is also found even in a single particle, as shown in Figure S4b (Supporting Information). Different contrast region contains different amount of cobalt. A high cobalt con-tent of up to 24.27% was found, where the nominal amount is ≈10%.

The variation of local cobalt content demands for more local-ized study. As such, SAED has been employed in different regions of SPS and MPS with different contrast. Figure 4a shows the Bright Field (BF) image of SPS. From three different regions, SAED has been performed, and all of them have

matched with hexagonal layered phase (R-3m space group). For MPS, different regions from the same grain revealed that the high cobalt zone is predominantly in spinel phase (Fd-3m space group), and the other zone is in hexagonal layered phase (R-3m space group).

With the presence of spinel phase in MPS, it is projected that the electrochemical performance of MPS will be improved. Figure 5 shows the electrochemical performance of SPS and MPS. Figure 5a shows the formation cycle performance of both SPS and MPS (0.1 C, 3.0–4.5 V). The charge/discharge capacity and Coulombic efficiency for both SPS and MPS is similar with MPS being slightly higher. Although spinel phase is present in MPS, the amount of spinel present in MPS is low enough to decrease the cathode capacity. Rather, because of high voltage operation (4.5 V), MPS has shown a better capacity due to the presence of stable spinel phase.

Figure 5b shows the rate capability of the SPS and MPS. The rate capability is improved for MPS due to the presence of spinel phase in MPS. Spinel phase contains 3D open pathway for Li-ion, which facilities easy and fast Li-ion diffusion to and from the MPS cathode material during lithiation/delithi-ation. Even though the average particle size of MPS is a little higher than SPS (Figure 2), the presence of spinel phase has improved the Li-ion diffusion rate and thereby improved the rate performance.[31,35] Rate capability of MPS is improved at lower rates and comparable up to 4C. However, with increasing


Figure 3. SR-XRD pattern and their fitted results for a) SPS and b) MPS powder, and Raman spectra of c) SPS and d) MPS powder showing major peak positions.





Figure 4. HRTEM results of SPS and MPS. a) BF image of SPS and the SAED pattern from three different zones and b) BF image of MPS and the SAED pattern from two different contrast areas.

Figure 5. Electrochemical Performance of MPS and SPS: a) formation cycle voltage plot, b) rate capability (up to 8C), 1C/1C cycle data c) at room temperature, and d) at 60 °C. e) EIS results after 1st and 250th cycle at room temperature, and f) DSC Results of the SPS and MPS cathode after charging to 4.5 V.




rate (such as 8C), the average particle size effect plays a major role, and the benefit provided by the spinel phase become less important. Because of the higher surface area, Li-ion dif-fusion is improved, and SPS showed higher performance at higher rate. Nevertheless, the presence of spinel phase in MPS improved the rate performance.

Figure 5c shows the galvanic charge/discharge cycling result of SPS and MPS over 250 cycles (1C/1C cycling, 3.0–4.5 V). In the first cycle, the capacity of SPS and MPS is similar; SPS is having 177.64 mAh g−1, and MPS 177.67 mAh g−1. After 250 cycles, the capacity of MPS is 132.92 mAh g−1, whereas, for SPS, the capacity is 127.31 mAh g−1. MPS could retain a 74.81% of its initial capacity after 250 cycles, whereas SPS capacity retention was 71.67%. Figure 5d shows the high temperature (60 °C) charge/discharge cycling (1C/1C cycling, 3.0–4.5 V) results. Capacity after the first cycle is 183.2 mAh g−1 for SPS. The increased capacity comparing to room temperature cycling is due to the increase in Li-ion diffusion rate resulting from the temperature rise. The first cycle capacity of MPS is increased even more (191.54 mAh g−1). In addition to the temperature effect—which effectively increased the diffusion rate—the pres-ence of spinel phase has also aided in the Li-ion diffusion in MPS. After 50 cycles of charge/discharge, the capacity retention of MPS is 93.48%, whereas, for SPS, the retention has dropped to 79.12%. The high capacity retention of MPS is attributed to its high structural and thermal stability from the spinel phase. During prolonged cycling and continuous charge/discharge, the cathode particles usually break down and degrade to other phases. Hence, the capacity fading appears, as found for SPS. Then again, because the spinel phase is more stable, structural transformation is less for MPS. Thus, MPS is structurally, elec-trochemically, and thermally more stable.

The excellent cycle life and good rate capability of MPS are also related to its interfacial charge-transfer process and Li-ion diffusion kinetics in the cathode material. Figure 5e showed the electrochemical impedance spectroscopy (EIS) of aged cells (250 cycles for 1C/1C at room temperature). EIS result showed that the surface resistance of SPS and MPS is similar after 1C/1C cycling for 250 cycles. The major difference in EIS result of SPS and MPS is in their charge transfer resistance (Rct). Rct of SPS is more than 3 times higher than that of MPS. This is again attributed to the spinel phase on the MPS material sur-face. The 3D pathways in the MPS Fd-3m spinel phase readily transfer the Li-ion and thus show significantly less amount of resistance in charge transfer layer.

Thermal stability of Ni-based cathodes is of crucial impor-tance. Substantial amount of oxygen generation occurs in Ni-based cathodes when lithium deintercalation results at higher temperatures. For this reason, we have studied the thermal sta-bility of both SPS and MPS using differential scanning calo-rimetry (DSC), as illustrated in Figure 5f. Both SPS and MPS showed similar degradation temperature for oxygen evolu-tion. The amount of heat generation is directly representative of amounts of oxygen liberated from the lattice. For SPS, the amount of heat generated was 130.8 J g−1, whereas, for MPS, the total heat is 84.0 J g−1 which is 36% less than SPS. This result shows that the MPS shows improved structural and thermal stability which is again attributed to the structurally stable spinel phase in MPS.

Aged cathodes (1C/1C for 250 cycles at room temperature) were studied with XRD for SPS and MPS in Figure 6a,b, respectively. It is well-known that structural degradation occurs during cycling, and it is enhanced by unwanted side reactions from the electrolyte. Consequently, we found that both the SPS and the MPS aged cathodes show the formation of some other phases (peaks from 2θ = 20° to 35°). The amount of phase transition from layered phase (R-3m) to other phases by aging can be quantified by comparing the XRD peak intensity ratio of the reflections from (003) and the reflection from (104). In the fresh SPS cathode, the (003)/(104) intensity ratio is found to be 1.33. However, after aging the ratio is changed to 1.90. On the other hand, the (003)/(104) intensity ratio for MPS is almost unchanged by aging of 250 cycles; the ratio for fresh MPS cathode is 1.66 and for the aged one is 1.70. Thus, the change in (003)/(104) intensity ratio for SPS is 43%, comparing to 2% in MPS. This clearly indicates that the layered phase in MPS is more stable than the same phase in SPS. The possible reason is that the spinel phase in MPS has stabilized the cathode struc-ture and thereby suppressed the structural transition.

Concerning the surface chemical nature of SPS and MPS cathodes due to aging, X-ray photoelectron spectroscopy (XPS) analysis has been carried out on fresh and aged cathodes after 50 cycles of charge/discharge at 60 °C. XPS data were collected from the surface of the cathodes as well as after sputtering for 5, 15, 30, and 60 min with 2 keV Ar+ ion beam to obtain the chemical information at different depths of the cathodes. Figure S5a,b (Supporting Information) shows the normalized Co 2p, Mn 2p, and Ni 2p core level XPS spectra of fresh SPS and MPS cathodes, respectively, after 60 min sputtering. It is found that the oxidation states of Co, Mn, and Ni in both the SPS and MPS are similar for fresh cathodes; Co is predomi-nantly in 2+, Mn is in a combination of 2+ and 3+, and Ni is also in combination of 2+ and 3+ oxidation states. In Ni 2p spectra, the F Auger peak (coming from the PVDF binder in cathodes) is also overlapped with the Ni2+ and Ni3+ peaks, so it is difficult to obtain the quantitative information. After the aging at 60 °C, Co2p, and Mn 2p spectra did not show any sig-nificant change in their spectra, as shown in Figure S5b (Sup-porting Information). However, there is a change in Ni 2p spectra of SPS cathode relative to that of MPS cathode. Both spectra (Ni2p of SPS and MPS) contain a broad major peak indicating the presence of Ni2+ and Ni3+ peaks with overlapped F Auger peaks. According to previous studies.[31,35,46–48] the for-mation of Ni2+ indicates structural instability of the SPS, which is associated with the formation of Ni2+ species from side reac-tions with the electrolytes and spontaneous reduction. In our results, the SPS cathode may have more Ni2+ compared with MPS cathode. However, it is not feasible to confirm it quanti-tatively due to the presence of F Auger peaks in Ni 2p spectra.

It is also important to explore the role and nature of fluo-rine in these cathodes using XPS. Figure 6c,d shows the deconvoluted F 1s core level XPS spectra of fresh and cycled cathodes on the surface and at a 60 nm depth. Analyzing the fresh cathodes before sputtering, Figure 6c, the primary source of fluorine is PVDF. There is also a small contribution from LiF, which is likely to be a result of localized interaction between lithium and fluorine in PVDF on the cathode.[49] The removal of F from PVDF due to the X-ray beam damage has





been observed before.[50] In MPS fresh cathode, fluorine has interacted with other species as well and formed MF and CF in addition to LiF. The amount of LiF is lower in MPS than SPS. This is probably because MPS has a deficiency of lithium, which may have caused the F− to interact with the other spe-cies. With the sputtering depth increased, MF and CF also appear on SPS with the decrease in the amount of PVDF for both samples. The increased fluorination of Li and metals and the presence of C–F with the Ar+ sputtering can be associated with the differential sputtering of F in PVDF due to the high energy Ar+ ion beam.[51,52]

The F 1s spectra of aged samples cycled at 60 °C show the presence of same species on the surface along with some PFO due to the decomposition of electrolyte salt (LiPF6), as shown in Figure 6d. SPS still shows significantly higher LiF on the surface than MPS. The amounts of PFO/CF and MF are higher for MPS. This is probably because of the presence

of LiPF6 during cycling supplies extra F−, but since the amount of lithium is less in MPS, there is enough leftover F− to form more MF and CF than SPS.[53] With the Ar+ sputtering, the amount of PVDF and LiF decreased, while the amount of MF and PFO/CF increased as a combination of both actual composition of SEI layer at each depth and preferential sput-tering of F in PVDF. Even after sputtering, SPS shows more LiF than in MPS. The full XPS depth profiles of F species are shown in Figure S5c (Supporting Information). The change in the number of different F species in SPS and MPS after aging surely suggests that the SEI layer is different for SPS than MPS; probably the SEI is thinner for MPS because of the less amount of LiF. At the same time, the loss of lithium in SEI layer is less. Together, they have impacted on the electrochemical perfor-mance and has shown superior cycling stability at 60 °C.

The use of nanoporous precursor always holds a question about the mechanical strength of the cathode material. This is


Figure 6. XRD results of aged a) SPS and b) MPS cathode After cycling at 1C at 25 °C for 250 cycles, and F 1s core level XPS spectra of SPS fresh, MPS fresh, SPS aged, and MPS aged at 60 °C c) before and d) after sputtering.





particularly important from the industrial point of view since the cathode materials are commonly calendered to achieve maximum energy density. Consequently, we have studied the mechanical properties of MPS by indenting the cathode material using nanoindenter. Figure S6 (Supporting Informa-tion) summarizes the nanoindentation results of the SPS and MPS cathodes. As it shows, the change in both modulus and hardness is only trivial while comparing SPS with MPS. SPS showed a reduced modulus of 10.66 GPa, while MPS showed 10.70 GPa of the same. Also, the hardness of SPS is 113.50 GPa comparing to 112.31 GPa in MPS. The nanoindentation results reported are merely decisive, since they represent the average values of the modulus and hardness at different loadings (1, 5, and 10 mN). Thus, to see the effect of calendering on the electrodes, the cathodes were pressed at 5 and 8 tons, and their topography was observed. Figure S7a,b (Supporting Informa-tion) revealed the topography of SPS and MPS cathode mate-rials, respectively after pressing at 8 tons. After pressing, the SPS particles were broken as shown in Figure S7a (Supporting Information). Comparatively, MPS particles did not break much as found in Figure S7b (Supporting Information). This is also observed while pressing the cathode materials with 5 tons of pressure. However, when pressed, both SPS and MPS reached the same thickness, i.e., same target density (Table S4, Sup-porting Information). Thus, the strength of MPS is comparably higher than SPS and is suitable for calendering.

3. Conclusion

In conclusion, we have synthesized multiphase lithium defi-cient cathode materials for lithium-ion batteries in a proce-dure readily transferrable to the industrial scale. We have used a shorter calcination time that can significantly reduce the overall electrical cost of the cathode material to more than 35%. SR-XRD and SAED show that SPS contains only layered phase (space group R-3m), whereas MPS contain both the layered and spinel phase (space group Fd-3m). Because of the multiphase structure, MPS is more stable than SPS at high voltage and high temperature. After cycling, MPS showed significantly reduced phase transition compared to SPS. Capacity retention after 250 cycles at room temperature was 75% for MPS, com-paring to 72% in SPS. MPS also showed high rate capability up to 8C, comparing to SPS. At high temperature (60 °C) cycling of 50 cycles, MPS showed high capacity retention of 93.5% comparing to 79.1% with SPS. The high performance of MPS is attributed to the multiphase structure which is structurally more stable at higher voltages.

4. Experimental SectionMaterials and Methods: Nickel hydroxide, cobalt hydroxide, citric

acid, (ACROS Organics), manganese oxide (Strem Chemical, Inc.), and lithium carbonate were all collected as an analytical grade. The commercial polypropylene separator, Celgard 2400 (denoted as Celgard), was provided by Celgard company. The thickness of the Celgard separator was ≈25 µm. The electrolyte, 1 m lithium hexafluorophosphate (LiPF6) in 1:1 (v/v) ratio of ethylene carbonate and ethylene methylene carbonate was collected from Panax Etec. (Korea) and used as it is. For

the commercial NMC cathode material (SPS), NMC (1/1/1) cathode material was collected from Umicore (MX7h) and was used as it is. All other chemicals were of analytical grade and used as received.

The MPS cathode material was synthesized by calcination of the metal hydroxide precursor. The metal hydroxide precursor was first synthesized by a spray-drying method. Transition metal oxide/hydroxides (Ni(OH)2, Co(OH)2, Mn3O4) and a chelating agent (citric acid) were thoroughly mixed via a planetary ball mill (MSE Supplies). The ball to powder ratio was maintained to be 10:1 and a wet milling was employed. The milling time was 2 h with an rpm of 800. The resulting slurry was sprayed in a spray dryer using a spray nozzle with an inlet temperature of 200 °C. The outlet temperature was kept at 100 °C. The resulting hydroxide precursor was sieved for micrometer-sized particles. The metal hydroxide precursor is then mixed with Li2CO3 and was calcined for 467 min to obtain the final NMC (1/1/1) MPS cathode material.

Physicochemical Characterization: The morphology of the precursor, SPS, and MPS was observed using a field emission SEM (Carl Zeiss Auriga-BU FIB FESEM). During the SEM observation, EDS was also employed to obtain elemental composition of the material. To obtain the compositional information of the metals including lithium, ICP-OES was employed using a Varian 710-ES ICP spectrometer. To obtain bonding information and to see any possible difference between SPS and MPS, Raman spectroscopy was utilized. A Horiba Raman confocal microscope was used to analyze the Raman spectra using a 532 nm laser to observe any vibrational mode of SPS and MPS. SR-XRD was utilized using the National Synchrotron Light Source at Brookhaven National Lab. The characteristic diffraction patterns were used to identify the materials and different phases. The diffraction patterns were further studied by refinement to confirm multiple phases and to estimate corresponding lattice parameters.

TEM (FEI Tecnai F20) was employed to see the particle topography and the compositional information (using EDS). In addition to that, SAED study was performed on different regions of the material to check for various phases. XRD technique was employed on the aged and fresh cathodes to see the effect of electrochemical cycling on the structure of the material. A Rigaku Miniflex benchtop diffractometer was used for XRD study using a Cu Kα radiation of 1.54 056 Å. XPS was performed using a Kratos Axis Ultra DLD spectrometer to obtain the chemical information. Al Kα monochromatic X-ray source (1486.6 eV) was used and the X-ray source was operated at 130 W. XPS data were collected on the fresh and aged cathodes as well as after sputtering for 5, 15, 30, and 60 min with 2 KeV Ar+ ion beam to obtain the chemical information at different depths of the cathodes.

The surface area of the materials and the porosity information was obtained using the BET technique. A Micromeritics ASAP-2020 (Nitrogen sorption at 77 K) analyzer was used for surface area and porosity measurement. The samples were dried in at 80 °C for 4 h in vacuum to remove any residual moisture and then transferred to the analyzer for BET analysis. Nitrogen adsorption isotherms measurement was carried out in a relative pressure range from 0.04 to 0.25.

The mechanical properties of the SPS and MPS were studied by indenting the cathode material using a Hysitron TI 950 Triboindenter nanoindenter. A Berkovich probe was used to indent the cathode materials carefully. The samples were mounted on an epoxy stub, and a variety of loads (1, 5, or 10 mN) was applied for 2 s, with 5 s loading and 5 s unloading time. From the load versus displacement curve and the indentation area, the reduced Young’s modulus and the hardness of the cathode materials were determined. To see the effect of calendering on the cathodes, they were pressed using a Dake B-10 press and their topography was observed using SEM. The electrodes were placed between two cylindrical (1 in. diameter × 1 in. length) steel bars and were kept under a pressure of 5 or 8 tons for 1 min.

The thermal stability of the both SPS and MPS cathode were studied using DSC. A Q-200 DSC from TA Instruments was utilized. Both cathodes were heated up to 400 °C at a heating rate of 5 °C min−1 in a nitrogen environment.

Electrochemical Characterization: The electrochemical performance of the SPS and MPS cathode were evaluated by assembling half





cells (2032 type coin cells). A 0.25 mm Li foil was used as a counter electrode. The PANAX electrolyte and the Celgard separator were used to assemble the half cells. The SPS and MPS cathode were prepared by mixing the cathode material (90%), carbon black (6%), and PVDF (4%) in N-methyl pyrrolidone (NMP) and coating the resulting slurry on an aluminum current collector foil, followed by drying in a vacuum oven at 120 °C for 12 h. A battery testing system (TOSCAT 3100, Toyo System Co., Japan) was used to measure the cycling and C-rate performance. Electrochemical cycling was performed at 1C/1C rate for up to 250 cycles at room temperature. Also, a 1C/1C rate cycling was carried out at high temperature (60 °C) for up to 50 cycles. by varying the C-rate from 1 to 8C in a voltage window of 3–4.5 V. AC impedance measurement of the cells was investigated using the EIS measurement by applying an AC voltage in the frequency range of 0.01 Hz to 100 kHz using a Biologic electrochemical workstation.

Supporting InformationSupporting Information is available from the Wiley Online Library or from the author.

AcknowledgementsThe authors gratefully acknowledge support from Joint School of Nanoscience and Nanoengineering, a member of Southeastern Nanotechnology Infrastructure Corridor (SENIC) and National Nanotechnology Coordinated Infrastructure (NNCI), which is supported by the National Science Foundation (No. ECCS-1542174).

Conflict of InterestThe authors declare no conflict of interest.

Keywordsheterostructures, high energy-density, lithium-deficiency, lithium-ion batteries, multiphase cathode

Received: February 28, 2017Revised: May 9, 2017

Published online: June 23, 2017

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