An Unavoidable Challenge for Ni-rich Positive Electrode Materials for

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An Unavoidable Challenge for Ni-rich Positive Electrode Materials for Lithium-Ion Batteries Hongyang Li, Aaron Liu, Ning Zhang, Yiqiao Wang, Shuo Yin, Haohan Wu, and Jeff R. Dahn Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.9b02372 • Publication Date (Web): 12 Aug 2019 Downloaded from pubs.acs.org on August 12, 2019

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Chemistry of Materials

An Unavoidable Challenge for Ni-rich Positive Electrode Materials for Lithium-ion Batteries Hongyang Li a, Aaron Liu b, Ning Zhang a,c, Yiqiao Wang d, Shuo Yin d, Haohan Wu d,e and Jeff R. Dahn *,a a Physics

b

Department of Chemistry, Dalhousie University, Halifax, NS B3H 4R2, Canada c

d

and Atmosphere Science, Dalhousie University, Halifax, NS B3H 3J5, Canada

School of Metallurgy, Northeastern University, Shenyang 110819, China

Hunan Zoomwe Zhengyuan Advanced Material Trade Co., Ltd, Changsha 410000, China

e Hoffmann

Institute of Advanced Materials, Shenzhen Polytechnic, Shenzhen 518000, China

*corresponding author - [email protected]. Tel.: 001-902-494-2991; Fax: 001-902-494-5191

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Abstract LiNi1-x-yCoxAlyO2 (NCA) and LiNi1-x-yMnxCoyO2 (NMC) materials are widely used in electric vehicle and energy storage applications. Derived from LiNiO2 (LNO), NCA and NMC materials with various chemistries were developed by replacing Ni with different cations. Many studies on the failure mechanisms of NCA and NMC materials have attributed the cell degradation to the anisotropic volume change of particles. In this work, it is shown that for Ni-rich layered transition metal oxide materials, regardless of composition, the unit cell volumes change in an almost identical manner during delithiation. Half-cell cycling data collected from 26 sets of Nirich materials with different compositions allows a relationship between capacity retention and accessible capacity to be observed. This relationship can be correlated to the unit cell volume change during the lithiation-delithiation process.

This work suggests a universal failure

mechanism for Ni-rich positive electrode materials which must be overcome.


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Introduction Ni-rich layered lithium transition metal oxides are one of the most promising positive electrode materials for lithium ion batteries.1 yMnxCoyO2(NMC) and

Many derivatives of LiNiO2, such as LiNi1-x-

LiNi1-x-yCoxAlyO2 (NCA) have been successfully commercialized and are

of great interest to both industry and academia.2–4 For layered lithium transition metal oxide cathode materials, many failure mechanism studies attribute some portion of cell degradation to micro-cracks which form in the secondary particles causing impedance growth and active material loss.

5–11

As a result of cracking, electrolyte

infiltration into the particle interior can further worsen the cell performance. Such cracking issues were identified in the studies of NCA materials by Y. Makimura et al.9 It is believed that the secondary particle cracking issue is highly related to the anisotropic volume change during charge and discharge, which has been observed in both NCA and NMC materials.12,13 H. Ryu et al.14 showed that in the NMC series, the problem of micro-crack generation during cycling becomes more severe as the Ni content increases. However, a recent study by W. Li et al.15 show that for NMC materials, the lattice collapse is a universal phenomenon almost entirely dependent on Li utilization, and not Ni content. W. Li et al. point out that one reason for less crack generation in low Ni content NMC materials is the lower degree of delithiation for the same upper cutoff voltage. Particle cracking is widely suspected to be one cause of cathode material degradation and there are many efforts focusing on resolving the cracking issues.16–19

It is very important to

understand the unit cell volume change and how it affects the charge-discharge cycling performance of Ni-rich positive electrode materials. In this work, from in-situ X-ray diffraction



(XRD) measurements on a variety of Ni-rich positive electrode materials, including conventional NMC, NCA, and LiNiO2 materials, it is shown that regardless of the composition, the unit cell volume of LiNiO2 derivative cathode materials change in an almost identical manner with lithium content. From 26 sets of half–cell cycling data collected with different Ni-rich positive electrode materials, an interesting relationship between the capacity retention and cell accessible capacity was found. For materials with a 1st C/20 discharge capacity higher than ~210 mA/g, the capacity retention after 50 cycles dropped rapidly as the accessible capacity increased. This abrupt drop in capacity retention coincides for all materials with a unit cell volume contraction upon delithiation, which shows a sudden drop beyond ~76% of full delithiation or when x = 0.24 in LixMO2. It is tempting to speculate that there is an intrinsic correlation behind the loss in capacity retention and the more severe volume contraction as more capacity is accessed. Recently, Jung et al. 20 have shown that oxygen release from a variety of NMC materials occurs at the point where 81% (very close to 76%) of the lithium is removed. It is possible that this increase in oxygen evolution is related to fresh surfaces created by microcracking induced by volume contraction. This work will hopefully help the battery community understand one common failure mechanism of Ni-rich cathode materials and interpret data sensibly. Experimental Ni-rich positive electrode materials Metal hydroxide precursors used for synthesizing samples 1-16 (listed in Table S1) were acquired from Zoomwe Zhengyuan Advanced Material Co., Ltd (China); precursors used for synthesizing samples 17-21 were made at Dalhousie University using co-precipitation in a continuously stirred tank reactor (CSTR) (Brunswick Scientific/Eppendorf BioFlo 310). More 4 ACS Paragon Plus Environment

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details of precursor synthesis at Dalhousie University have been reported by J. Li et al..21 The experimental details of the lithiation (heating) step have been reported in references.13,22 The NCA material used for sample 22 was provided by Ecopro. Samples 23-26 were extracted as electrodes from dry (no electrolyte) pouch cells provided by reputable suppliers called vendor (samples 23-26). Scanning electron microscopy imaging (SEM) SEM imaging was conducted using a NanoScience Phenom Pro G2 Desktop Scanning Electron Microscope with a backscattered electron detector. Samples were prepared by mounting the powders onto adhesive carbon tape. The images of samples were taken with an accelerating voltage of 5 kV and a current of 0.6 nA. Powder X-ray diffraction (XRD) Powder X-ray diffraction (XRD) was conducted to study the structure of the materials using a Siemens D5000 diffractometer equipped with a Cu target X-ray tube and a diffracted beam monochromator. Diffraction patterns were collected in the scattering angle (2θ) range of 15–70o at 0.02o intervals with a dwell time of 3 s. A 1o divergence slit, 1o anti-scattering slit and 0.2 mm receiving slit were used for the measurements.

The collected XRD patterns were refined

(Rietveld method) using “Rietica”.23 Coin cells Synthesized positive electrode materials were handled with great care to avoid exposure to moist air. All samples after sintering were moved directly to an argon-filled glove box for storage. Small amounts of powder were taken from the glove box for electrode preparation. Electrode 5 ACS Paragon Plus Environment


preparation was done in the laboratory and took a maximum of three hours. Electrodes were transferred into the glove box after they were made. The positive electrode materials were mixed with polyvinylidene difluoride (PVDF) and Super-S carbon black having a mass ratio of 92:4:4 in N-methyl-2-pyrrolidone (NMP) to make slurry. Single side coated electrodes were made by casting the slurry onto aluminum foil with a 150 μm notch bar spreader. The electrodes were dried in oven at 120oC for 3 hours. Dried electrodes were then calendared at a pressure of 2000 atm. Coin cell electrodes were punched (1.2 cm in diameter) and further dried under vacuum at 120oC for 14 hours before coin cell fabrication. The loading of electrode material was 10-12 mg/cm2. The electrode fabrication procedure closely followed that described by Marks et al.24 Coin cells of samples 23-26 were punched from single side coated regions of electrodes taken from dry (no electrolyte) 402035-size pouch cells provided by Vendor 1. The dry pouch cells had been vacuum sealed in the dry room of Vendor 1 prior to shipping to Canada. EC/DEC electrolyte (1.0 M LiPF6 in EC:DEC (1:2 v/v)), FEC/DMC electrolyte (1.0 M LiPF6 in FEC:DMC (1:4 v/v)), and FEC/DTD-EC:DEC electrolyte (2 wt.% FEC and 1 wt.% DTD in EC/DEC electrolyte) were used for the half coin cell cycling tests. Different electrolytes were used as we learned over time which electrolyte was superior for these tests. The FEC:DMC 1:4 electrolyte is most preferred. Standard 2325 coin cells were assembled in an argon-filled glovebox. Each coin cell had a positive electrode and a Li foil negative electrode with two layers of separators (Celgard #2300) in between.

Galvanostatic charge/discharge cycling was

conducted with E-One Moli Energy Canada battery testing systems. Results and Discussion


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Figure 1 shows three representative SEM images of the positive electrode materials used in this work. Images are shown of LiNiO2 made starting with a Zoomwe Ni(OH)2 precursors (a), LiNi0.95Mg0.05O2 made starting with a Ni0.95Mg0.05(OH)2 precursor made at Dalhousie Univ. (b), and LiNi0.80Co0.15Al0.05O2 from Ecopro (c). As NMC and NCA materials were all developed on the basis of LiNiO2, the failure mechanism of LiNiO2 was first studied in this work. A previous study demonstrated that LiNiO2 goes through a kinetic hindrance region (KH), H1 phase to monoclinic phase (M) transition, single phase M region, M to H2 phase transition, and H2 to H3 phase transition as lithium is removed from the material.25 To investigate the impacts of the phase transitions, five sets of half cells were cycled for 50 cycles at a rate of ~C/5 with different lower cutoff voltages (LCV) and different upper cutoff voltages (UCV). Before and after the cycling, two cycles over a full voltage range (3 - 4.3 V) at a rate of C/20 were performed. FEC/DMC electrolyte (1.0 M LiPF6 in FEC:DMC (1:4 v/v)) was used as the electrolyte. Figure 2a shows the specific capacity as a function of cycle number, and Figure 2b shows the corresponding normalized capacity retention during cycling. Figure S1 shows the differential capacity vs. voltage (dQ/dV vs. V) of LixNiO2 with the KH region, and phase transition peaks marked. Figures 2c1-2c5 show dQ/dV vs. V for the first C/5 cycling of half cells in black. Figure 2c shows that cells with a full cycling range have all the KH, H1, M, H2 and H3 regions included, while cells cycled between 3.55-4.3V, 3.65-4.3V, 3-4.1V, and 3.55-4.1V avoid KH, the KH region and the H1-H2 transition, the H2-H3 transition, and both the KH region and the H2-H3 transition, respectively. Figure 2 shows that cells with an UCV restricted to 4.1 V maintained good capacity retention, while the capacities of cells cycled with an UCV of 4.3 V, regardless of the choice of LCV, lost capacity at a similar rapid rate during cycling. Figure 2 shows that the going through the KH 7 ACS Paragon Plus Environment


region and the H1-M-H2 phase transitions does not result in rapid capacity loss, while the H2-H3 phase transition, which happens between 4.1-4.3 V is highly detrimental to the LiNiO2 cathode material. Figures 2c1-2c5 show dQ/dV vs. V of 50th cycle in red. Cells cycled with an UCV of 4.3 V show more changes to the dQ/dV curve after cycling compared to before cycling as shown in Figures 2c1-2c3 compared to cells with an UCV of 4.1 V as shown in Figures 2c4 and 2c5. Figures 2d1-2d5 show dQ/dV vs. V of full voltage range C/20 cycles before and after cycling. For cells that went through the H2-H3 phase transitions during cycling, more polarization at top of discharge, and more reduction of the dQ/dV peaks at ~3.8 V can be observed (marked by blue circles).

By contrast, cells cycled with an UCV restricted to 4.1 V have much smaller

polarization development and the peaks in dQ/dV vs. V are well preserved after cycling, which indicates less material degradation. The detrimental impact of the H2-H3 phase transition is believed to be caused by the abrupt anisotropic unit cell volume contraction,25 which can cause stress and strain generation due to lattice mismatch. The results presented here are consistent with a study by C. Yoon et al.,26 which showed the cracking of LiNiO2 particles when cycled to 4.2V and 4.3V. In a previous work by the authors, it was shown that the phase transitions in LixNiO2 can be effectively suppressed by partially substituting Ni atoms with Al, Mg, or Mn atoms, which can randomly “deactivate” lithium sites and thus disturb the lithium orderings thought to be responsible for the transitions.13 The abrupt unit cell volume contraction due to the mismatch between H2 and H3 phases could be suppressed by the elimination of the H2-H3 phase transition, but a continuous rapid unit cell volume change (as a single phase) at high state of charge (SOC) was still observed.13 To probe the impact of this unit cell volume change, half coin cells made 8 ACS Paragon Plus Environment

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with LiNi0.95Mg0.05O2 (NMg9505), LiNi0.95Al0.05O2 (NA9505), and LiNi0.9Co0.05Al0.05O2 (NCA900505) cathode materials were cycled with different UCVs. Figure S2 shows the XRD patterns of LiNi0.95Mg0.05O2, LiNi0.95Al0.05O2, and LiNi0.9Co0.05Al0.05O2. All cells were filled with FEC/DMC electrolyte and cycled at 30oC with a rate of ~C/5 for 50 cycles. Before and after the cycling, two cycles at a rate of C/20 were performed. Figures 3a – 3c show the initial C/20 cycling dQ/dV vs. V of NiMg9505 cells, NA9505 cells, and NCA900505 cells, respectively. The unit cell volume change of each cathode material as a function of cell voltage is plotted in the same corresponding graph, and the data was replotted from previous work.13 Figure 3 shows that cells with an UCV of 4.4 V have a full dQ/dV peak at 4.2 V, which can be considered to be a feature reminiscent of the H2-H3 phase transition peak observed in LiNiO2. Accompanying the 4.2 V peak, there is a continuous sharp drop in unit cell volume. For cells cycled up to 4.3 V, NMg9505 cells have only about 3/4 of the 4.2 V dQ/dV peak accessed, while NA9505 and NCA900505 cells still have most of the 4.2 V dQ/dV peak accessed. For NA9505 and NCA900505 cells, the 3-4.1 V cycling completely excludes the 4.2 V dQ/dV peak. Figures 3d – 3i show the specific capacity and normalized capacity of half coin cells as a function of cycle number. For all of the three cathode materials, with an increase of UCV, more capacity can be accessed, larger volume changes occur and the capacity retention becomes worse. We plotted the 50-cycle capacity retention of the half coin cells including LiNiO2 as a function of their 1st C/20 discharge capacity in Figure 4a.

NMg9505 cells, NA9505 cells, and

NCA900505 cells are marked by circles, diamonds and triangles, and LiNiO2 (LNO) is marked by a magenta star. Figure 4a displays an interesting pattern: the blue circle marks a region in which cells show stable capacity retention as C/20 discharge capacity increases, and the red 9 ACS Paragon Plus Environment


circle marks a region in which the capacity retention drops quickly as more C/20 capacity can be accessed. To figure out if this pattern also holds for other cell chemistries, 26 sets of cells were made and their half-cell testing results listed in Table I were plotted in the same way shown in Figure 4b. The half-cell testing data was collected in the authors’ lab. In this table, 16 different cathode materials, including lab made materials and commercial materials, are included. Three types of electrolyte were used due to different original testing purposes. All the cells were cycled at 30oC with the same LCV of 3.0 V and different UCVs ranging from 4.1 V to 4.4 V. Similar to Figure 4a, Figure 4b demonstrates the same trend of capacity retention change with increase in accessible C/20 discharge capacity. This pattern shows that regardless of the cell chemistry and electrolyte choice, the half-cell cycling capacity retention is mainly governed by the accessible capacity: as more capacity is delivered, an inevitable drop in capacity retention is observed. It appears that it may be impossible to reach the upper right corner of this graph (our target!) where one would have the highest capacity and also excellent capacity retention. So this represents a real challenge for Ni-rich positive electrode materials. To further assess the validity of the curve, we added two sets of half-cell cycling data adopted from the literature as data points 27 and 28.27

Samples 27 and 28 refer to NCM90

(LiNi0.901Co0.051Mn0.048O2) and NCM80 (LiNi0.796Co0.106Mn0..098O2) in reference 27, respectively. The half-cells in ref. 27 were cycled with a rate of C/2 instead of the C/5 rate used here and the reversible discharge capacity was measured with a rate of C/10 instead of C/20 used here. It is impressive that the data collected from a different research facility with multiple differences in experimental conditions (material synthesis, electrode fabrication, electrolyte formulation, etc.) also falls on the trend shown in Figure 4b.


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To examine the validity of the trend in Figure 4b, studies of the dQ/dV vs. V curves of cells shown in Figure 3 were made as examples. Figures 5a-5h plot dQ/dV vs. V of the 1st C/5 cycles of all the cells in black and the 50th cycles in red. Figures S3a-S3h show expanded views of dQ/dV vs. V of the 1st and 50th cycles. Comparing the dQ/dV vs. V curves before and after cycling shows that NA9505 and NCA900505 cells cycled with an UCV of 4.4 V developed large polarization during cycling, while NMg9505 shows more moderate polarization. With an UCV of 4.3 V, the polarization was smaller in all cases. In the case that the 4.2 V peak was excluded, little polarization was developed in NA9505 and NCA900505 cells. The analysis of the dQ/dV vs. V curves shows that the capacity loss observed in the half cells is mostly due to polarization. A question that needs to be answered is what caused the polarization growth? Polarization can be affected by electrolyte oxidation on existing surfaces,28 surface impurities,29 oxygen loss from the materials,20,30 as well as structural fractures which can expose new surfaces to electrolyte.7

Here, the comparison between NMg9505, NA9505 and

NCA900505 shows that with the same UCV, the polarization developed in NMg9505 cells is always the least, which suggests that electrolyte oxidation on existing surfaces may not be the dominant factor causing polarization growth. In terms of surface impurities, Figure S4a shows a blow-up view of the 1st cycle voltage profiles for NMg9505, NA9505, and NCA900505 cells. N.V. Faenza et al. showed that the open circuit voltage and the voltage overshoot in the beginning of the 1st cycle are good indicators of cell impedance caused by surface impurities.29 Figure S4a shows that NMg9505 and NA9505 cells have some small voltage overshoot in the beginning of the 1st cycle, while NCA900505 does not.. This observation suggests that there is no obvious correlation between the surface impurities and the cell cycling performance at least at the minor levels of the impurities in these samples. Figure 4a shows that symbols in the region 11 ACS Paragon Plus Environment


indicated by the red circle represent cells that have almost complete 4.2 V dQ/dV peaks, while cells in the region indicated by the blue circle only have partial or no 4.2 V dQ/dV peaks. Figures 3a-3c show that the biggest difference before and after the 4.2 V peak is the unit cell volume. This strongly suggests that the half-cell capacity fade and polarization growth is related to the unit cell volume change, which can cause particle cracking and reactions of electrolyte on fresh surfaces coupled with oxygen loss to the electrolyte.30. Figure 4b suggests that the capacity retention trend is a universal behavior, which is not greatly affected by cathode composition or electrolyte choice. Figures 6a-6c show the a-axis lattice, caxis lattice, and unit cell volume as a function of cell voltage, collected from in-situ XRD experiments in previous works.12,13,31 Cells with higher Ni content experience greater c-axis lattice collapse and unit cell volume change when cycled to the same UCV. However, Figure 6d shows that when the normalized unit cell volume is plotted as a function of delithiation degree (x in Li1-xMO2), regardless of the Ni content which ranges from 50% to 95% and different cation substituents, all the materials show an almost identical trend of unit cell volume change. The onset of sharp unit cell volume change is virtually the same as the onset of the H2-H3 phase transition in LNO, which suggests that the unit cell volume contraction is an intrinsic property of LNO derivatives.

Figure 6e shows each cell’s capacity retention vs. the corresponding

delithiation degree (x in Li1-xMO2) calculated based the 1st C/20 discharge capacity, and in the same graph, the normalized unit cell volume change profile of NCA900505 is plotted as a representative. Figure 6e shows that the profile of unit cell volume change follows the pattern of the capacity retention versus delivered capacity curve. It is tempting to speculate that the universal unit cell volume change is the mechanism behind the reduction of capacity retention with delivered capacity. When the cells are cycled with less 12 ACS Paragon Plus Environment

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capacity, the volume change of the unit cell falls into the blue region of Figure 4, and relatively good capacity retention can be achieved. Once the capacity delivered by the cathode is higher than a certain value, the red region in Figure 4 is entered where much greater anisotropic volume change is inevitably experienced by the electrode materials. Resulting from the more aggressive structural change during each cycle, cathode material degrades faster and a much steeper capacity retention drop is expected. Figure 6e shows an outlier labeled as No.26. Table I shows that cells No.25 and No.26 were made with polycrystalline (PC) NMC811 electrodes provided by Vendor 1, and it is known by the authors that the PC-NMC811 cathode materials were coated with materials not disclosed by Vendor 1. Besides, cells No.10 made with LiNi0.99Mg0.01O2 surprisingly show that the 1% Mg cation substitution can greatly improve the cycling stability. . The role of Mg substitution is still debatable and it has been reported that Mg atoms, at low Mg content tend to stay in the lithium layer and act as “pillars” to enhance structural stability.

32–34.

Figure 6e correlates cathode failure more to the inevitable unit cell volume contraction, but less to the choice of cathode composition, an important route for moving to the top right of Figure 6e should be enhancing the structural stability of the material. Single crystal cathode materials represent a good approach to enhance particle integrity during cycling,35 although the single crystal cathode materials in this study do not stand out. It is possible that an appropriate electrolyte system is required for the single crystal cathode to demonstrate its advantages.36 Gradient core-shell materials as promoted by the research group of Y.K. Sun et al.37,38 may also represent a path toward the right corner of Figure 6e. A surface with a high dopant content changing smoothly to a core of extremely Ni-rich material may be successful. As many efforts are being made to improve the performance of positive electrode materials, it is hoped that Figure 6e can help researchers make more sensible evaluations. 13 ACS Paragon Plus Environment



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Conclusions Based on half coin cell cycling of LiNiO2 with different UCVs and LCVs, the H2-H3 phase transition was found to be a major cause of poor capacity retention. With cation substitutions, the derivatives of LNO, including NMC and NCA materials, can avoid this phase transition. However, the single phase unit cell contraction in the same composition range was found to be a universal behavior dependent on delithiation degree. The half coin cell cycling data collected from 26 different sets of cells showed that capacity retention worsens as accessible capacity increases, leading to an inevitable challenge for Ni-rich positive electrode materials. Most of the cells follow the trend in Figure 6e regardless of differences in cathode chemical composition, electrolyte, etc. By combining the capacity retention versus delivered capacity curve with the universal unit cell volume change, it is proposed that the anisotropic unit cell volume change is a contributing failure mechanism of Ni-rich positive electrode materials. This work is believed of to be of value for researchers working to understand the intrinsic properties of Ni-rich cathode materials, and may help them make further effective improvements. Acknowledgements The authors thank NSERC and Tesla Canada for the funding of this work under the auspices of the Industrial Research Chairs program.

Hongyang Li thanks the Nova Scotia Graduate

Scholarship program for financial support. Ning Zhang thanks the China Scholarship Council for generous support. Aaron Liu thanks the Walter C. Sumner Foundation for financial support. Supporting Information: Further electrochemical and X-ray diffraction data for interested readers.



References (1)

Myung, S. T.; Maglia, F.; Park, K. J.; Yoon, C. S.; Lamp, P.; Kim, S. J.; Sun, Y. K. Nickel-Rich Layered Cathode Materials for Automotive Lithium-Ion Batteries: Achievements and Perspectives. ACS Energy Lett. 2017, 2 (1), 196–223. https://doi.org/10.1021/acsenergylett.6b00594.

(2)

Kobayashi, Y.; Tabuchi, M.; Miyashiro, H.; Kuriyama, N. A New Design of Highly Reversible LiNiO2: Defect Formation in Transition Metal Site. J. Power Sources 2017, 364, 156–162. https://doi.org/10.1016/j.jpowsour.2017.08.027.

(3)

Noh, H. J.; Youn, S.; Yoon, C. S.; Sun, Y. K. Comparison of the Structural and Electrochemical Properties of Layered Li[NixCoyMnz]O2 (X = 1/3, 0.5, 0.6, 0.7, 0.8 and 0.85) Cathode Material for Lithium-Ion Batteries. J. Power Sources 2013, 233, 121–130. https://doi.org/10.1016/j.jpowsour.2013.01.063.

(4)

Lee, K. K.; Yoon, W. S.; Kim, K. B.; Lee, K. Y.; Hong, S. T. Characterization of LiNi0.85Co0.10M0.05O2 (M = Al, Fe) as a Cathode Material for Lithium Secondary Batteries. J. Power Sources 2001, 97-98, 308–312. https://doi.org/10.1016/S03787753(01)00516-X.

(5)

Kim, J.-H.; Kim, S. J.; Yuk, T.; Kim, J.; Yoon, C. S.; Sun, Y.-K. Variation of Electronic Conductivity within Secondary Particles Revealing a Capacity-Fading Mechanism of Layered Ni-Rich Cathode. ACS Energy Lett. 2018, 3002–3007. https://doi.org/10.1021/acsenergylett.8b02043.

(6)

Yoon, C. S.; Ryu, H. H.; Park, G. T.; Kim, J. H.; Kim, K. H.; Sun, Y. K. Extracting 16 ACS Paragon Plus Environment

Page 16 of 32

Page 17 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Maximum Capacity from Ni-Rich Li[Ni0.95Co0.025Mn0.025]O2 Cathodes for HighEnergy-Density Lithium-Ion Batteries. J. Mater. Chem. A 2018, 6 (9), 4126–4132. https://doi.org/10.1039/c7ta11346c. (7)

Ryu, H.-H.; Park, K.-J.; Yoon, C. S.; Sun, Y.-K. Capacity Fading of Ni-Rich Li[NixCoyMn1-X-y]O2 (0.6 ≤ X ≤ 0.95) Cathodes for High Energy Density Lithium-Ion Batteries: Bulk or Surface Degradation? Chem. Mater. 2018, 2, acs.chemmater.7b05269. https://doi.org/10.1021/acs.chemmater.7b05269.

(8)

Muto, S.; Sasano, Y.; Tatsumi, K.; Sasaki, T.; Horibuchi, K.; Takeuchi, Y.; Ukyo, Y. Capacity-Fading Mechanisms of LiNiO[sub 2]-Based Lithium-Ion Batteries. J. Electrochem. Soc. 2009, 156 (5), A371. https://doi.org/10.1149/1.3076137.

(9)

Makimura, Y.; Zheng, S.; Ikuhara, Y.; Ukyo, Y. Microstructural Observation of LiNi0.8Co0.15Al0.05O2 after Charge and Discharge by Scanning Transmission Electron Microscopy. J. Electrochem. Soc. 2012, 159 (7), A1070–A1073. https://doi.org/10.1149/2.073207jes.

(10)

Zhang, H.; Omenya, F.; Yan, P.; Luo, L.; Whittingham, M. S.; Wang, C.; Zhou, G. RockSalt Growth-Induced (003) Cracking in a Layered Positive Electrode for Li-Ion Batteries. ACS Energy Lett. 2017, 2 (11), 2607–2615. https://doi.org/10.1021/acsenergylett.7b00907.

(11)

Liu, H.; Wolf, M.; Karki, K.; Yu, Y. S.; Stach, E. A.; Cabana, J.; Chapman, K. W.; Chupas, P. J. Intergranular Cracking as a Major Cause of Long-Term Capacity Fading of Layered Cathodes. Nano Lett. 2017, 17 (6), 3452–3457. https://doi.org/10.1021/acs.nanolett.7b00379.



(12)

Weber, R.; Fell, C. R.; Dahn, J. R.; Hy, S. Operando X-Ray Diffraction Study of Polycrystalline and Single-Crystal LixNi0.5Mn0.3Co0.2O2. J. Electrochem. Soc. 2017, 164 (13), A2992–A2999. https://doi.org/10.1149/2.0441713jes.

(13)

Li, H.; Cormier, M.; Zhang, N.; Inglis, J.; Li, J.; Dahn, J. R. Is Cobalt Needed in Ni-Rich Positive Electrode Materials for Lithium Ion Batteries? J. Electrochem. Soc. 2019, 166 (4), A429–A439. https://doi.org/10.1149/2.1381902jes.

(14)

Ryu, H. H.; Park, K. J.; Yoon, C. S.; Sun, Y. K. Capacity Fading of Ni-Rich li[NixCoyMn1-X-y]O2(0.6 ≤ X ≤ 0.95) Cathodes for High-Energy-Density Lithium-Ion Batteries: Bulk or Surface Degradation? Chem. Mater. 2018, 30 (3), 1155–1163. https://doi.org/10.1021/acs.chemmater.7b05269.

(15)

Li, W.; Yaghoobnejad Asl, H.; Xie, Q.; Manthiram, A. Collapse of LiNi1-X-yCoxMnyO2 Lattice at Deep Charge Irrespective of Nickel Content in Lithium-Ion Batteries. J. Am. Chem. Soc. 2019, jacs.8b13798. https://doi.org/10.1021/jacs.8b13798.

(16)

Kim, U. H.; Jun, D. W.; Park, K. J.; Zhang, Q.; Kaghazchi, P.; Aurbach, D.; Major, D. T.; Goobes, G.; Dixit, M.; Leifer, N.; et al. Pushing the Limit of Layered Transition Metal Oxide Cathodes for High-Energy Density Rechargeable Li Ion Batteries. Energy Environ. Sci. 2018, 11 (5), 1271–1279. https://doi.org/10.1039/c8ee00227d.

(17)

Kim, H.; Kim, M. G.; Jeong, H. Y.; Nam, H.; Cho, J. A New Coating Method for Alleviating Surface Degradation of LiNi0.6Co0.2Mn0.2O2 Cathode Material: Nanoscale Surface Treatment of Primary Particles. Nano Lett. 2015, 15 (3), 2111–2119. https://doi.org/10.1021/acs.nanolett.5b00045.


Page 18 of 32

Page 19 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(18)

Park, K. J.; Jung, H. G.; Kuo, L. Y.; Kaghazchi, P.; Yoon, C. S.; Sun, Y. K. Improved Cycling Stability of Li[Ni0.90Co0.05Mn0.05]O2Through Microstructure Modification by Boron Doping for Li-Ion Batteries. Adv. Energy Mater. 2018, 8 (25). https://doi.org/10.1002/aenm.201801202.

(19)

Yoon, C. S.; Park, K. J.; Kim, U. H.; Kang, K. H.; Ryu, H. H.; Sun, Y. K. High-Energy Ni-Rich Li[NixCoyMn1-X-y]O2Cathodes via Compositional Partitioning for NextGeneration Electric Vehicles. Chem. Mater. 2017, 29 (24), 10436–10445. https://doi.org/10.1021/acs.chemmater.7b04047.

(20)

Jung, R.; Strobl, P.; Maglia, F.; Stinner, C.; Gasteiger, H. A. Temperature Dependence of Oxygen Release from LiNi 0.6 Mn 0.2 Co 0.2 O 2 (NMC622) Cathode Materials for LiIon Batteries . J. Electrochem. Soc. 2018, 165 (11), A2869–A2879. https://doi.org/10.1149/2.1261811jes.

(21)

Li, J.; Camardese, J.; Glazier, S.; Dahn, J. R. Structural and Electrochemical Study of the Li − Mn − Ni Oxide System within the Layered Single Phase Region. Chem. Mater. 2014, 26 (24), 7059–7066. https://doi.org/10.1021/cm503505b.

(22)

Li, J.; Zhang, N.; Li, H.; Liu, A.; Wang, Y.; Yin, S.; Wu, H.; Dahn, J. R.. Impact of the Synthesis Conditions on the Performance of LiNiXCoYA ZO 2 with High Ni and Low Co Content . J. Electrochem. Soc. 2018, 165 (14), A3544–A3557. https://doi.org/10.1149/2.0931814jes.

(23)

B.A. Hunter, C. J. H. Australian Nuclear Science and Technology of X-Ray and Neutron; 1998.



(24)

Marks, T.; Trussler, S.; Smith, A. J.; Xiong, D.; Dahn, J. R. A Guide to Li-Ion Coin-Cell Electrode Making for Academic Researchers. J. Electrochem. Soc. 2011, 158 (1), A51. https://doi.org/10.1149/1.3515072.

(25)

Li, H.; Zhang, N.; Li, J.; Dahn, J. R. Updating the Structure and Electrochemistry of Li x NiO 2 for 0 ≤ X ≤ 1. J. Electrochem. Soc. 2018, 165 (13), A2985–A2993. https://doi.org/10.1149/2.0381813jes.

(26)

Yoon, C. S.; Jun, D.-W.; Myung, S.-T.; Sun, Y.-K. Structural Stability of LiNiO 2 Cycled above 4.2 V. ACS Energy Lett. 2017, 2 (5), 1150–1155. https://doi.org/10.1021/acsenergylett.7b00304.

(27)

Ryu, H. H.; Park, G. T.; Yoon, C. S.; Sun, Y. K. Microstructural Degradation of Ni-Rich Li[NixCoyMn1−x−y]O2 Cathodes During Accelerated Calendar Aging. Small 2018, 14 (45). https://doi.org/10.1002/smll.201803179.

(28)

Xu, K. Electrolytes and Interphases in Li-Ion Batteries and Beyond. Chem. Rev. 2014, 114 (23), 11503–11618. https://doi.org/10.1021/cr500003w.

(29)

Faenza, N. V.; Bruce, L.; Lebens-Higgins, Z. W.; Plitz, I.; Pereira, N.; Piper, L. F. J.; Amatucci, G. G. Editors’ Choice—Growth of Ambient Induced Surface Impurity Species on Layered Positive Electrode Materials and Impact on Electrochemical Performance. J. Electrochem. Soc. 2017, 164 (14), A3727–A3741. https://doi.org/10.1149/2.0921714jes.

(30)

Jung, R.; Metzger, M.; Maglia, F.; Stinner, C.; Gasteiger, H. A. Oxygen Release and Its Effect on the Cycling Stability of LiNi X Mn Y Co Z O 2 (NMC) Cathode Materials for Li-Ion Batteries . J. Electrochem. Soc. 2017, 164 (7), A1361–A1377. 20 ACS Paragon Plus Environment

Page 20 of 32

Page 21 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


https://doi.org/10.1149/2.0021707jes. (31)

Li, J.; Downie, L. E.; Ma, L.; Qiu, W.; Dahn, J. R. Study of the Failure Mechanisms of LiNi0.8Mn0.1Co0.1O2 Cathode Material for Lithium Ion Batteries. J. Electrochem. Soc. 2015, 162 (7), A1401–A1408. https://doi.org/10.1149/2.1011507jes.

(32)

Xie, Q.; Li, W.; Manthiram, A. A Mg-Doped High-Nickel Layered Oxide Cathode Enabling Safer, High-Energy-Density Li-Ion Batteries. Chem. Mater. 2019, 31 (3), 938– 946. https://doi.org/10.1021/acs.chemmater.8b03900.

(33)

Chang, C. C.; Kim, J. Y.; Kumta, P. N. Synthesis and Electrochemical Characterization of Divalent Cation-Incorporated Lithium Nickel Oxide. J. Electrochem. Soc. 2000, 147 (5), 1722–1729. https://doi.org/Doi 10.1149/1.1393424.

(34)

Pouillerie, C.; Croguennec, L.; Biensan, P.; Willmann, P.; Delmas, C. Synthesis and Characterization of New LiNi1−yMgyO2 Positive Electrode Materials for Lithium‐Ion Batteries. J. Electrochem. Soc. 2000, 147, 2061–2069. https://doi.org/10.1149/1.1393486.

(35)

Kim, J.; Lee, H.; Cha, H.; Yoon, M.; Park, M.; Cho, J. Prospect and Reality of Ni-Rich Cathode for Commercialization. Adv. Energy Mater. 2017, 1702028, 1–25. https://doi.org/10.1002/aenm.201702028.

(36)

Li, J.; Cameron, A. R.; Li, H.; Glazier, S.; Xiong, D.; Chatzidakis, M.; Allen, J.; Botton, G. A.; Dahn, J. R. Comparison of Single Crystal and Polycrystalline for High Voltage Li-Ion Cells. J. Electrochem. Soc. 2017, 164 (7), 1534–1544. https://doi.org/10.1149/2.0991707jes.



(37)

Sun, Y. K.; Myung, S. T.; Park, B. C.; Prakash, J.; Belharouak, I.; Amine, K. High-Energy Cathode Material for Long-Life and Safe Lithium Batteries. Nat. Mater. 2009, 8 (4), 320– 324. https://doi.org/10.1038/nmat2418.

(38)

Myung, S. T.; Noh, H. J.; Yoon, S. J.; Lee, E. J.; Sun, Y. K. Progress in High-Capacity Core-Shell Cathode Materials for Rechargeable Lithium Batteries. J. Phys. Chem. Lett. 2014, 5 (4), 671–679. https://doi.org/10.1021/jz402691z.


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Figure 1.

Figure 1. SEM images of (a) LiNiO2 made using a Ni(OH)2 precursor from Zoomwe, (b) LiNi0.95Mg0.05O2 made using a Ni0.95Mg0.05(OH)2 precursor made at Dalhousie University, and (c) LiNi0.80Co0.15Al0.05O2 from Ecopro (c).



Figure 2.

Figure 2. (a) Specific capacity and (b) normalized capacity of LiNiO2 with different UCVs and LCVs as a function of cycle number. (c1 – c5) Differential capacity vs. voltage (dQ/dV vs. V) of first (black) and last (red) C/5 cycles for LiNiO2 half coin cells. (d1 – d5) dQ/dV vs. V of first (black) and last (red) C/20 cycles for LiNiO2 half coin cells. Testing was made at 30oC.


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Figure 3

Figure 3. dQ/dV vs. V and unit cell volume as a function of cell voltage for (a) NMg9505, (b) NA9505, and (c) NCA900505 cells. Specific capacity (d – f) and normalized capacity (g – i) as a function of cycle number for NMg9505 cells, NA9505 cells, and NCA900505 cells, respectively. All cell testing was made at 30oC.



Figure 4

Figure 4. Capacity retention after 50 cycles (C/5) as a function of 1st C/20 discharge capacity for (a) NMg9505, NA9505 and NCA900505 cells, and (b) cells listed in Table I.


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Figure 5.

Figure 5. (a – h) dQ/dV vs. V of the 1st C/5 cycles for all the cells in black and the 50th cycles in red. Solid and dashed lines in the same color represent pair cells. The testing was all made at 30oC. In panels a, b and c, the upper cutoff potential as 4.3 V, on panels d, e and f it was 4.2 V while in panels g and h, it was 4.1 V

An expanded view of the 4.2 V peak region is given in

Figure S3.



Page 28 of 32

Figure 6

Figure 6. (a) a-axis lattice parameter, (b) c-axis lattice parameter, and (c) unit cell volume as a function

of

cell

voltage

for

LiNi0.95Al0.05O2,

LiNi0.95Mn0.05O2,

LiNi0.95Mg0.05O2,

LiNi0.9Co0.05Al0.05O2, LiNi0.8Mn0.1Co0.1O2, and LiNi0.5Mn0.3Co0.2O2. (d) Normalized unit cell 28 ACS Paragon Plus Environment

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volume as a function of x in LixMO2. (e) Capacity retention after 50 cycles (C/5) as a function of x in Li1-xMO2 for cells listed in Table I.


Chemistry of Materials 1 2 3 4 5 6Sample 7Number 8 1 9 2 10 11 3 12 4 13 14 5 15 6 16 7 17 18 8 19 9 20 21 10 22 11 23 12 24 25 13 26 14 27 28 15 29 16 30 17 31 32 18 33 19 34 35 20 36 37 21 38 22 39 23 40 41 24 42 25 43 44 26 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 30 of 32

Table I: Summary of half coin cell cycling tests. Compositions (%)

0.973

267.7

243.0

x in Li1xNiO2* **

Material Maker ****

UCV (V)

220.9

173.3

0.7844

0.8869

Zoomwe-Dal

4.3

LNO

FEC DMC

100

NCAMg89/6/3/2

FEC DMC

89

6

3

2

0.714

226.5

195.1

184.5

179.9

0.9751

0.7119

Zoomwe-Dal

4.3

NCAMg92/3/3/2

FEC DMC

92

3

3

2

0.9

234.1

201.3

191.7

186.6

0.9735

0.7347

Zoomwe-Dal

4.3

NCAMg95/0/3/2

FEC DMC

95

3

2

1.485

236.8

204.1

192.6

178.0

0.9239

0.7450

Zoomwe-Dal

4.3

NCAMg89/6/3/2

FEC DMC

89

6

3

2

0.714

233.9

204.3

196.3

191.0

0.9732

0.7454

Zoomwe-Dal

4.4

NCAMg92/3/3/2

FEC DMC

92

3

3

2

0.9

226.7

193.3

183.4

175.7

0.9579

0.7055

Zoomwe-Dal

4.25

NCAMg95/0/3/2

FEC DMC

95

3

2

1.485

226.5

190.1

178.2

164.6

0.9235

0.6938

Zoomwe-Dal

4.25

NMg95/5

EC DEC

95

5

1.36

234.1

196.4

183.5

178.5

0.9733

0.7168

Zoomwe-Dal

4.3

NMg97.5/2.5

FEC DMC

97.5

2.5

0.987

242.2

226.3

211.0

192.5

0.9126

0.8258

Zoomwe-Dal

4.3

NMg99/1

FEC DMC

99

1

0.788

250.1

235.2

216.7

189.8

0.876

0.858

Zoomwe-Dal

4.3

NCA90/5/5

FEC DMC

90

5

5

1.095

185.4

165.6

155.2

145.1

0.9349

0.6044

Zoomwe-Dal

4.1

NCA90/5/5

FEC DMC

90

5

5

1.095

235.2

215.9

204.3

179.0

0.8763

0.7878

Zoomwe-Dal

4.3

NCA90/5/5

FEC DMC

90

5

5

1.095

241.4

222.0

208.6

181.4

0.8697

0.8103

Zoomwe-Dal

4.4

NA95/5

FEC DMC

95

5

1.219

181.7

159.1

149.5

146.1

0.9771

0.5806

Zoomwe-Dal

4.1

NA95/5

FEC DMC

95

5

1.219

244.6

218.7

206.6

193.4

0.9362

0.7982

Zoomwe-Dal

4.3

NA95/5

FEC DMC

95

5

1.219

244.7

226.3

214.9

192.9

0.8977

0.8260

Zoomwe-Dal

4.4

NMg95/05

FEC DMC

95

5

1.658

229.3

203.0

183.4

178.9

0.9759

0.7409

Dal

4.3

NMg95/05

FEC DMC

95

5

1.658

238.3

214.5

202.6

189.7

0.9359

0.7828

Dal

4.4

NA98.75/1.25

FEC DMC

98.75

1.5

259.4

240.9

223.2

187.6

0.8407

0.8792

Dal

4.3

NMC622_PC

2FEC DTDECDEC

60

20

20

2.5

209.5

197.0

187.4

173.0

0.9231

0.7191

Dal

4.4

NMC622_SC*

2FEC DTDECDEC

60

20

20

2.1

211.2

192.9

183.1

172.8

0.9438

0.7041

Dal

4.4

NCA80/15/05

FEC DMC

80

15

1.3

224.1

217.5

208.9

196.5

0.9407

0.7937

Ecopro

4.4

NMC811_SC*

FEC DMC

80

10

10

N/A

230.0

208.6

199.2

189.4

0.9504

0.7613

Vendor 1

4.3

NMC811_SC*

FEC DMC

80

10

10

N/A

238.7

217.3

208.9

195.0

0.9333

0.7929

Vendor 1

4.4

NMC811_PC

FEC DMC

80

10

10

N/A

228.0

216.0

205.6

196.0

0.9531

0.7884

Vendor 1

4.3

NMC811_PC

FEC DMC

80

10

10

N/A

235.8

226.6

220.5

209.8

0.9515

0.8269

Vendor 1

4.4

SC: single crystal C

Mn

Retenti on after 50 cycles

Ni

**

Mg

50th C/5 Dicshar ge

Electrolyte

PC: polycrystalline

Al

C/20 Charg e

Sample Name

*

Co

Capacity (mAh g-1) C/20 1st C/5 Dischar Dischar ge ge

Ni in Li Layer (%)

1.3

5

20 Discharge capacity

***

x=

****

Zoomwe-Dal: Materials made at Dal with Zoomwe precursors;

274 mAh g ―1


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Dal: Materials made at Dal with Dal made precursors; Ecopro: Materials provided by Ecopro; Vendor 1: Dry pouch cell provided by Vendor 1



Table of Contents Graphic.


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An Unavoidable Challenge for Ni-rich Positive Electrode Materials for

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