Intergranular Cracking as a Major Cause of Long-Term Capacity

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Inter-granular cracking as a major cause of long-term capacity fading of layered cathodes Hao Liu, Mark Wolf, Khim Karki, Young-Sang Yu, Eric A. Stach, Jordi Cabana, Karena W Chapman, and Peter J. Chupas Nano Lett., Just Accepted Manuscript • Publication Date (Web): 26 May 2017 Downloaded from http://pubs.acs.org on May 27, 2017

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Inter-granular cracking as a major cause of longterm capacity fading of layered cathodes Hao Liu†, Mark Wolf§, Khim Karki‡, Young-Sang Yu§,#, Eric A. Stach‡, Jordi Cabana§, Karena W. Chapman†*, Peter J. Chupas¶* †

X-ray Science Division, Advanced Photon Source, Argonne National Laboratory, 9700 South Cass Ave., Argonne, IL, 60439, United States. §

Department of Chemistry, University of Illinois at Chicago, Chicago, Illinois 60607, United States.

‡

Center for Function Nanomaterials, Brookhaven National Laboratory, Upton, NY 11973-5000, United States.

#

Advanced Light Source, Lawrence Berkeley National Laboratory, Berkeley, California 94720,

USA. ¶

Science Directorate, Advanced Photon Source, Argonne National Laboratory, 9700 South Cass Ave., Argonne, IL, 60439, United States.

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TABLE OF CONTENTS GRAPHIC

ABSTRACT

Capacity fading has limited commercial layered Li-ion battery electrodes to 90 cycles is the development of the “sluggish” NCA population, which showed reduced variation in the Li composition compared to the active group during charge and discharge.

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Absence of such heterogeneous Li composition in the initial cycles (first two cycles shown in Figure S4) indicates that this heterogeneity was developed only after extended cycling and not due to pre-existing regions with poorer electronic and ionic conduction than the rest of the electrode. Specifically, the “sluggish” population constitutes nearly 40% of the entire NCA population

(Figure S11b), which is not commensurate with the non-existent, pre-existing

inhomogeneity of the electrode. Although Li-compositional heterogeneity has been observed in a partially-charged LiNi1/3Mn1/3Co1/3O2 secondary particle even after long time of relaxation following low-rate charge19, this heterogeneity spans over a much smaller compositional range (0.1 Li per f.u.) than what was observed in the present study (~0.4 Li per f.u.) and only leads to symmetrical broadening of the Bragg reflections. Therefore, the heterogeneous Li composition and reaction kinetics observed at cycles 92 and 93 must arise from extended cycling. While heterogeneous states can develop between secondary particles in different parts of the electrode, for example, as a function of depth within the cell,20 this manifests as a uniform rather than bimodal Li composition distribution. Here, the distinct duality of reaction rates observed can only be attributed to heterogeneity within secondary particles as a result of heterogeneous electronic and ionic conductivity. Although degradation can be associated with a multitude of ageing mechanisms21, only the mechanisms that can modify conductivities within secondary particles are most relevant to the fade in capacity in the present work. In fact, changes in conductivity can be achieved only by changing the grain boundaries, through which electrons and ions are conducted between primary particles. Specifically, the primary particles at the surface of the secondary aggregates are well connected both electronically and ionically, while electronic and ionic transport to sub-surface primary particles is indirect, via grain boundaries, and can be compromised. Extensive internal cracking in cycled secondary particles (charged to 7


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4.5V after 93 cycles) was revealed in virtual slices of tomograms from transmission X-ray microscopy (Figure 4b) (at ~60 nm resolution). These cross-sectional images reflect the internal structure of the secondary electrode particles. Cracks are evident as areas of reduced density, that is, as darker lines that are uniformly distributed within each secondary particle. Given that the primary particle size did not change after 93 cycles (from XRD peak width analysis), the only consistent interpretation of the crack patterns is that the cracks are formed along grain (primary particle) boundaries, i.e. inter-granular cracking. In contrast, such cracking was clearly absent in pristine particles, which were found to be dense and free of voids (Figure 4a). Inter-granular cracking, at the boundaries between primary particles, is a well-recognized consequence of cyclic anisotropic changes in lattice dimension and is a common cause of thermal fatigue.22, 23 Here, the anisotropic changes in lattice dimension accompany Li extraction and insertion. Cracking at the grain boundaries compromise the connectivity and the reaction kinetics of buried particles. Inter-granular crack formation reduces the contact surface area between neighboring primary particles and increases the impedance for electronic and ionic transport. Formation of a NiO-like, rock-salt type structure at the particle surface has been reported for Ni-rich compounds cycled to high voltages3, 6, 17, 18, 24; the NiO-like structure has poor ionic and electronic conductivity and could lead to capacity fading. Here, the surface structural transformation is mainly associated with the first cycle, which is indicated by the 5% loss in the NCA phase fraction after the first cycle and is consistent with a previous transmission electron microscopy observation6. While the growth of the rock-salt type layer with progressive cycling has been proposed to account for the capacity loss24, we observed no effective growth of the surface layer after the first cycle (Figure S2), indicating that the capacity loss may not 8


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necessarily arise from the growth of the rock-salt type surface layer. In fact, only a moderate increase in the charge transfer resistance by a factor of 2 has been reported for an NCA electrode held at 4.5V for 5 weeks25; in comparison, soaking LiNi0.4Mn0.4Co0.18Ti0.02O2 in the electrolyte for a week increases the charge transfer resistance also by a factor of 224. The very little capacity loss of the active NCA particles, where chemical degradation also occurs, suggests that chemical degradation only makes a minor contribution to the capacity loss. Moreover, the homogeneous reaction during the second cycle (Figure S4) suggests that the formation of the rock-salt type surface layer alone does not lead to heterogeneous reaction. In contrast, inter-granular cracking allows the electrolyte to seep into the secondary particles through the cracks, exposing buried surfaces of the primary particles to the electrolyte. This can induce further structural and compositional changes at the rock-salt type surface layer via electrolyte decomposition3, oxygen release26 and metal dissolution27, which modify the conductivity of the surface layer. As new surfaces are progressively exposed to the electrolyte, heterogeneous conductivity will develop between primary particles and lead to heterogeneous reaction. Therefore, the heterogeneous Licomposition and reaction kinetics developed between primary particles at extended cycles can ultimately be attributed to inter-granular cracking within the secondary particle aggregates. Here, the use of slow, rather than accelerated cycling rates, avoids compositional inhomogeneity due to rate-induced compositional gradient across the electrode and the primary and secondary particles. Under high rate conditions, gradients within the secondary and primary particles can further contribute to inter-granular cracking and, even, trans-granular cracking (fracture of primary particles).28 While most operando12, 13 and post-mortem3, 6 studies occur in this high-rate regime, where there are multiple driving forces for cracking, the slow-rate regime studied here demonstrates that this phenomenon is a fundamental and unavoidable contribution to capacity 9


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fade in commercial systems. In situ acoustic emission29 and electron microscopy12 studies indicate that a substantial number of inter-granular cracks are initiated during the first cycle. However, the heterogeneous reaction kinetics do not develop until inter-granular cracks propagate under continued cycling, leading to increased separation and possible isolation of primary particles from the secondary particles, as has been observed after ten to a few hundred cycles3, 6, 30. To decouple the mechanical from the chemical contribution to the capacity loss, we cycled NCA electrodes to the same upper cut-off voltage at 4.5 V, which induces the same level of chemical change, but to different lower cut-off voltages (4.25 V, 4.05 V and 2.5 V), which induces different levels of lattice volume change (Table 1 and Figure S12). The particular values of different lower cut-off voltages were selected such that cycling within 2.5 – 4.05 V, 4.05 – 4.25 V and 4.25 – 4.5 V induces approximately the same amount of volume change (~2%). The current rate is dependent on the voltage range: C/10 (28 mA/g) for 2.5 – 4.05 V, C/50 (5.6 mA/g) for 4.05 – 4.25 V, and C/100 (2.8 mA/g) for 4.25 – 4.5 V. This ensures a uniform rate of volume change across the full voltage window (2.5 – 4.5 V). For electrodes cycled within limited voltage windows, full discharge to 2.5V was performed for select cycles (the 19th, 50th and 98th cycles for electrode A and the 20th and 48th cycles for electrode B in Figure 5a) to compare the capacity loss with the electrode constantly cycled in the full voltage range (2.5 – 4.5V, electrode C in Figure 5a). The electrode cycled within a larger voltage window shows a much more rapid fade in capacity with increasing number of cycles (Figure 5a). It is important to note that, even for the same amount of cycling time, the electrode cycled between 2.5 – 4.5V showed consistently lower discharge capacity than those cycled only in the high voltage regime (Figure 5b and Table

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1), further demonstrating the principal contribution of mechanical processes to the capacity loss (for at least 4.5V). The capacity of NCA fades by secondary particle fatigue, as summarized in Figure 4c, where the cyclic stress induced by Li extraction and insertion promotes the growth of inter-granular cracks. Inter-granular cracks form and propagate progressively within the secondary particles as they are cycled, leading to reduced electronic and ionic transport for buried primary particles, the development of reaction heterogeneity and capacity fade. Widening of such inter-granular cracks could ultimately expose the surface embedded in the grain boundary to the electrolyte, forming a passivation layer3,

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and further increasing the electrical and ionic resistance at the grain

boundary. Ultimately this could lead to complete isolation of some primary particles with no electrochemical activity; at high rates, where there are additional contributions to cracking, an inactive phase has been observed in an extensively-cycled electrode.13 The capacity fade mechanism observed here is fundamental to non-cubic electrode compounds that exist as polycrystalline aggregates. The aggregate morphology of NCA is dictated by the coprecipitation synthesis method,31 which is widely used to synthesize layered electrodes with mixed transition metal stoichiometry with the general chemical formula of LiNixMnyCo1-x-yO2. For cubic-phase materials, where the lattice changes during cycling are isotropic, inter-granular crack formation can be mitigated. Indeed, the cubic spinel LiNi0.5Mn1.5O4 exhibits no appreciable fade in capacity after 100 cycles32 despite a high cut-off voltage on charge (4.95 V) and a similar volume change to NCA during cycling. However, gradients induced by high rate of cycling will induce lattice mismatches between primary particles and give rise to inter-granular cracking. Electrodes with the best lifetimes have less cracking17,

30, 33

: Strategies to deliver the best

lifetimes, should minimize strain at grain boundaries due to anisotropic changes in lattice 11


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dimension. This can be achieved in multiple ways: (i) by eliminating secondary particles aggregate structure or manipulating the aggregate structure using smaller primary particles in the same secondary particles to reduce the strain developed at each grain boundary, (ii) by selecting voltage ranges where anisotropic changes in lattice dimension are minimized, and (iii) by mechanically mitigating lattice changes by coating secondary particles. This underlies the known trend that while a wide voltage yields a higher capacity initially, this gain is eroded by a faster fade in capacity;1, 2 the wide voltage range that yields a higher initial Li utilization and capacity also induce larger lattice changes that accelerates inter-granular cracking and, ultimately, capacity fade. It is important to note that chemical degradation, including surface reconstruction, is highly dependent on the upper cut-off voltage. Charging to higher voltages (>4.7V)24, 25 will induce a much greater increase in the charge transfer resistance than to 4.5V, so chemical degradation will probably become the more dominant contributor to capacity fading at very high voltages. By following the electrode’s reaction process during extended cycles via quantitative operando measurements, we have successfully identified the dominant capacity fading mechanism of a commercial electrode material as the formation and growth of inter-granular cracks in secondary electrode particles. The mechanical origin of the degradation process linked to the microstructure of the electrode particles, suggests that the capacity fade can be minimized by controlling the morphology of the primary and secondary electrode particles. We believe that long-term quantitative operando studies, such as present here, hold the key to deciphering failure and degradation in battery electrode materials.

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ASSOCIATED CONTENT Supporting Information Available: Experimental methods, the nature of peak splitting, the length scale of phase segregation, Figure S1-S12, Table S1. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] (K.W.C.) *E-mail: [email protected] (P.J.C.)

Conflict of Interest Disclosure The authors declare no competing financial interests.

ACKNOWLEDGEMENT This research is supported as part of the NorthEast Center for Chemical Energy Storage (NECCES), an Energy Frontier Research Center funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Award Number DE-SC0012583. YSY also acknowledges support from an Advanced Light Source Collaborative Postdoctoral Fellowship. This research used resources of the Advanced Photon Source, a U.S. Department of Energy (DOE) Office of Science User Facility operated for the DOE Office of Science by Argonne National Laboratory under Contract No. DE-AC02-06CH11357. Use of the Stanford Synchrotron Radiation Lightsource, SLAC National Accelerator Laboratory, is supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Contract 13


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No. DE-AC02-76SF00515. SEM was performed at Center for Functional Nanomaterials, which is a U.S. DOE Office of Science Facility, at Brookhaven National Laboratory under Contract No. DE-SC0012704. REFERENCES 1. Yin, S. C.; Rho, Y. H.; Swainson, I.; Nazar, L. F. Chem. Mater. 2006, 18, 1901-1910. 2. Robert, R.; Novak, P. J. Electrochem. Soc. 2015, 162, A1823-A1828. 3. Watanabe, S.; Kinoshita, M.; Hosokawa, T.; Morigaki, K.; Nakura, K. J. Power Sources 2014, 258, 210-217. 4. Abraham, D. P.; Twesten, R. D.; Balasubramanian, M.; Petrov, I.; McBreen, J.; Amine, K. Electrochem. Commun. 2002, 4, 620-625. 5. Itou, Y.; Ukyo, Y. J. Power Sources 2005, 146, 39-44. 6. Makimura, Y.; Zheng, S.; Ikuhara, Y.; Ukyo, Y. J. Electrochem. Soc. 2012, 159, A1070A1073. 7. Ogata, K.; Salager, E.; Kerr, C. J.; Fraser, A. E.; Ducati, C.; Morris, A. J.; Hofmann, S.; Grey, C. P. Nat. Commun. 2014, 5. 8. Ebner, M.; Marone, F.; Stampanoni, M.; Wood, V. Science 2013, 342, 716-720. 9. Pietsch, P.; Westhoff, D.; Feinauer, J.; Eller, J.; Marone, F.; Stampanoni, M.; Schmidt, V.; Wood, V. Nat. Commun. 2016, 7, 12909. 10. Huang, J. Y.; Zhong, L.; Wang, C. M.; Sullivan, J. P.; Xu, W.; Zhang, L. Q.; Mao, S. X.; Hudak, N. S.; Liu, X. H.; Subramanian, A.; Fan, H.; Qi, L.; Kushima, A.; Li, J. Science 2010, 330, 1515-1520. 11. Lei, J.; McLarnon, F.; Kostecki, R. J. Phys. Chem. B 2005, 109, 952-957. 12. Miller, D. J.; Proff, C.; Wen, J. G.; Abraham, D. P.; Bareño, J. Adv. Energy Mater. 2013, 3, 1098-1103. 13. Kleiner, K.; Dixon, D.; Jakes, P.; Melke, J.; Yavuz, M.; Roth, C.; Nikolowski, K.; Liebau, V.; Ehrenberg, H. J. Power Sources 2015, 273, 70-82. 14. C/n corresponds to the current required to charge or discharge the battery to its theoretical cpacity in n hours 15. Borkiewicz, O. J.; Shyam, B.; Wiaderek, K. M.; Kurtz, C.; Chupas, P. J.; Chapman, K. W. J. Appl. Crystallogr. 2012, 45, 1261-1269. 16. Abraham, D. P.; Twesten, R. D.; Balasubramanian, M.; Kropf, J.; Fischer, D.; McBreen, J.; Petrov, I.; Amine, K. J. Electrochem. Soc. 2003, 150, A1450. 17. Watanabe, S.; Hosokawa, T.; Morigaki, K. i.; Kinoshita, M.; Nakura, K. ECS Transactions 2012, 41, 65-74. 18. Zheng, S.; Huang, R.; Makimura, Y.; Ukyo, Y.; Fisher, C. A. J.; Hirayama, T.; Ikuhara, Y. J. Electrochem. Soc. 2011, 158, A357. 19. Gent, W. E.; Li, Y.; Ahn, S.; Lim, J.; Liu, Y.; Wise, A. M.; Gopal, C. B.; Mueller, D. N.; Davis, R.; Weker, J. N.; Park, J.-H.; Doo, S.-K.; Chueh, W. C. Adv. Mater. 2016, 28, 6631-6638. 20. Strobridge, F. C.; Orvananos, B.; Croft, M.; Yu, H. C.; Robert, R.; Liu, H.; Zhong, Z.; Connolley, T.; Drakopoulos, M.; Thornton, K.; Grey, C. P. Chem. Mater. 2015, 27, 2374-2386. 14


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21. Vetter, J.; Novák, P.; Wagner, M. R.; Veit, C.; Möller, K. C.; Besenhard, J. O.; Winter, M.; Wohlfahrt-Mehrens, M.; Vogler, C.; Hammouche, A. J. Power Sources 2005, 147, 269-281. 22. Tvergaard, V.; Hutchinson, J. W. J. Am. Ceram. Soc. 1988, 71, 157-166. 23. Kitaoka, S.; Matsushima, Y.; Chen, C.; Awaji, H. J. Am. Ceram. Soc. 2004, 87, 906-913. 24. Lin, F.; Markus, I. M.; Nordlund, D.; Weng, T.-C.; Asta, M. D.; Xin, H. L.; Doeff, M. M. Nat. Commun. 2014, 5. 25. Sallis, S.; Pereira, N.; Mukherjee, P.; Quackenbush, N. F.; Faenza, N.; Schlueter, C.; Lee, T. L.; Yang, W. L.; Cosandey, F.; Amatucci, G. G.; Piper, L. F. J. Appl. Phys. Lett. 2016, 108, 263902. 26. Karki, K.; Huang, Y.; Hwang, S.; Gamalski, A. D.; Whittingham, M. S.; Zhou, G.; Stach, E. A. ACS Applied Materials & Interfaces 2016, 8, 27762-27771. 27. Gilbert, J. A.; Bareño, J.; Spila, T.; Trask, S. E.; Miller, D. J.; Polzin, B. J.; Jansen, A. N.; Abraham, D. P. J. Electrochem. Soc. 2016, 164, A6054-A6065. 28. Zhao, K.; Pharr, M.; Vlassak, J. J.; Suo, Z. J. Appl. Phys. 2010, 108, 073517. 29. Woodford, W. H.; Carter, W. C.; Chiang, Y.-M. Energ. Environ. Sci. 2012, 5, 8014. 30. Lee, S.-H.; Yoon, C. S.; Amine, K.; Sun, Y.-K. J. Power Sources 2013, 234, 201-207. 31. Lee, M. H.; Kang, Y. J.; Myung, S. T.; Sun, Y. K. Electrochim. Acta 2004, 50, 939-948. 32. Wu, H. M.; Belharouak, I.; Deng, H.; Abouimrane, A.; Sun, Y. K.; Amine, K. J. Electrochem. Soc. 2009, 156, A1047. 33. Kim, H.; Kim, M. G.; Jeong, H. Y.; Nam, H.; Cho, J. Nano Lett. 2015, 15, 2111-2119.

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FIGURES

Figure 1. Scanning electron micrograph of uncycled NCA primary and secondary particles. The white dashed line delineates the boundary of a primary particle.

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Figure 2. The discharge capacity of NCA and the X-ray diffraction (XRD) patterns at different states of charge. (a) The specific discharge capacity of the NCA electrode cycled in the voltage window of 2.7-4.5 V at a current rate of 14 mA/g (black dots). The red dots represent the specific discharge capacity, quantified from operando XRD patterns. (b) The (113) reflections at different states of charge obtained from operando XRD. At the end of the 92nd charge, the reflection has a long tail towards lower 2θ angle, which is not evident at the end of the 2nd charge.

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Figure 3. The NCA electrode reaction captured by operando XRD during the 92nd and 93rd cycles reveals the presence of two NCA populations within electrode that react at different rates: an active population that retains the behavior of in the original electrode and a sluggish population in which the reaction is retarded. (a) Plots of the (003) and (113) reflections from operando XRD with contour levels of constant intensity (black). (b) Distribution of the Li composition of the NCA electrode based on analysis of the diffraction data. The contour level corresponds to a population density at 0.05. The centroids of the active (dashed green line) and sluggish (dashed purple line) populations are shown. (c) The corresponding voltage profile (red) and the average Li composition (blue) of the electrode quantified from operando XRD. The blank regions corresponds to a brief loss of the synchrotron X-ray beam.

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Figure 4. Virtual slices of transmission X-ray tomograms collected for NCA (a) in its pristine state and (b) charged to 4.5 V following 93 cycles (2.7 – 4.5 V, C/20) showing substantial fracturing after extended cycling. (c) A schematic illustrating the process leading to degradation in a secondary NCA particle. The arrow in each primary particle represents the different crystal orientations within the aggregate. The first cycle initiates inter-granular crack formation and the formation of the passivation layer on the NCA surface exposed to the electrolyte. Further cycles promote the growth of inter-granular cracks.

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Figure 5. A comparison of the reduction in the specific discharge capacity for NCA electrodes cycled to different lower cut-off voltages shown as a function of (a) the cycle number and (b) the cycling time shows that mechanical degradation is more important than high voltage chemical degradation pathways. There is less capacity loss for electrodes cycled over a more limited range in the high voltage regime (electrodes A and B) compared to cycling over the full voltage range (electrode C), that is, with larger changes in lattice dimension.

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Table 1. Specific discharge capacity between 2.5-4.5V at select cycles for electrodes cycled to different lower cut-off voltages. The electrodes were cycled in the respective voltage windows for cycles other than those shown in the Table.

Voltage window

Volume change (%)

Cycle number

Elapsed time (hr)

Discharge capacity (mAh/g)

4.25 4.50V

1.6

19

244

203.7

50 98

671 1220

200.9 197.6

20

542

202.0

48

1220

199.0

6

242

203.9

14 18 32

530 674 1190

200.5 197.2 190.7

4.05 4.50V 2.50 4.50V

3.0

5.4

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Figure 2. The discharge capacity of NCA and the X-ray diffraction (XRD) patterns at different states of charge. (a) The specific discharge capacity of the NCA electrode cycled in the voltage window of 2.7-4.5 V at a current rate of 14 mA/g (black dots). The red dots represent the specific discharge capacity, quantified from operando XRD patterns. (b) The (113) reflections at different states of charge obtained from operando XRD. At the end of the 92nd charge, the reflection has a long tail towards lower 2θ angle, which is not evident at the end of the 2nd charge. 118x47mm (300 x 300 DPI)


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Figure 3. The NCA electrode reaction captured by operando XRD during the 92nd and 93rd cycles reveals the presence of two NCA populations within electrode that react at different rates: an active population that retains the behavior of in the original electrode and a sluggish population in which the reaction is retarded. (a) Plots of the (003) and (113) reflections from operando XRD with contour levels of constant intensity (black). (b) Distribution of the Li composition of the NCA electrode based on analysis of the diffraction data. The contour level corresponds to a population density at 0.05. The centroids of the active (dashed green line) and sluggish (dashed purple line) populations are shown. (c) The corresponding voltage profile (red) and the average Li composition (blue) of the electrode quantified from operando XRD. The blank regions corresponds to a brief loss of the synchrotron X-ray beam. 120x82mm (300 x 300 DPI)


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Figure 4. Virtual slices of transmission X-ray tomograms collected for NCA (a) in its pristine state and (b) charged to 4.5 V following 93 cycles (2.7 – 4.5 V, C/20) showing substantial fracturing after extended cycling. (c) A schematic illustrating the process leading to degradation in a secondary NCA particle. The arrow in each primary particle represents the different crystal orientations within the aggregate. The first cycle initiates inter-granular crack formation and the formation of the passivation layer on the NCA surface exposed to the electrolyte. Further cycles promote the growth of inter-granular cracks. 120x87mm (300 x 300 DPI)


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Figure 5. A comparison of the reduction in the specific discharge capacity for NCA electrodes cycled to different lower cut-off voltages shown as a function of (a) the cycle number and (b) the cycling time shows that mechanical degradation is more important than high voltage chemical degradation pathways. There is less capacity loss for electrodes cycled over a more limited range in the high voltage regime (electrodes A and B) compared to cycling over the full voltage range (electrode C), that is, with larger changes in lattice dimension. 84x116mm (300 x 300 DPI)


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Table of Content graphics 76x40mm (300 x 300 DPI)


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Intergranular Cracking as a Major Cause of Long-Term Capacity

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