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Materials by Design: Tailored Morphology and Structures of Carbon Anodes for Enhanced Battery Safety Ryan A Adams, Aashutosh N. Mistry, Partha P Mukherjee, and Vilas G. Pol ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b02921 • Publication Date (Web): 20 Mar 2019 Downloaded from http://pubs.acs.org on March 21, 2019

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ACS Applied Materials & Interfaces

Materials by Design: Tailored Morphology and Structures of Carbon Anodes for Enhanced Battery Safety Ryan A. Adams1§, Aashutosh N. Mistry2§*, Partha P. Mukherjee2† and Vilas G. Pol1‡ 1

Davidson School of Chemical Engineering, Purdue University, West Lafayette, IN 47907, United States

2

School of Mechanical Engineering, Purdue University, West Lafayette, IN 47907, United States

*

ORCID: 0000 – 0002 – 4359 – 4975 ORCID: 0000 – 0001 – 7900 – 7261 ‡ Email: [email protected] § These authors contributed equally to this work. †

ACS Paragon Plus Environment

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Abstract Next generation Li-ion battery technology awaits materials that not only store more electrochemical energy at finite rates but exhibit superior control over side reactions and better thermal stability. Herein, we hypothesize that designing an appropriate particle morphology can provide a well-balanced set of physicochemical interactions. Given the anode-centric nature of primary degradation modes, we investigate three different carbon particles – commercial graphite, spherical, and spiky carbon, and analyze the correlation between particle geometry and functionality. Intercalation dynamics, side reaction rates, self-heating, and thermal abuse behavior have been studied. It is revealed that the spherical particle outperforms an irregular one (commercial graphite) under thermal abuse conditions, as it eliminates unstructured inhomogeneities. A spiky particle with ordered protrusions exhibits smaller intercalation resistance and attenuated side reactions, thus outlining the benefits of controlled stochasticity. Such findings emphasize the importance of tailoring particle morphology to proffer selectivity among multimodal interactions.

Keywords particle morphology, intercalation dynamics, side reactions, thermal abuse tolerance, inhomogeneity


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1. Introduction Despite the long history of commercialization1,2, next generation lithium ion battery technology faces three key challenges: (i) insufficient specific performance (ii) cycle life and (iii) safety risks3– 8.

Scientifically, these directly relate to intercalation dynamics, uncontrolled side reactions, and

internal heat generation. Lithium ion batteries are fundamentally multiscale systems where physicochemical interactions are strongly correlated across multiple length scales9. Traditional approaches have been quite disparate in their treatment of these concerns. For example, thermal management systems are often employed to constrain global temperature rise and represent a celllevel solution. Alternatively, strategies such as electrolyte additives, functionalized binder, and particle coatings have also been employed to alter the behavior of the electrode-electrolyte interface (to mitigate degradation) and/or the bulk electrolyte response (to modulate thermal runaway). For example, polymeric coating of graphite anode enhanced the cycling stability and decreased the first cycle capacity loss by mitigating electrolyte decomposition at the electrode surface10. However, such “add-on solutions” often compromise the primary battery function, i.e., electrochemical energy storage11–18. Active material (e.g., graphite) is responsible for assimilating the electrochemical energy19,20 as well as the irreversible changes21–23 (in part). Hence, if intrinsic active material characteristics are targeted, it has the potential to considerably alter the energy storage landscape24– 33.

Past studies reveal a nonmonotonic dependence on active material dimensions where larger

particles are prone to mechanical degradation arising from intercalation induced stress and smaller nanoscale morphologies promote chemical degradation, for example, excessive solid electrolyte interphase (SEI) formation34–41. Here we investigate three distinct anode particle morphologies (Figure 1(a) - (c)) in terms of their multimodal interactions. Since in state-of-the-art Li-ion cells,


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degradation (intercalation-induced stress, SEI formation, Li plating etc.) is primarily anodecentric, carbon (conventional anode) is studied.

Figure 1: Different carbon-types studied here: (a) commercial graphite (b) spherical carbon and (c) spiky carbon. (b) and (c) are prepared in-house. Their (d) crystalline structure is examined by X-ray diffraction and (e) Raman spectroscopy. (f) Multimodal interactions related to particle morphology. The commercial synthetic graphite has irregular particles (Figure 1 (a)). Such geometrical irregularity fosters nonuniform intercalation dynamics and in turn localized reactions42,43. However, spherical particles (Figure 1 (b)), given the absence of directional preference, should avoid the hot spot formation. A spiky particle (Figure 1 (c)) represents an interesting possibility where controlled irregularity is introduced, while maintaining the spherical symmetry at large. Relevant materials and methods are summarized in the Supporting Information. Specific care has


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been exercised to ensure a proper comparison across these materials, e.g., using comparable electrode active material loading masses. Experimental results are supplemented with mathematical analysis to understand inaccessible aspects such as reaction distribution and hot spot formation.

2. Materials Characterization The morphologies of the three tested carbon materials are revealed via SEM imaging in Figure 1 (a) - (c). Commercial synthetic graphite particles (Figure 1 (a)) showcase irregular and rough surfaces, with widely varying particle shape and size. In comparison, ‘spherical carbon’ particles (Figure 1 (b)), synthesized via autogenic reaction, exhibit smooth surfaces, with an average particle diameter of 2.70 μm ± 0.15 μm. These spherical particles are highly uniform in size and morphology, due to the nucleation and growth mechanism of the corresponding autogenic synthesis44. Carbon particles derived from the pyrolysis of bee pollen (Figure 1 (c)), referred to as ‘spiky carbon’ elsewhere in the text, have an average particle diameter of 14.4 μm ± 0.7 μm, and consist of a spherical core coated with numerous spikes. The designed morphologies of these micron-sized carbon particles were achieved by autogenic reaction of mesitylene to produce spherical carbon and pyrolysis of bee pollen biomass precursor to obtain spiky carbon material, as illustrated in prior studies33,44–46. Given the constraints on particle synthesis, the spherical particle dimensions are different than the other two morphologies. The analysis of intercalation response discussed subsequently assist in isolating the purely morphological effects. Furthermore, the nitrogen sorption isotherms were measured as shown in Figure S1. Table S1 summarizes the as obtained BET (Brunauer-Emmett-Teller) surface area for each material as well as the density functional theory (DFT) pore volume47. Comparable surface areas of 3.8 m2 g-1


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(graphite), 1.5 m2 g-1 (spherical carbon), and 4.5 m2 g-1 (spiky carbon) were observed, with increasing values for the less smooth particle morphologies. The graphite consists of mostly carbon (99 wt. %), whereas the spherical carbon contains 3 wt. % of hydrogen and oxygen heteroatoms from the mesitylene precursor. The spiky carbon consists of 94 wt. % carbon, with oxygen heteroatoms (2.8 wt. %) originating from the biomass source. For carbon anodes, the degree of crystallinity and carbon structuring plays a significant role in the lithium storage mechanism and electrochemical behavior and thus is important to characterize48–50. The XRD pattern for graphite in Figure 1 (d) contains the characteristic (002) peak for the layered graphene sheets at a 2θ of 26.4°, resulting in an inter-lattice spacing of 0.347 nm via Bragg’s law. In contrast, the spherical and spiky carbons demonstrate broad (002) peaks, indicating amorphous carbon characterized by defects, heteroatoms, and sp3 hybridization, as expected from the low pyrolysis temperature51. The (002) peak for spherical carbon is centered at 25.1°, for an average inter-lattice spacing of 0.363 nm, in contrast to 24.5° and 0.371 nm for spiky carbon. Raman spectra were collected for these carbon materials, as shown in Figure 1 (e). The G band 1580 cm-1 correlated to sp2 hybridized carbon and the overtone 2D band at 2700 cm-1 are prominent in graphite as expected. In contrast, the D band at 1330 cm-1 related to sp3 hybridized non-graphitic carbon (due to heteroatoms, edge defects, etc.) is much more prominent for the spiky and spherical carbon.

3. Electrochemical Implications To compare the electrochemical performance of these carbon materials, lithium ion half cells were constructed, with the charge-discharge voltage profiles for the first two cycles (C/10 rate) shown in Figure 2 (a)-(c). The graphite exhibits its characteristic low voltage intercalation plateau and a reversible capacity of 345 mAh g-1 (approaching the theoretical capacity of 372 mAh g-1). In


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comparison, spherical and spiky carbon anodes display sloping voltage profiles and reversible capacities of 300 mAh g-1 and 240 mAh g-1, respectively. The first cycle Coulombic inefficiency (demonstrated by the red arrow) is attributed to SEI layer formation on the carbon anode surface. This occurs due to electrochemical degradation of surrounding electrolyte on the lithiated graphite surface until quasi-equilibrium is established (often referred to as the formation cycle capacity loss). On the other hand, SEI growth upon repeated cycling later is related to an electrochemical reaction. Graphite and spherical carbon show similar irreversible capacities of 19% and 28% respectively, while spiky carbon has a higher value of 50%. The first cycle capacity loss is known to depend on surface area as well as carbon structuring, favoring increased SEI layer formation for amorphous carbon materials.52–54 Thus, the much higher first cycle loss of spiky carbon as compared to graphite can be attributed to its increased surface area, heteroatoms, and hard carbon structuring.


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Figure 2: Electrochemical trends upon cycling. Initial cycles refer to the formation of a stable SEI film and accordingly large capacity variations are seen across all carbon types (a) graphite (b) spherical and (c) spiky. These tests are carried out at C/10. (d) Constant current cycling quantifies the capacity decay due to SEI growth on repeated cycling. (e) Capacity retention trends differ as SEI formation and intercalation dynamics vary across different particle types Electrochemical impedance spectra of (f) graphite, (g) spherical carbon, and (h) spiky carbon. Cells were aged at 50°C for 100 cycles at 1C rate, and EIS was collected at 200 mV (vs. Li/Li+). Intercalation stability over repeated cycling is further probed using constant current cycling at a C/2 rate as shown in Figure 2 (d). Close to 100% Coulombic efficiencies are observed, pointing to structural stability. By cycle 100, specific capacities of 240 mAh g-1, 250 mAh g-1, and 230 mAh g-1 are obtained for graphite, spherical carbon, and spiky carbon, respectively. A separate rate study was performed, with the capacity retention (normalized to C/10) plotted in Figure 2 (e). The spiky carbon exhibits consistently higher rate capability than spherical carbon. Further full cells with


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LiCoO2 cathode are tested to identify how anode morphological variations translate to a full cell setting, with similar trends observed in Figure S2. Electrochemical impedance spectroscopy (EIS) analysis was performed for each carbon material before and after cell aging (Figure S3), with spectra shown in Figure 2 (f) - (h). Modeling to an equivalent circuit (Figure S4) revealed that both the graphite and spherical carbon materials saw a rise in cell impedance after aging, stemming from the charge-transfer resistance and SEI resistance components (Table S2). Comparatively, the spiky carbon material exhibits minor change, indicating that the surface protrusions enhance lithiation kinetics and efficiency in harsh cycling conditions.

Figure 3: A spiky particle has two attributes (a) spike amplitude and (b) spike count. A spherical particle is a special instance of a spiky particle with vanishingly small spike amplitude. (c) Such geometrical features bring about complexities in the intercalation response which in turn affects side reaction rates. As revealed by Figure 2(b) and (c) (as well as Figure 1(d) and (e)) the intrinsic material characteristics for the two are similar, and in turn the electrochemical differences between the two in Figure 2(d) and (e) largely stem from geometrical features such as (i) different intercalation lengths (ii) morphology. To understand this intricate correlation between surface features and intercalation dynamics, electrochemical dynamics was modeled for the spiky particles (Figure 3). These spikes (surface features) have two distinct attributes – length and count. Varying these two,


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one can obtain a range of particulate structures as shown in Figure 3 (a) and (b). Figure 3 (c) compares potential evolution during intercalation for a spherical (smooth) and a spiky particle. These intercalations are carried out at identical rates (4C). A high rate (here 4C) is chosen to demonstrate the severity of these effects. The potentials are less than the open circuit potential (OCP) for a given state of (overall) lithiation as the intercalation flux is inwards. Both the profiles qualitatively match the OCP, with the potential of a spiky particle being closer to OCP. This indicates better electrochemical utilization, in terms of both the voltage as well as lithium storage. Carbon-based electrodes face two distinct modes of chemical degradation: i.

Solid electrolyte interphase (SEI) formation – 0.2 V reaction potential

ii.

Lithium plating – 0.0 V reaction potential

At any instant, the following identity holds (integration is performed over particle surface):

I app    iintercalation  iside  dS

(1)

S

where intercalation current refers to lithiation (a reversible reaction):

iintercalation  k Cs Ce  C

max s

F  F  U   F  U     2FRT  0 2 RT  Cs  e  e   iintercalation e 2 RT  e 2 RT     

(2)

and SEI formation reaction – an irreversible side reaction: 0 iside  iSEI e



F  U SEI  2 RT

(3)

Since the OCP for SEI reaction (Eq. (3)) is around 0.2 V, as soon as the particle potential drops 0 below 0.2 V, the SEI layer starts forming. The exchange current density, iSEI , and OCP, U SEI , are

constants. Additionally, given the high electrical conductivity of carbon-based materials, no appreciable gradient exists in the solid phase potential,  . These assumptions amount to uniform SEI formation at different locations on the particle surface and the respective flux changes over


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time due to variation in the particle potential,  . Figure 3 (c) also sketches this side reaction flux during intercalation. The total side reaction current is I 2   iside dS . As the potential for the S

spherical particle is consistently lower than that for the spiky particle, the side reaction flux is higher for the spherical particle (Figure 3 (c)). Moreover, since the spherical particle voltage drops to zero earlier, the onset of Li plating is earlier for the smooth spherical geometry. This is quite interesting since the introduction of these ordered geometrical features appears to improve intercalation behavior as well as subside the side reactions. Note that the two particles have similar volumes so the observed trends in Figure 3 (c) clearly stem from morphological differences. These findings match with Figure 2 (e), where experimentally the spherical particle consistently underperforms compared to the spiky one. Also note that Figure 2 (b) and (c) quantify the initial (chemical) SEI formation, while SEI reaction (3) corresponds to capacity fade over repeated cycling. The capacity loss during formation cycling is a more complex phenomena, and accounts for the differences in Figure 2 (d), while Figure 2 (e) presents intercalation rate dependence once the formation cycles are over (i.e., normalized with respect to the stabilized capacity). To further understand the intercalation dynamics of a spiky particle, surface and bulk fields are sketched in Figure 4. First, consider the concentration distribution on the particle surface at 20% lithiation, as shown in Figure 4 (a). Surface points are identified by a dimensionless indicator where -1 refers to the valley, 0 to mean surface and 1 to peak locations. In general, the peak locations intercalate more than the valley points. This points to preferential intercalation as peak locations locally have more surface area available for reactions. Interestingly, not all the points at identical radial distances have identical intercalation. This is because the lateral dimensions of the spikes are not identical and vary based on their surface location (refer to Figure 4 (c) inset). This gives rise to overlapping lithiation fronts and breaks the symmetry. Next, refer to the local


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0 intercalation flux expression (2), where both the exchange current density, iintercalation , and open

circuit potential, U , are dependent on local surface concentration. Also, given the high electronic conductivity, the entire particle is at a constant electric potential,  (less than the smallest OCP during lithiation). At a given lithiation extent, the current distribution is dictated by the OCP variation as exponential terms represent a stronger nonlinearity (Eq. (2)). As shown in Figure 3 (c), OCP decreases with lithiation. Hence, valley locations have a higher OCP and in turn a greater overpotential. This leads to a higher intercalation flux in valleys, with Figure 4 (b) showing this distribution. As found in the concentration distribution, the intercalation flux also exhibits a nonmonotonic dependence on the radial location of the surface point. A clearer trend is obtained if intercalation flux is plotted against surface concentration (Figure 4 (c)), where locations that are less intercalated have a higher intercalation flux. Figure 4 (c) to (e) demonstrate the evolution of this flux – concentration relation with lithiation. Less intercalated locations consistently have a higher flux but the qualitative nature changes. As a reference to the intercalation extent, Figure 4 (f) presents the potential and side reaction flux evolution during the respective operation. Figure 4 (g) to (j) provide a visual picture of the intercalation front propagation in the interior portions of the particles. The normalized max lithium concentration, x  C / C , is plotted on vertical and horizontal planes. It appears that the

intercalation front qualitatively follows a shape similar to the surface features. Note that these calculations only reveal the active material centered phenomena. In typical porous electrodes, secondary species as well as the electrolyte pore network contribute to overall dynamics19,55–58. Revisit Figure 3 (c) where the potential profiles largely differ during the transition stages for graphite. Graphite characteristically has three voltage plateaus, each of which correspond to a particular lithiation window. For a spherical particle, surface concentration is uniform, and the


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voltage transitions take place simultaneously for all the points on the surface (recall that the OCP is a function of surface concentration). On the other hand, for a particle with surface protrusions, preferential intercalation takes place, and consequently about the transition points on the OCP profile (Figure 3 (c)), surface inhomogeneity of the state of lithiation causes a prolonged transition. When the all the lithiation points are in the same plateau, the behavioral difference originates from difference in surface area. In either of the stages, the disparity grows with the operating current.

Figure 4: Intercalation nonuniformity due to geometrical features. (a) Concentration and (b) intercalation flux distribution on particle surface at 20% lithiation. (c) Corresponding intercalation flux as a function of concentration and surface map (inset). (d) and (e) show the progression of flux – concentration maps at higher lithiation extents, 40%, and 80%, respectively. (f) Potential evolution and contribution from side reaction. Concentration evolution upon lithiation in the interior of the particle (g) initial (h) 20% (i) 40% and (j) 80% lithiation. Vertical and horizontal sections are shown in (g) through (j). 4. Inhomogeneity in Self-heating When an intercalation reaction takes place, a portion of the reaction enthalpy gets stored electrochemically (correlated to OCP evolution with lithiation) and part dissipates as heat. This remaining portion is referred to as the reversible heat and is quantified as:


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qreversible  iintercalationT

U T

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(4)

Additionally, to sustain a finite rate of electrochemical reactions, a penalty is incurred in the form of reaction overpotential that eventually manifests as heat generation at the active interface:

qkinetic  iintercalation

(5)

Figure 5: Progression of (a) – (d) local overpotential (e) – (h) kinetic and (i) – (l) reversible heat generation rates as lithiation takes place. The inhomogeneity translates to all the different fields and is observable during intermediate stages of lithiation. For comparison, states of an equivalent spherical particle are also shown. A spherical particle has a uniform surface state.


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Intercalation flux  ii ntercalation  , entropic coefficient  U / T  , and overpotential   all depend on the local surface concentration, and in turn the heat generation rates also exhibit a spatial dependence. Figure 5 presents this progression with lithiation. First, consider the evolution of overpotential as shown in Figure 5 (a) to (d). At the start of intercalation, the concentration field is uniform and in turn, a constant overpotential exists (Figure 5 (a)). Initially, the exchange current density is very low due to a quadratic dependence on lithium concentration (Eq. (2)). This leads to a fairly high reaction overpotential (~50 mV). Exchange current density reaches its maxima at 50% lithiation, hence going from Figure 5 (a) to (c), the average color shifts towards smaller overpotentials. The overpotential inhomogeneity is visible in these maps. Again beyond 50% lithiation, exchange current density decreases and equivalently average overpotential rises (Figure 5 (d)). The remaining plots in Figure 5 demonstrate the evolution of kinetic heat (Figure 5 (e) to (h)) and reversible heat (Figure 5 (i) to (l)). For comparison, corresponding heat generation rate distribution of an equivalent spherical particle is also sketched simultaneously. As mentioned earlier, at the onset of intercalation, the concentration field is uniform and in turn, every other field variable is spatially invariant. Hence, the respective data points overlap in Figure 5 (e) and (i). As lithiation proceeds, heat generation rates for the spiky particle exhibit spatial variations, while spherical particle continues to exhibit unique values (Figure 5 (f) to (h) and Figure 5 (j) to (l)). Given the nonlinear coupling between the lithiation fields and the heat generation modes, the average heat generation rate in the spiky particle is different than the singular value for the spherical particle. Also, the peak or the valley locations do not necessarily demonstrate the extreme behavior, since as the intercalation progresses, diffusion fronts propagate and break such an ordering. Note that the spread in heat generation rate values changes non-monotonically with the extent of lithiation. A spiky particle does not necessarily experience a higher heat generation than


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its spherical counterpart during intercalation (Figure 5 (g) vs. (h)). The evolution of kinetic and reversible heat generation rates is also nontrivial. Such complexities give rise to behavioral transitions as the rate of operation is varied since qkinetic and qreversible scale differently with C-rate. As the present discussion focuses on the particle scale thermal response, transport heat related with the electrolyte transport resistance in the pore-network is not discussed here20,59. The self-heating signature for the spiky particle, quantitatively described as Qparticle = Qkinetic + Qreversible, exhibits a quadratic departure from the corresponding (equal volume) spherical particle. At the start of lithiation, the concentration distribution is uniform and in turn little disparity exists between the two in Figure 5 (e) and (i). The interfacial stochasticity present in the spiky particle causes nonuniform intercalation resistance, which is apparent through the preferential intercalation patterns. This effectively leads to distributions in both the heat generation terms, as is apparent in Figure 5 (f), (g), (j), and (k). The reversible heating can become negative based on the direction of entropy exchange and give rise to situations where the total particle heat for a spiky particle is less than the equivalent spherical active material (Figure 5 (g), (k) in contrast to (f), (j)). Towards the end of lithiation, when most of the low intercalation resistance locations are nearly lithiated, the lagging spots get filled predominantly. Such a reversal in lithiation preference is responsible for the decreased field inhomogeneities in Figures 4 (e) and 5 (d), (h), (l).

5. Thermal Abuse Behavior The previous discussion compared the self-heating60 rates of different particle morphologies. On the other hand, there are certain thermally activated side reactions that can trigger at higher temperatures. Given their exothermic nature, if allowed to carry on, these reactions prove to be autocatalytic where heat generated due to one reaction increases the local temperature and further


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triggers another reaction. In the worst case, if the generated heat is not dissipated quickly, it leads to extremely high temperatures and possible cell combustion (often called thermal runaway). Safety concerns for lithium ion batteries arise from such poor thermal stability of electrode and electrolyte materials61,62. Typically, thermal runaway initiates from degradation reactions occurring on the graphite anode, where metastable species in the SEI undergo exothermic breakdown, initiating as early as 90°C63–65. Hence, the anode’s thermal stability, in turn, affects the thermal runaway response of the entire cell.

Figure 6: (a) Differential scanning calorimetry (DSC) captures the thermally activated reactions (endothermic and exothermic) of active material. This reaction information is further useful in estimating thermal abuse response of cells composed of the same active material. (b) DSC responses of the three carbon materials and corresponding (c) oven test trends. Note that a larger spike in the heat flow does not necessarily mean poorer thermal stability as reaction rate constant can change the rate of heat generation.


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In order to characterize these thermally activated reactions for the three carbon types, differential scanning calorimetry (DSC) analysis was carried out. Figure 6 (c) presents the DSC profiles for these materials, with two characteristic exothermic peaks shown. The first one corresponds to the decomposition of metastable components in the SEI. At higher temperatures, the intercalated lithium starts reacting with the electrolyte in another exothermic reaction. DSC information is somewhat difficult to interpret, as the temperature continuously increases (at a constant rate) and in turn, the next reaction may get triggered before the previous reaction proceeds to completion. Given this complexity, past studies resort to accelerated rate calorimetry (ARC) in order to fingerprint these thermally activated reactions. We have developed an inverse problem formulation to consistently interpret DSC data. Figure 6 (a) shows the DSC data for commercial graphite, along with the interpreted trend based on the inverse analysis. This gives us reaction information, essentially, the rate constant, activation energy (related to onset temperature) and reaction heat. While DSC comments on an individual material’s thermal instability, this does not translate monotonically to the battery scale. For example, a battery with a lot of extra inert material will absorb all the heat generated due to thermally activated reactions and would never initiate the thermal runaway. Given the sufficient information regarding the abuse reactions, the cell-level thermal abuse tolerance can be predicted66. Figure 6 (a) shows the thermal runaway response of a fully charged cylindrical cell with graphite anode and lithium cobalt oxide LiCoO2 cathode. Similar predictions have been carried out for other carbon materials by properly interpreting their DSC data (Figure 6 (b)). Figure 6 (c) presents the thermal abuse response of these materials. For a consistent comparison, identical active material loadings and cell dimensions are used. In other words, the observed differences purely stem from different carbon-types and is not contaminated by unequal loadings for example. Figure 6 (c) reveals that the spherical carbon is the safest


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material as the smallest amount of heat is generated as well as the onset temperature for the thermal ramp is the highest. Also, notice a small kink in the temperature profiles (Figure 6 (c)) for the spherical and spiky carbons. This results from an appreciable heat generation during SEI decomposition (visible in Figure 6 (b)). But since the trigger temperatures for the second reaction between the intercalated Li and electrolyte are much higher, this does not trigger an early onset thermal runaway. In other words, sharper and taller peaks on DSC profile does not necessarily mean an inferior thermal stability. Thermal stability relies not only on reaction enthalpies but also on reaction rates and activation energies.

6. Conclusions Active material, the phase responsible for electrochemical energy storage, exhibits varying particle geometries. The physical aspects of the active particles dictate the distribution of intercalation flux and associated physicochemical complexations. Here we probe such morphological dependence of multimodal interactions responsible for observable battery characteristics, namely, performance, degradation and safety. Using complimentary experimental and analytical tools, we studied three distinct anode particle types: (i) irregular graphite, (ii) spherical carbon and (iii) spiky carbon. The irregular graphite exhibits poor rate dependence given the nonuniformity caused by particle geometry, as compared to a spherical carbon. Introducing controlled disorder, i.e., a spiky particle, results in a noncanonical response where improved particle – electrolyte contact reduces the overpotential, enhances intercalation and attenuates the side reactions leading to better performance and slower capacity fade. Consequently, the self heating response also scales nonmonotonically with the particle surface features. However, the spiky particle demonstrates


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marginally poor thermal abuse response. In summary, tailored particle morphology provides a unique avenue to tuning particle scale energy storage through geometrical alterations.

Supporting Information Materials, methods, and additional results are presented in Supporting Information. This includes: S1. Materials S2. Electrode Preparation and Electrochemical Testing S3. Thermal and Materials Characterization S4. Interpreting DSC Trends 34,67,68 S5. Electrochemical Dynamics S6. Thermal Abuse Predictions 69–71 Acknowledgments The A-C010 LiCoO2 cathodes were produced at the U.S. Department of Energy’s (DOE) CAMP (Cell Analysis, Modeling, and Prototyping) Facility, Argonne National Laboratory. The CAMP Facility is fully supported by the DOE Vehicle Technologies Program (VTP) within the core funding of the Applied Battery Research (ABR) for Transportation Program. VP and PPM wish to thank the Office of Naval Research (ONR) for supporting lithium ion battery safety work under the grant N00014-18-1-2397, and Dr. Michele Anderson from ONR for supporting this work. RA thanks the Bilsland Dissertation Fellowship for support. References (1) (2) (3) (4) (5) (6) (7) (8) (9) (10) (11) (12) (13)

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