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Daniela Molina Piper , Thomas A. Yersak , Seoung-Bum Son , Seul Cham Kim , Chan Soon Kang , Kyu Hwan Oh , Chunmei Ban , Anne C. Dillon , and Se-Hee Lee *

Conformal Coatings of Cyclized-PAN for Mechanically Resilient Si nano-Composite Anodes

By cyclizing commercially available polyacrylonitrile (PAN), we show that it is possible to conformally coat nanoparticles of Si with a conjugated polymer. We utilize cyclized-PAN both as a binder and conductive additive because of its good mechanical resiliency to accommodate silicon’s (Si) large expansion as well as its good ionic and electronic conductivity. By the 150 th cycle, our nano-Si/cyclized-PAN composite anodes exhibit a specifi c charge capacity of nearly 1500 mAh g − 1 with a coulombic effi -ciency (CE) approaching 100%. Because Si is naturally abundant and has such a high achievable specifi c capacity, [ 1 , 2 ] the next generation of Lithium-ion batteries will inevitably incorporate an advanced Si based anode like that presented here. At room temperature, Si can accommodate 3.75 mole Li per mole of Si (Li 15 Si 4 ) for a theoretical capacity of 3579 mAh g − 1 . [ 3 , 4 ] Despite these advantages, progress towards a commercially viable Si anode has been impeded by Si’s rapid capacity fade, poor ionic transport and low CE.

Rapid capacity fade is a primary cause for the delay of Si’s commercialization. At room temperature, a volume expansion of 300% occurs upon full lithiation to Li 15 Si 4 . [ 5–7 ] Such a mas-sive volumetric change can result in cracking and pulveriza-tion of the Si particles, which then leads to the interruption of electronic transport pathways and the electrochemical isola-tion of pulverized particles. [ 2 ] To better accommodate the large stresses and strains generated upon lithiation, utilization of nanoparticles, nanowires, and 3D nano-porous structures have been studied. Nano-Si based structures have the added benefi t of reducing the average Li + diffusion length into Si for faster charge and discharge rates. [ 8–10 ] However, nano-Si structures often suffer from poor CE due to the continual formation of a solid electrolyte interphase (SEI) layer at the large nano-Si/elec-trolyte interface. Adding to the frustration, many of these clever nano-Si based engineering solutions are rarely appropriate for

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D. Molina Piper, T. A. Yersak, S.-B. Son, Prof. S.-H. Lee Department of Mechanical Engineering University of Colorado at Boulder, 80309 USA E-mail: sehee.lee@colorado.edu D. Molina Piper, Dr. C. Ban, Dr. A. C. DillonNational Renewable Energy Laboratory 1617 Cole Boulevard, 80401 USA S.-B. Son, S. C. Kim, C. S. Kang, Prof. K. H. OhDepartment of Material Science and Engineering Seoul National University, 151-742 Korea

DOI: 10.1002/aenm.201200850

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commercialization because they may require expensive syn-thesis techniques. [ 11–15 ]

Yet, recent work has shown that Si-C nano-composites may be promising candidates for viable, inexpensive, stable and effi cient high capacity anodes. Si-C conventional composites are typically prepared by carbonizing precursors [ 16–18 ] or by mechanically mixing Si with carbon. [ 19 , 20 ] The result are com-posites of Si particles embedded in carbon matrices. Unfor-tunately, these carbon matrices cannot accommodate Si’s large volumetric changes because of their tendency for brittle failure. Cracking of these carbon matrices during cycling interrupts electronic conduction pathways and exposes fresh Si to the electrolyte which results in further SEI formation. Several studies using electronically conductive polymer based matrices or coatings have demonstrated better results. [ 21–25 ] For example, G. Liu et al. developed a cathodically (n-type) doped polymer binder that maintains good electronic conduc-tivity when subjected to the reducing environment of anodes (0.01–1 V vs. Li/Li + ). The authors abstained from using com-mercially available polymers because these polymers have a p-type conductive character that is thought to disappear in the anodic potential range. [ 26 ] Despite G. Liu et al.’s fi ndings, we found it possible to develop conformal coatings using the com-mercially available polymer, PAN, as a precursor. By limiting the pyrolysis temperature to 300–500 ° C in an inert environ-ment, the cyclization of PAN proceeds without carbonization. Limiting the pyrolysis to only cyclization maintains PAN’s polymeric properties while still introducing delocalized sp 2 π bonding for intrinsic electronic conductivity. As we will show, the superior performance of our nano-Si/cyclized-PAN com-posite electrodes is enabled by the unique material properties of cyclized-PAN coatings.

To confi rm the cyclization of PAN we utilized differential scanning calorimetry (DSC), X-ray photoelectron spectroscopy (XPS) and Raman spectroscopy. During this part of the study, we focused our attention upon thick (1–5 μ m) PAN fi lms with no active material component. It is not appropriate to charac-terize the mechanical properties of these fi lms because these results would not be representative of our thin ( ∼ 5 nm) PAN conformal coatings on actual nano-Si particles. However, these thick micron fi lms are more than adequate for understanding the molecular structure of our thinner coatings.

DSC results reveal an exothermic peak at 321.7 ° C with an onset temperature at about 296.9 ° C (Figure S1). This peak is consistent with the cyclization of PAN’s nitrile groups at tem-peratures between 200–350 ° C in an inert atmosphere. [ 27–29 ]

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Figure 1 . (a) Structural representation of the cyanic group (N1), pyridinic group (N2), and substitutional graphite group (N3). These groups account for the π resonances observed in our samples’ N 1s spectra. (b) N 1s spectra of PAN fi lms treated at temperatures of 25–500 ° C. For treatment temperatures ≥ 300 ° C, the spectra broadens according to the presence N2 and N3 π resonances. (c) Deconvolution of the 300 ° C sample’s N 1s spectra includes only the N2 and N3 π resonances and confi rms the cyclization of PAN. (d) Deconvolution of the 500 ° C sample’s N 1s spectra indicates that the N3 π resonance grows in area and that PAN samples treated at high temperatures begin to “graphitize.”

Because we heat treat at such a low temperature in Argon for an extended period of time, PAN will not carbonize. Instead, it develops a pyridine-based conjugated polymer.

To confi rm the presence of nitrogen domains we use XPS ( Figure 1 ). Depending on the sample, deconvolution of the N 1s spectra may give different π resonances based on the particular bonding environments of nitrogen present. For the cyanic group (N1, C�N) there is a resonance at 399.0 eV, while the pyridinic group (N2, C�N = C) has a resonance at 398.4 eV, and the substitutional graphite group (N3, N coordinated with three C atoms) has a resonance at 399.9 eV (Figure 1 a). [ 30–32 ] The N 1s spectra for samples of 1–5 μ m thick PAN fi lms treated at 25, 100, 200, 300, 400 and 500 ° C are presented in Figure 1 b. The N 1s spectra indicates the presence of only cyanic groups for treat-ment temperatures ≤ 200 ° C, but broadens to include pyridinic and substitutional graphite groups for treatment temperatures ≥ 300 ° C. The disappearance of the cyanic group π resonance at treatment temperatures ≥ 300 ° C is consistent with our DCS results that indicate the onset of cyclization occurs at 296.9 ° C. An optimized Gaussian-Lorentzian fi t of the 300 ° C sample’s N 1s spectra excludes the cyanic group suggesting complete cycli-zation of the sample (Figure 1 c). The ratio of the N2 to N3 areas for this sample is 0.79 and the N3 π resonance has a maximum at 399.6 eV, which is 0.3 eV lower than reported values. [ 30–32 ] We conclude that this sample is likely comprised of pyridine

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ring chains with a signifi cant degree of cross-linking. [ 30 ] When treatment temperatures are increased beyond 300 ° C the N3 π resonance continues to broaden. At 500 ° C, the ratio of the N2 to N3 areas decreases to 0.68 and the N3 π resonance maximum increases to 399.9 eV (Figure 1 d). A more intense N3 π resonance indicates the evolution of a more graphite-like substitutional structure. If treat-ment temperatures are increased beyond 500 ° C, we expect that the N3 π resonance would come to dominate the N 1s spectra and eventually disappear completely as car-bonization progresses to completion. [ 34 , 35 ] At high temperatures, carbonization and graphitization of PAN will occur and our coatings would lose the elastic properties we require to accommodate Si’s volumetric expansion. Such high heat treatments are known to develop materials with up to 99% carbon content and have been used to syn-thesize other Si-C composites. [ 18 , 22 , 36 ]

A pyridine-based conjugated polymer sug-gests that our cyclized-PAN fi lms will have intrinsic electronic conductivity. To provide further evidence for cyclized-PAN’s delocal-ized sp 2 π bonding directly, we use Raman spectroscopy. Figure 2 a presents the Raman spectra for micron thick fi lms of untreated-PAN, PAN cyclized at 300 ° C, and PAN cyclized at 500 ° C. The Raman spectrum for the untreated-PAN (orange profi le) exhibits no peaks, whereas the spectra for PAN fi lms cyclized at 300 ° C (cyan profi le) and

500 ° C (black profi le) exhibit two Raman shifts at 1600 cm − 1 (G band) and 1360 cm − 1 (D band). These two bands are attrib-uted to delocalized sp 2 π bonding. [ 37 ] The observation of both the D and G bands attests to the existence of disordered and ordered structural confi gurations, respectively. [ 38 ] To charac-terize the degree of order we compared the relative intensity of the G and D bands. To do this, each peak was fi tted with a Gaussian-Lorentzian function [ 39 ] (Figure 2 b and 2 c). The ratio of the D to G band intensities (I D /I G ) was calculated using the fi tted data and found to be 2.66 and 2.50 for PAN cyclized at 300 ° C and 500 ° C, respectively. Consistent with the litera-ture, we have found that I D /I G decreases with increasing heat treatment temperature. [ 40 , 41 ] A stronger relative G band for the sample treated at 500 ° C indicates a higher degree of order than that for the sample treated at 300 ° C. Like graphite, pyridine rings have delocalized sp 2 π bonding which enables good elec-tronic conductivity. [ 42 , 43 ] Because the sample treated at 300 ° C has a higher I D /I G ratio than the samples treated at 500 ° C, we choose to focus the remainder of our material characterization on samples cyclized at 300 ° C.

To characterize the microstructure of our nano-Si/cyclized-PAN electrodes we used transmission electron microscopy (TEM) and electron energy loss spectroscopy (EELS). TEM sam-ples were prepared using a focused ion beam (FIB) equipped with a mobile air-lock chamber. [ 10 ] Micrographs of the TEM sample

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Figure 3 . (a) TEM micrograph of the uncycled nano-Si/cyclized-PAN electrode with (b) EELS elemental mapping confi rming that we achieve a conformal 5 nm thick cyclized-PAN coating. (c) Raman spectra for an untreated electrode (red) and for an electrode treated at 300 ° C (cyan). The delocalized sp 2 π D and G bands present in the spectra for the heat-treated sample confi rm cyclization of PAN coatings.

Figure 2 . (a) Raman spectra of untreated PAN fi lms (orange), PAN fi lms treated at 300 ° C (cyan) and PAN fi lms treated at 500˚C (black). The observation of graphite D and G bands confi rms delocalized sp 2 π bonding and a cyclic structure for heat-treated samples. (b), (c) Fitted D and G bands used to calculate the I D /I G ratio for the 300 ° C and 500 ° C samples, respectively. A smaller ratio at 500 ° C (2.50) than at 300 ° C (2.66) indicates that a more ordered structure is achieved at higher heat treatment temperatures.

preparation by Ga + ion milling are presented in Figure S2. Figure 3 also presents a TEM micrograph (Figure 3 a) and an EELS (Figure 3 b) elemental map of an uncycled nano-Si/cyclized-PAN electrode. We observe a conformal thin coating ( ∼ 5 nm) of cyclized-PAN (cyan) on nano-Si particles (red). The coating provides an intimately linked conductive network that connects nano-Si particles throughout the electrode. To con-fi rm that the PAN coating contains delocalized sp 2 π bonding, we used Raman to characterize our electrodes. Figure 3 e pre-sents the Raman spectra for nano-Si electrodes with coatings of untreated PAN (red profi le) and PAN cyclized at 300 ° C (cyan profi le). As expected, both samples exhibit a shift attributed to Si at 520 cm − 1 . And like the Raman analysis of 1–5 μ m PAN fi lms, only the electrode treated at 300 ° C exhibits the D and G delocalized sp 2 π bands. From this result we conclude that nano-Si is coated with an electronically conductive and mechan-ically resilient cyclized-PAN conformal coating.

Results of our electrochemical characterization are presented in Figure 4 . Figure 4 a presents the cyclic stability of our nano-Si/cyclized-PAN anode treated at 300 ° C. The electrode was run at a rate of C/20 for the fi rst 10 cycles and a rate of C/10 for all subsequent cycles. At cycle 150 our nano-Si/cyclized-PAN

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cell exhibits a specifi c charge capacity of nearly 1500 mAh g − 1 with a CE approaching 100%. 1500 mAh g − 1 corresponds to an electrode volumetric capacity of 620 mAh cm − 3 considering an

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Figure 4 . (a) Cyclic capacity (blue circles) and CE (blue squares) of nano-Si/cyclized-PAN electrodes run at a rate of C/20 for the fi rst 10 cycles and at C/10 for subsequent cycles. A conventional nano-Si electrode (red circles) was cycled at a rate of C/10 as a control cell. After 150 cycles, our nano-Si/cyclized-PAN composite electrode has a specifi c charge capacity of nearly 1500 mAh g − 1 and a CE approaching 100%, whereas the con-ventional electrode fails after 3 cycles. (b) Voltage profi les of our nano-Si/cyclized-PAN electrode showing a minimal overpotential when cycled at C/10 compared to C/20. (c) A rate test demonstrates that our nano-Si/cyclized-PAN electrodes can achieve a specifi c charge capacity in excess of 2300 mAh g − 1 at a rate of 5C.

initial electrode thickness of 9.08 μ m ( Figure 5 a). Such good cycle life and CE is evidence that the cyclized-PAN coating has good mechanical resiliency. Figure 4 b displays the 1 st , 3 rd , 10 th , 30 th , 40 th , and 50 th voltage profi les of our nano-Si/cyclized-PAN anode. The fi rst cycle profi le shows a specifi c charge capacity of 2585 mAh g − 1 , which is equivalent to the extraction of 2.7 mole Li per mole of Si. At cycle 50, the specifi c charge capacity is 2078 mAh g − 1 (2.2 mole Li).

In lieu of a direct electrical conductivity measurement of the electrode, we used AC-impedance spectroscopy to measure its ohmic resistance. It was found to be 4.10 Ω , from which 1.12 Ω

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is attributed to our nano-Si/cyclized-PAN electrode fi lm and the rest results from the liquid electrolyte (Figure S3). We also conducted a rate test to study the transport properties of our cyclized PAN fi lms (Figure 4 c). At a rate of 5C, our nano-Si/cyclized-PAN electrode exhibits a specifi c charge capacity in excess of 2300 mAh g − 1 . If the coating did not provide adequate ionic and electronic transport, it is reasonable to conclude that we would not have been able to achieve such high charging capacities at a rate of 5C. At this time, an ionic transport mech-anism for the cyclized-PAN coatings has yet to be elucidated. However, other amorphous, thin coatings have been shown to be ionically conductive despite being ionically resistive in the bulk form. [ 44 ]

We characterized the microstructural properties of our nano-Si/cyclized-PAN electrodes using TEM to verify the mechanical resiliency of our cyclized-PAN coating. TEM samples were prepared using the same FIB procedure as mentioned before. Figure 5 a and 5 b present TEM micrographs of electrode cross-sections before and after initial lithiation, respectively. Before cycling, the electrode has good porosity and an overall thick-ness of 9.08 μ m. After full lithiation at a rate of C/20, the elec-trode expands in overall thickness to 12.75 μ m. Si’s expansion is largely accommodated by the porosity of the electrode such that the electrode only expands by 40%. However, the porous structure reopens upon delithiation after 20 cycles (Figure 5 c) indicating that ionic transport pathways are maintained.

To confi rm that cyclized-PAN coatings conform to the strains of lithiated Si, an electrode was recovered after its 20 th cycle. This electrode was cycled at a slower rate of C/30 for the fi rst cycle and C/20 for subsequent cycles to fully lithiate Si and expose the cyclized-PAN ∼ 5 nm coating to a condition of maximum stress and strain. Even after extensive cycling, the cyclized-PAN coating and the Si nanoparticles display no evi-dence of cracking nor does the cyclized-PAN/Si interface show signs of delamination (Figure 5 d and 5 e). The superior resil-iency of our cyclized-PAN coating successfully accommodates Si expansion where previously reported dense carbon matrices would have experienced brittle failure. We believe that the con-formal cyclized-PAN coating isolates nano-Si from the organic liquid electrolyte and inhibits nano-Si aggregation during cycling. Isolation of nano-Si from the electrolyte prevents the parasitic formation of a SEI and improves CE. Material aggre-gation is a frequent cause of capacity fade in electrodes utilizing nano-particles, [ 45 , 46 ] but our cyclized-PAN coating confi nes nano-Si for good capacity retention.

Cyclized-PAN conformal coatings address several chronic issues that have impeded the commercialization of Si-based electrodes. When designing the coating, we wished to have a material with elastic polymeric properties as well as good con-ductivity. By cyclizing PAN at temperatures between 300–500 ° C and avoiding carbonization at temperatures > 800 ° C, we obtain a pyridine-based conjugated polymer that accommodates Si’s volumetric expansion during lithiation. Pyridine also has delo-calized sp 2 π bonding for intrinsic electronic conductivity. Good ionic conductivity of cyclized-PAN coatings is assumed based upon good electrochemical performance at fast cycling rates. And for increased electrode energy density, cyclized-PAN coat-ings serve as both conductive additive and binder singly. Our approach is attractive for commercialization for two reasons.

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Figure 5 . (a), (b) TEM micrographs of a nano-Si/cyclized-PAN electrode cross-sections before and after initial lithiation, respectively. Si’s expansion is largely accommodated by the porosity of the electrode such that the electrode only expands by 40%. (c) The porous structure reopens upon delithia-tion after 20 cycles indicating that ionic transport pathways are maintained. (d), (e) TEM micrographs of a delithiated nano-Si/cyclized-PAN electrode collected after its 20 th cycle. We observe that the cyclized-PAN coating does not delaminate from Si particles nor does it crack.

First, our coating is based upon a commercially available and low-cost polymer precursor. And second, the process described is compatible with commercial electrode slurry manufacturing methods and is adaptable to other anode or cathode materials.

Experimental Section Material Preparation : PAN (Mw = 150,000 g mol − 1 , Sigma-Aldrich)

was mixed with 50 nm diameter crystalline Si particles (Alpha Aesar) in a 3:7 mass ratio respectively using a mortar and pestle. The mixture was then dissolved in N, N-dimethylformamide (DMF, 99%) such that 87.5 wt% of the fi nal solution/suspension was solvent. The solution/suspension was then mixed via magnetic stirring for about 12 hours to produce a viscous slurry that was bladed onto a copper foil current collector and dried at 80 ° C for about 3 hours. The subsequent electrode fi lm was then calendared before a PAN cyclization heat treatment at 300 ° C inside an Ar-fi lled glove box for about 12 hours. Electrodes were punched (½” punch) and weighed to obtain the mass of nano-Si active material. A conventional cell (depicted in Figure 3 a as nSi:AB:PVDF) was used as our control cell for a performance analysis and for comparison purposes. The same nano-Si powder used for our nano-Si/cyclized-PAN, acetylene black (AB), polyvinylidene (PVDF), and N-methyl pyrolidinone were mixed into a slurry that was bladed onto a copper foil current collector to make the electrodes for our control cell. Electrochemical measurements were all normalized based on the mass of nano-Si in each electrode (typically 0.5–0.6 mg).

Differential scanning calorimetry : DSC curves of PAN-DMF solution were obtained using a DSC S-650 by heating from 25 to 500 ° C in Argon at a heating rate of 20 ° C/min.

X-ray photoelectron spectroscopy : XPS spectra were obtained on an AXIS His 165 and ULTRA spectrometer (Kratos) to determining the

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binding confi gurations and the chemical state of the elements present in PAN fi lms heat treated under Argon at various temperatures. The nomenclature used to describe each π resonance of the N 1s spectra came from previous literature. [ 30 , 34 ]

Electrochemical Measurement : Electrochemical measurements were carried out using an Arbin TM 2000 battery test station. All cells were assembled in an Ar-fi lled glove box using our prepared nano-Si/cyclized-PAN electrodes as the working electrode and lithium metal foil as the counter electrode. The electrolyte was 1M LiPF 6 dissolved in a 1:1 (volume ratio) mixture of ethylene carbonate (EC) and diethyl carbonate (DEC), the separator was a glass micro-fi ber disk (Whatman TM GF/F) and the shell was a stainless steel CR2032 coin cell (VWR Inter.). We used a constant current constant voltage (CCCV) testing scheme to cycle our cells. The cells were discharged (lithiated) and charged (delithiated) with various cycling currents between 0.05 and 1 V (vs. Li/Li + ). The conducted rate study was carried out with charging rates ranging from C/20-5C. The discharge rates were started at C/20, increased to C/10 and maintained at this rate for subsequent cycling. Charging was conducted with constant current (CC) and discharge was conducted with CCCV cycling parameters.

Material Characterization : Raman spectroscopy measurements were performed with a Jasco NRS-3100 (532 nm excitation). A FIB (FEI, NOVA200 dual beam system) equipped with and air-lock chamber is used for TEM sample preparation. [ 10 ] TEM and EELS analysis was performed with a FEI Tecnai F20 operated at 200 keV. A detailed description of our TEM and EELS characterization procedures can be found elsewhere. [ 6 ]

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

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Acknowledgements This work was funded by the U.S. Department of Energy under subcontract number NFT-8-88527-01 through the DOE offi ce of Energy Effi ciency and Renewable Energy Offi ce of the Vehicle Technology Program, by a grant from the Fundamental R&D Program for Technology of World Premier Materials funded by the Ministry of Knowledge Economy, Republic of Korea (10037919), the National Science Foundation (NSF, DMR-1206462), and the National Science Foundation graduate research fellowship program (NSF-GRFP).

Received: October 23, 2012Published online: March 7, 2013

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