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Xin Zhao , Cary M. Hayner , Mayfair C. Kung , and Harold H. Kung *

In-Plane Vacancy-Enabled High-Power Si–Graphene Composite Electrode for Lithium-Ion Batteries

ION

The current insatiable appetite for more effi cient electrical

energy storage devices with a fl exible and/or compact confi gu-ration, fed by the proliferation of portable electronics with ever increasing functional complexity, will be even more diffi cult to satisfy when electrifi ed vehicles become the preferred mode of transportation, and energy harvesting from intermittent sources evolves into the practiced norm. [ 1–4 ] Incorporation of the attrac-tive features of supercapacitors with high rate performance and rechargeable batteries with high energy densities into a single unit would enable the design of high-capacity energy storage devices for sustainable power delivery. [ 5–7 ] Thus far, however, this goal has proved to be diffi cult to attain. Strategies used to boost the power capability of electrode materials for Li-ion bat-teries generally involve reducing the domain size of the active charge-storage material (e.g., Si) in the electrode to shorten the ion diffusion paths, such as by fabricating vertically aligned 1-D nanostructures or ultrathin coatings on nanofoams and metallic mesh. [ 8 , 9 ] Generally, such an approach suffers from lim-ited overall charge storage capacity due to a low mass fraction of the active component in the electrode. Fabricating porous electrode frameworks using sacrifi cial templates is another approach, [ 10–12 ] but these porous electrodes introduce different problems, such as confi gurational infl exibility imposed by mechanical fragility and a dramatic drop in volumetric energy density consequential of reduced packing densities. Here we demonstrate that graphene sheets possessing a high density of in-plane, -sized carbon vacancies can be transformed into a fl exible, 3-D conducting graphenic scaffold with excellent cross-plane ion diffusivity and tolerance to structural deformation. When employed as a structural platform to incorporate high storage-capacity materials, such as Si, a stable, self-supporting composite electrode with enhanced accessible interior and high rate capacity is obtained, representing an attractive electrode candidate towards high-performance Li-ion batteries. Further-more, the composite is ductile and offers confi gurational fl ex-ibility, and can be produced by cost-effective processes.

Our 3-D graphenic scaffold was constructed with aligned graphene sheets, derived from exfoliated graphene oxide sheets into which in-plane, nm-sized carbon vacancies were intro-duced by a facile wet chemical method. It confers a combina-tion of advantageous features over reported electrodes systems:

© 2011 WILEY-VCH Verlag GAdv. Energy Mater. 2011, 1, 1079–1084

Dr. X. Zhao , C. M. Hayner , Prof. M. C Kung , Prof. H. H. Kung Department of Chemical and Biological Engineering Northwestern University Evanston, Illinois, 60208, USA E-mail: hkung@northwestern.edu

DOI: 10.1002/aenm.201100426

i) Facile ion transport throughout the structure, enabled by new diffusion channels created with in-plane carbon vacancies. This overcomes the characteristic high resistance of graphene mate-rial for Li ion transport due to their extreme width-to-thickness aspect ratio and inter-sheet aggregation. [ 13–17 ] ii) Superior elec-trical conductivity and high packing density derived from the compact structure of interconnecting graphitic domains. [ 18 , 19 ] iii) Sustained structural integrity as a result of the high fl ex-ibility of the graphene sheets to accommodate large volume variations of high-capacity embedded storage materials during charge/discharge cycle. [ 17 , 20–22 ] These advantages were demon-strated with a self-supporting, fl exible electrode consisted of Si nanoparticles embedded in a 3-D graphenic scaffold, as shown schematically in Figure 1 . This Si–graphene composite elec-trode achieved an unprecedented reversible capacity of around 1100 mAh g − 1 at 8 A g − 1 , a rate equivalent to full discharge in 8 min, or approximately 3200 mAh g − 1 at 1 A g − 1 , repeatable up to 99.9% between cycles for over 150 cycles.

The graphenic scaffold was constructed by fi rst exfolia-ting graphite fl akes into graphene oxide (GO) sheets. [ 23 ] After washing, suspensions of GO were sonicated in aqueous HNO 3 solutions in a procedure similar to the ones used to generate defects on carbon nanotube sidewalls [ 24 ] and scissor GO into polyaromatic hydrocarbons, [ 25 ] except that a lower acid con-centration and a reduced sonication period rendered our method a milder version. Four different acid concentrations were investigated (see the Supporting Information for details), and the resulting samples with induced structural defects are

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Figure 1 . A schematic drawing (not to scale) of a section of a composite electrode material constructed with a graphenic scaffold with in-plane carbon vacancy defects. The graphene sheets with these holey defects are displaced from each other for clarity. Electrochemically active com-ponents, for example, Si nanoparticles (large spheres), are sandwiched between graphene sheets, and these composites are structurally inte-grated with a 3-D graphenic network of interconnecting graphitic domains formed by reconstituting these graphene sheets. Li ions (small spheres) can diffuse easily across graphene sheets throughout the structure by passing through the in-plane vacancy defects.

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Figure 2 . TEM images of Pd-stained graphene oxide samples: a) GO, b) deGO-I, c) deGO-II, d) deGO-III, and, e) deGO-IV. f) A high magnifi cation image of a highlighted region in (d), showing approximately 3 nm Pd particles in ring-like arrangements. The crystalline regions of graphene oxide are indicated by arrows.

labeled deGO-I, -II, -III, and -IV, respectively. To obtain a Si-deGO composite paper, an aqueous mixture of deGO and Si nanoparti-cles (approximately 50 nm diameter) was fi l-tered and dried.

The acid-sonication treatment created carbon vacancies with carboxylate groups decorating the defect edge sites. [ 24 ] Palladium nanoparticles, formed by reduction of Pd ions that bind to these carboxylates during examination with electron microscopy, were used as a vehicle to examine the dependence of defect generation on the severity of treat-ment. Figure 2 a–e show bright-fi eld TEM images of GO and deGOs, and Figure S1 in the Supporting Information shows the com-plementary Z -contrast images. As expected, the densities of Pd particles and larger pores on deGO increased with increasing severity of acid treatment. A high magnifi cation image of deGO-III (Figure 2 f) shows Pd clus-ters arranged in ring formations, consistent with their location at the perimeter of the window-like porous defects. Many observable defects were 10–20 nm in diameter. These observations are in agreement with X-ray photoelectron spectroscopic (XPS) measure-ments which showed that the acid-oxidation treatment increased the peak area of oxygen-containing functional groups relative to aro-matic carbon (Figure S2).

Thermal reduction of deGO to form a 3-D graphenic scaffold (deG) was accom-plished in a fl ow of either Ar or 10% H 2 in Ar at 700 ° C for 1 h with similar results. XPS confi rmed removal of most of the O in the sample (Figure S3 and S4). [ 26 , 27 ] The largest porous defects observed on thermally reduced deGO were roughly < 10 nm, 20 nm, 100 nm, and > 100 nm for samples I, II, III, and IV, respectively (Figure S5).

Both deG and Si-deG papers were highly conducting (Table S1 and S2), fl exible, and remained integral when bent ( Figure 3 a and b). In fact, the Si-deG papers were more duc-

tile than the deG papers, possibly because of the presence of roughly 30% by volume of loosely assembled Si nanoparticles in Si-deG that can be deformed readily. X-ray diffraction con-fi rmed the presence of both crystalline graphite domains and disordered regions, as indicated by a sharp peak at 2 θ = 26.5 ° on top of a broad hump (Figure S6). There was a slight shift of the hump towards lower angle as the acid concentration increased, implying reduced van der Waals attraction between the reconstituted graphene sheets with large numbers of carbon vacancies. The XRD pattern of an Si–G sample showed the crystalline graphite diffraction peak shifted to 2 θ = 26.3 ° , consistent with Si nanoparticles embedded between graphene layers (Figure S7). The Si contents in these Si–deG papers were estimated to be 65–70 wt-% (Figure S8), based on weight losses

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upon graphene combustion in thermogravimetric analysis. SEM images revealed homogeneously dispersed Si nanoparticles sandwiched between reassembled graphene sheets, along with visible porosity on the basal planes of deG (Figure 3 c and d). Synchrotron nonresonant X-ray Raman scattering (XRS) of Si–deG fi lms showed the absence of silicon carbide phase.

The Raman spectra of the reduced Si–deG and Si–G samples showed little change in the intensities of the D bands (approxi-mately 1350 cm − 1 ) relative to the G bands (approximately 1580 cm − 1 ) (Figure S9). The intensity ratios of the D/G bands corresponded to an approximately 4 nm sized ordered graph-itic domain calculated using the empirical Tuinstra-Koenig rela-tion. [ 28 ] This is consistent with the fact that samples containing in-plane defects remained intact, mechanically strong, and

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Figure 3 . Microstructures of Si–graphene composite papers. a) Digital image of a Si–deGO-III paper. b) Digital image a Si–deG-III paper. c) SEM image of the top surface of a Si–deG-III paper. d) SEM image of the cross-section of a Si–deG-III paper, the inset shows Si nanoparticles embedded between graphene sheets uniformly. e) SEM image of the top surface of Si–deG-III paper at higher magnifi cation, with circles highlighting the in-plane defects. f) HRTEM image of crushed Si–deG-III showing the distribution of approximately 50 nm Si nanoparticles in the graphene sheets, the inset shows the circled region at higher magnifi cation. The Si nanopar-ticles have an amorphous SiO x surface layer with a thickness of 2–3 nm outside the metallic crystalline core. A reconstituted graphitic phase composed of graphene is also identifi ed.

electrically conducting, and the TEM images showing defects dispersed in the graphene planes.

Electrochemical properties were characterized using a half-cell arrangement with a Li foil as the counter electrode. After 5 charge/discharge cycles at 1 A g − 1 between 0.02–1.5 V (vs. Li/Li + ), an approximately 5 μ m thick Ar-reduced Si–deG-II paper exhibited a reversible capacity close to its theoretical value of approximately 3200 mAh g − 1 (see the Experimental Section). Thereafter, the capacity loss was only 0.14% per cycle, such

© 2011 WILEY-VCH Verlag GmbH & Co. KGaA, WeinhAdv. Energy Mater. 2011, 1, 1079–1084

that 83% of its theoretical capacity was main-tained after 150 cycles ( Figure 4 a). The initial irreversible capacity losses were mostly due to the formation of solid electrolyte interphase (SEI) [ 29 ] and reactions of Li with residual O and H-containing groups in graphene. [ 30 ] When tested at 8 A g − 1 , a reversible capacity of approximately 1100 mAh g − 1 was obtained, which degraded at approximately 0.34% per cycle for 150 cycles. This specifi c current cor-responded to a rate of approximately 8 C based on the accessible storage capacity, or approxi-mately 2.6 C based on the theoretical capacity (theoretical C -rate). The characteristic voltage plateaus for phase transformation of amor-phous Si at 0.3–0.17 V and 0.1–0.05 V (lithia-tion), and 0.2–0.4 V and 0.45–0.58 V (del-ithiation) [ 31 ] remained distinguishable even at such high rates (Figure 4 b). This illustrates that our Si–deG composites can deliver an exceptionally high energy output even at a rate one order of magnitude higher than for other reported Si–carbon composites. [ 17 , 32 , 33 ]

Capacity fade of the composite electrodes could be reduced by cycling through a nar-rower range of 0.1–0.55 V to minimize the destructive effect of volume variation and possible dissolution and re-formation of SEI at high voltages. At 4 A g − 1 (around 1.3 C theoretical), the capacity fade was as low as approximately 0.1% per cycle, and a revers-ible capacity of approximately 600 mAh g − 1 was maintained after 150 cycles. The high cycling stability and rate capability of Si–deG-II is demonstrated in Figure 4 c. Reversibility of storage capacity was observed when the rate was fi rst stepwise increased from 0.1 to 8 A g − 1 and then switched back.

The ability to maintain a high capacity at high rates for the Si–deG samples was attrib-uted to the much shortened Li diffusion path throughout the electrode, leading to a fully accessible interior and fast lithiation and del-ithiation reactions of Si nanoparticles. This was confi rmed by electrochemical imped-ance spectroscopy. Nyquist plots (Figure S10) of Si–deG showed a pronounced reduc-tion in both the charge-transfer resistance and the Warburg coeffi cient [ 34 ] compared with Si–G, whereas the constant phase ele-

ment increased gradually, implying a more capacitive interface (Figure S11, Table S3 and S4).

The defect size and density strongly infl uenced the electro-chemical performance of Si–deG composites, while reduction of Si–deGO in different atmospheres showed minor effects (Figure 4 and S12). Without the intentional defects, the capacity of Si–G diminished by around 90% when the rate was increased from C /30 to 2.6 C (theoretical). Si–deG-II was the most tolerant to high rates, maintaining 34% of the theoretical capacity at

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Figure 4 . Electrochemical characterization of approximately 5 μ m thick Si–deG and deG paper anodes. a) Specifi c delithiation capacity (solid) and cou-lombic effi ciency (open) of Ar-reduced Si–deG-II between 0.02–1.5 V at 1 and 8 A g − 1 ( C /3 and 2.6 C based on a theoretical capacity of 3052 mAh g − 1 ), and between 0.1–0.55 V at 4 A g − 1 (1.3 C ). b) Fifth-cycle charge/discharge curves of Ar-reduced Si–deG-II at 1, 4, and 8 A g − 1 . c) Specifi c delithiation capacity (solid) and coulombic effi ciency (open) of Ar-reduced Si–deG-II between 0.02–2.0 V at current densities of 0.1, 0.2, 0.4, 1, 2, 4, and 8 A g − 1 ( C /30, C /15, C /7.6, C /3, C /1.5, 1.3 C , and 2.6 C ) and recovery of the original capacity when switched back to 0.1 A g − 1 . d) Ar-reduced and H 2 /Ar-reduced Si–deG between 0.02–1.5 V at current densities of 0.1, 0.2, 0.4, 1, 2, 4, and 8 A g − 1 . e) Specifi c delithiation capacity (solid) and coulombic effi ciency (open) of Ar-reduced and H 2 /Ar-reduced deG-II between 0.02–2.0 V at current densities of 0.05, 0.1, 0.2, 0.5, 1, 1.5, and 2 A g − 1 ( C /7.4, C /3.7, C /1.8, 1.3 C , 2.7 C , 4 C , and 5.4 C based on a theoretical capacity of 372 mAh g − 1 of graphenic materials). Capacity is calculated using the total mass of each electrode.

8 A g − 1 . This suggests that there is an optimal pore size and/or density of defects, which might be attributed to a combination of factors including diffusivity of Li ions, ability of the porous graphene sheets to retain Si nanoparticles and accommodate their volume changes, as well as formation and nature of SEI at the defect sites. SEM images showed that at the optimal con-ditions, the Si–deG papers remained integral after rate tests (Figure S13).

While most of the data were collected using 5 μ m thick elec-trodes, similar benefi cial effects of the defects were observed

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with 20 μ m thick electrodes (Figure S12). That the benefi cial effects were due to in-plane defects and enhanced ion diffu-sivity throughout the composite paper was further supported by the surface area and pore size distribution information, shown in Figure S14. A sample with defects showed a higher surface area and, most importantly, much larger mesopore volume, consistent with easier access to more internal vol-umes of the composite paper. The benefi cial effect of in-plane defects was apparent even for the graphenic scaffold without Si. With induced defects, a capacity as high as 180 mAh g − 1 was

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obtained at a current of 2 A g − 1 for deG-III, compared to only 70 mAh g − 1 without defects (Figure 4 e). The coulombic effi cien-cies were very close to 100% with no obvious capacity fade over the 100 cycles tested.

Whereas the effect of atomic defects on the electronic and magnetic properties of graphene has captured much atten-tion, [ 35 , 36 ] we have demonstrated the benefi cial effects of nanometer- sized in-plane vacancies on ion transport for use in a 3-D graphenic scaffold that can be fabricated into hybrid materials with a hitherto unachievable combination of power capability and storage capacity for battery electrode applica-tions, without sacrifi cing its mechanical properties. Very signif-icantly, the preparation method employed can be easily adapted and scaled up to existing high-throughput processing protocols, such as spray-coating, inkjet printing, and roll-to-roll deposi-tion, for electrode manufacture. These high-performance, free-standing graphenic composite sheets with engineered porosity for cross-plane transport may also fi nd broad applications as support scaffolds and membranes for water desalination and remediation as well as in curvilinear electronics, catalytic, bio-fuel and biomimetric systems with new functionalities. [ 37 , 38 ]

Experimental Section Material Synthesis : GO was synthesized from fl ake graphite by

a modifi ed Hummers method. [ 23 ] Introduction of in-plane carbon defects was achieved by mixing approximately 0.1% w/w aqueous GO suspension with different amounts of 70% concentrated HNO 3 (GO suspension/70% HNO 3 volume ratio of 1:5 (sample I), 1:7.5 (II), 1:10 (III) and 1:12.5 (IV)), and sonicating in a bath sonicator for 1 h. The resulting deGO was fi ltered and washed using an Anodisc membrane fi lter, followed by air drying. The GO or deGO papers were reduced with either a fl ow of Ar (approximately 90 mL min − 1 ) or 10% hydrogen in Argon (approximately 100 mL min − 1 total fl ow) at 700 ° C for 1 h to form G or deG.

To form Si–graphene paper composites, an aqueous, homogeneous suspension of air-exposed Si nanoparticles and GO or deGO of the appropriate weight ratios was vacuum-fi ltered, dried, and thermally reduced as described above.

Structural Characterization : The as-prepared (de)GO and (de)G samples were characterized with fi eld emission scanning electron microscopy (FE-SEM), fi eld emission transmission electron microscopy (FE-TEM), X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), synchrotron nonresonant X-ray Raman scattering (XRS), Raman spectroscopy, and surface area and pore size distribution measurements. The Pd-stained (de)GO was prepared by mixing (de)GO with a dilute aqueous Pd(OAc) 2 solution. Thermogravimetric analysis (TGA) was performed on crushed Si–(de)G composite papers in air. Sheet resistance was measured with a four-point probe technique with an electrode separation of 1 mm using a Keithley 2400 Sourcemeter.

Electrochemical Tests : Half-cell tests were conducted using two-electrode coin cells with Li metal as the counter electrode, assembled in an argon-fi lled glove box, at various current densities, typically in the voltage range of 0.02–1.5 V vs Li/Li + . Electrochemical measurements of Si–(de)G composite papers were carried out using three-electrode Swagelok-type cells, with an independent Li metal reference electrode. Electrochemical cycling of Si–(de)G composite electrodes was performed using two different procedures, either at various current densities in the voltage range of 0.02–2 V vs Li/Li + or a constant current–constant voltage (CCCV) method. Electrochemical impedance spectroscopy (EIS) measurements were conducted on two-electrode Swagelok-type cells by applying an AC voltage of 10 mV amplitude and DC open circuit voltage (OCV) in the frequency range of 1 MHz–0.01 Hz at room temperature.

© 2011 WILEY-VCH Verlag GmbAdv. Energy Mater. 2011, 1, 1079–1084

Theoretical capacity of Si–G composite is calculated as: 4200 mAh g − 1 (Si theoretical capacity) × 70% (Si wt-%) + 372 mAh g − 1 (graphene theoretical capacity) × 30% (graphene wt-%) = 3052 mAh g − 1 .

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

Acknowledgements This research was supported by the U.S. Department of Energy, Basic Energy Sciences, grant DE-AC02-06CH11357 through the Center for Electrical Energy Storage, an Energy Frontier Research Center. We thank Dr. Shuyou Li (EPIC, NU), Dr. Xinqi Chen (KECK, NU), and Mr. Yuki Kusachi (Nissan) for helpful discussions.

Received: July 28, 2011 Revised: August 23, 2011

Published online: October 6, 2011

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