Materials Engineering of High-Performance Anodes as Layered

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Materials Engineering of High-Performance Anodes as Layered Nanoparticle Composites with Self-Organizing Conductive Networks Lehao Liu, Bong Gill Choi, Siu On Tung, Jing Lyu, Tiehu Li, Tingkai Zhao, and Nicholas A. Kotov J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b01105 • Publication Date (Web): 08 May 2018 Downloaded from http://pubs.acs.org on May 18, 2018

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The Journal of Physical Chemistry

Materials Engineering of High-Performance Anodes as Layered Composites with Self-Assembled Conductive Networks Lehao Liu†,‡,‫׏‬, Bong Gill Choi†,‖, Siu On Tung§, Jing Lyu†,‫׏‬, Tiehu Li‫׏‬, Tingkai Zhao‫׏‬, and Nicholas A. Kotov†,┴,* †Department

of Chemical Engineering, University of Michigan, Ann Arbor, Michigan 48109,

USA; ‡State Key Laboratory of Alternate Electrical Power System with Renewable Energy ‫׏‬ Sources, North China Electric Power University, Beijing 102206, PR China; School of Materials Science and Engineering, Northwestern Polytechnical University, Xi'an 710072, Shaanxi, PR China; ‖Department of Chemical Engineering, Kangwon National University, Samcheok 245-711, Republic of Korea; §Macromolecular Science and Engineering, University of Michigan, Ann Arbor, Michigan 48109, USA; ┴Biointerfaces Institute, University of Michigan, Ann Arbor, Michigan 48109, USA. *To whom correspondence should be addressed: [email protected].

ABSTRACT: The practical implementation of nanomaterials in high capacity batteries has been hindered by the large mechanical stresses during ion insertion/extraction processes that lead to the loss of physical integrity of the active layers. The challenge of combining the high ion storage capacity with resilience to deformations and efficient charge transport is common for nearly all battery technologies. Layer-by-layer (LBL) assembled nanocomposites were able to mitigate structural design challenges for materials requiring the combination of contrarian properties in the past. Herein, we show that materials engineering capabilities of LBL augmented by self-organization of nanoparticles (NPs) can be exploited for constructing multiscale composites for high capacity lithium ion anodes that mitigate the contrarian nature of three central parameters most relevant for advanced batteries: large intercalation capacity, high conductance, and robust mechanics. The LBL multilayers were made from three function-determining components, namely polyurethane (PU), copper nanoscale particles, and silicon mesoscale particles responsible for the high nanoscale toughness, efficient electron transport, and high lithium storage capacity, respectively. The nanocomposite anodes optimized in respect to the layer sequence and composition exhibited capacities as high as 1284 and 687 mAh/g at the 1st and 300th cycle, respectively, with a fading rate of 0.15% per cycle. An average Coulombic efficiencies were as high as 99.0—99.4% for 300 cycles at 1.0 C rate (4000 mA/g). Self-organization of copper NPs into three-dimensional (3D) networks with lattice-to-lattice connectivity taking place during LBL assembly enabled high electron transport efficiency responsible for high battery performance of these Si-based anodes. This study paves the way to finding a method for resolution of the general property conflict for materials utilized in for energy technologies. 1

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KEYWORDS: layer-by-layer assembly, self-assembly, copper nanoparticles, property correlations, silicon particles, contrarian properties, polyurethane, flexible nanocomposite films, anode materials, batteries.

1. INTRODUCTION High theoretical capacity of silicon (Si) for lithium storage makes it a promising anode material for lithium-ion batteries (LIBs).1–3 However, the practical implementation of Si anodes is hindered—even after many years of research--by poor cycle performance caused by large changes in volume of up to 300—400% in response to intercalation of Li+ ions into crystal lattice of Si.4–7

Cyclic insertion of

large volumes of lithium inevitably generates large mechanical stress and causes the loss of physical integrity of anodes.8–10 Considerable strides were made toward resolving this problem.2,4,5,9,11–14 However, simultaneous improvements in capacity, cycle life, and charge-discharge rate needed for advanced batteries with Si anodes require properties that seem to be antagonistic to each and this fundamental problem represents the key technological bottleneck for the energy storage technologies. Besides Si, the same problem of cyclic expansion-contraction cycles due to ion intercalation resulting in physical disintegration and delamination of active layers is observed many other anodes and cathodes. In more general terms, this materials engineering problem can be described as a property conflict between desirable functionality (i.e. capacity), transport properties, and nanoscale mechanics. Its commonality can also be appreciated from the similar materials design challenges identified for transparent conductors,15–18 electrolytes19–23 as well as other energy conversion materials including solar cells.24–26 2


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Several strategies were used in the specific case of Si anodes to mitigate the loss of physical integrity, electron transport, and associated with it capacity fade of batteries.5,9,13

Reduction of the dimensions of Si phase helps in resolving the

transport vs mechanics conflict because nanoscale versions of Si can relax stress during the volume expansion better than macroscale particles. Many Si nanostructures exemplified by anodes based on Si nanoparticles (NPs),27,28 nanowires (NWs),1,29 and nanotubes,28,30,31 have shown enhanced cycle lives.

However, coatings with high

volume fraction of Si nanostructures becomes poor electron conductors and their actual ability to store lithium ions is greatly reduced compared to their theoretical capacity due to emergence of electron transport problems caused by the tunneling gaps between neighboring NP or NWs.32

Furthermore, slurries of Si nanostructures

are difficult to deposit without phase segregation whereas in-situ growth of Si NWs presents integration challenges with the other battery components.27,29,33 Nanocarbons interdigitated with Si nanostructures can serve as conduits of electron transport and as a deformation resilient mechanical framework.

Si anodes

incorporating layer of nanoscale carbon layers on Si improved the cycle performance of LIBs.34–39

However, such materials design approach evokes new challenges

related to thermal and chemical compatibility of different battery components. The carbon layers tightly bound to Si surfaces were obtained by pyrolysis of polymers 7,39 at high temperatures or etching of SiO2 layer from the surface of Si by hydrofluoric acid.40,41 Three-dimensional (3D) porosity that helps to accommodate the structural 3



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changes involved in expansion/contraction cycles was also used in materials design of anodes

6,28,31,42–46

and cathodes.47–54

Nanoscale copper (Cu) or carbon are typically

used as electron transport components of the composites typically in the form of continuous coatings.55

Despite the complexity of manufacturing of 3D electrodes

and their integration with other battery components,5,42,44,56, these studies demonstrated a purposeful and viable fundamental materials engineering approach to resolution of the conflict between high capacity and mechanical robustness of anode materials. Layer-by-layer assembly (LBL) is a materials engineering technique based on sequential adsorption of high molecular weight components applicable both to solely organic,57–60 inorganic,61–63 and hybrid materials.16,64–67

Composite materials made

by LBL are known for their (1) wide range of building blocks that produce LBL films;16,57,62,68–70 (2) nanometer-level control of the materials architecture,71–75, (3) nearly perfect conformal deposition on surfaces of complex geometries.58,76–81, (4) effective suppression of spontaneous phase separations making possible materials that could be far away from their structural equilibrium82,83, and (5) the ability to combine self-assembly processes in dispersions with those on interfaces.16,67,84,85 As a consequence of these unusual features, the resulting materials display unusual combination of mutually restrictive materials properties such as high Young’s modulus and density86–90 or high electronic conductance and stretchability.48,82,91,92 Thus LBL-made materials make it possible to successfully obtain materials with functional, mechanical, and transport characteristics16,77,89,91,93–96 that might appear to 4


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be fundamentally impossible.97 Materials design based on the alternating layer motif inspired by biological nanocomposites, such as nacre,66,91,98–107 have been shown to improve performance of different battery components.73–75

The multilayer stacks were obtained by

magnetron sputtering, thermal evaporation, and ultrafiltration from the typical battery materials, such as coper,72 graphene,73 tin,73 manganese dioxide,74 and titanium75. In these nanocomposites, active components represented by Sn or Si offered high theoretical

capacity,

while

other

components

provided

electron

transport

characteristics. However, the rigidity of the stacks inevitably reduces its ability to accommodate large mechanical stress due to large volume expansion of Si or Sn components. Based on these body of knowledge, we hypothesized that one can utilize LBL as a general method for materials engineering for energy materials creating a composite multilayers from three inexpensive components so that each of them would be primarily responsible for a specific characteristic in the resulting nanocomposite. Early

works

on

LBL

composites

for

battery

materials

support

this

hypothesis.20,95,108–111 Thus, we decided to carry out systematic study focused on LBL-engineered anodes for LIBs using the technological bottleneck with Si anode as example.

As multilayer components we utilized water-soluble polyurethane (PU),

copper NPs, and Si mesoscale particles (MPs); the assembled materials inherited a measure of properties typical for three of these components: toughness of PU, efficient electron transport of Cu NPs, and high capacity of Si MPs. Importantly, 5



spontaneous formation of Cu NP into chains with lattice-to-lattice connectivity was observed during LBL assembly which greatly improved charge transport in the material.

We analyzed the performance of LIB cells with anodes made with

different sequences and composition demonstrating the protocol of structural optimization for the nanocomposites resulting in competitive charge storage performance. These findings demonstrate that LBL assembly can be a general materials engineering and manufacturing tool112–116 for a variety of materials for energy conversion and storage.117 2. METHODS 2.1. Materials and chemicals. Copper (II) sulfate, mercaptosuccinic acid, anhydrous hydrazine, Trizma base, hydrochloride acid (HCl), hydrofluoric acid (HF), sulfuric acid, hydrogen peroxide, and sodium hydroxide were purchased from Sigma-Aldrich. Si powder was purchased from Alfa Aesar. Positively-charged PU (30 wt%, molecular mass: 92,000 D) was obtained from HEPCE Chem, South Korea. Deionized (DI) water was obtained from a Barnstead E-pure water purification system. The coin cell case, separator, carbon black, polymer binder and electrolyte for the Li-ion battery assembly were purchased from MTI Corporation. 2.2. Synthesis of Cu NP and Si particle aqueous dispersions. (1). To prepare the Cu NP dispersion, 0.04 g copper sulfate and 0.08 g mercaptosuccinic acid were dissolved into 45 mL DI water and then heated at 60 oC under strong magnetic stirring for 30 min. Next, 0.6 mL 10 M sodium hydroxide solution was added under strong stirring for 30 min, after which 5 mL hydrazine was added into the mixture under strong 6


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stirring for an additional 30 min (sediments appeared after 40 min). The reaction solution was subsequently centrifuged at 5,000 rpm for 30 min, after which it was dispersed with DI water under ultrasonic treatment for 1 h to obtain a highly-concentrated Cu NP solution for LBL assembly. (2). To prepare the Si particle dispersion, 100 mg Si particles were sonicated for 12 h in 100 mL 0.01 M pH 8.5 Tris-HCl buffer solution, centrifuged at 8,000 rpm for 30 min, and then dispersed with DI water under ultrasonic treatment for 1 h to yield a highly-concentrated Si particle dispersion for LBL assembly. 2.3. LBL assembly of Cu, Si, and PU for anode composite films. Glass slides (25 mm × 75 mm, Fisher Scientific) were cleaned with piranha solution (sulfuric acid and hydrogen peroxide in a 3:1 volume ratio) overnight. Warning: piranha solution is dangerous and extremely reactive with organic substances so appropriate handling precautions are essential. (1). To prepare PU/Cu and PU/Si LBL films, the pretreated glass slides were immersed in 0.5 wt% PU solution for 5 min, rinsed with DI water for 1 min, and then gently dried with compressed air. Next, the slides were immersed in 1 mg/mL Cu NP solution (or 4 mg/mL Si particle solution) for 30 min, rinsed with DI water for 1 min, and again dried with compressed air. The films made by the entire alternating adsorption processes of PU and Cu (or Si) are denoted as (PU/Cu)N (or (PU/Si)N), where N is the number of pairs of LBL deposition cycles. (2). To prepare PU/Cu/Si LBL films, the pretreated glass slides were immersed in 0.5 wt% PU solution for 5 min, rinsed with DI water for 1 min, and then gently dried with compressed air. The slides were subsequently immersed into 4 mg/mL Si particle 7



solution for 30 min, rinsed with DI water, and again dried with compressed air. This deposition cycle constructed one bilayer of (PU/Si). The resultant films were then dipped into 0.5 wt% PU solution for 5 min, after which they were washed with DI water and dried with compressed air. The samples were subsequently immersed into 1 mg/mL Cu NP solution for 30 min, rinsed with DI water, and dried with compressed air. This cycle was repeated m times, making m bilayers of PU and Cu, denoted (PU/Cu)m. The films made by the entire alternating adsorption process of PU, Si, and Cu are denoted here as [(PU/Si)(PU/Cu)m]N, where the total bilayer number was (1+m)×N. After several tens of deposition cycles, the Cu and Si particle solutions needed to be changed to avoid aggregation. Free-standing films were obtained by peeling the films off the glass slides using 0.5 vol% hydrofluoric acid solution. Warning: hydrofluoric acid solution, even diluted, is very toxic so extreme precautions must be taken at all times. After thorough rinsing with DI water, the detached films were dried under vacuum at 50 oC for 1 day prior to measurements. The purchased Cu foils were also cleaned with the piranha solution for 30 sec and then with ethanol for the LBL assembly. After the piranha treatment, not only the color of the Cu foils changed obviously (Fig. S1), but also the resistivity decreased from 4.4×10–8 to 3.3×10–8 Ωꞏm (by RTS-8 4-point probes resistivity measurement system at room temperature of ~30 oC), which is close to the resistivity of the pure Cu (1.7×10–8 Ωꞏm at 20 oC). Correspondingly, the electrical conductivity of the pretreated Cu foils increased from 2.3×10–7 to 3.0×10–7 Sꞏm–1. 2.4. Assembly of lithium-ion coin cell battery. A two-electrode, CR2032 coin cell 8


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was constructed with a Celgard LIB polyethylene separator from MTI Corporation. A Li foil and a LBL film (both 1 cm × 1 cm) were used as counter/reference electrode and working electrode, respectively. 1 M lithium hexafluorophosphate (LiPF6) in a solution of 1:1 volume ratio of ethylene carbonate (EC) and dimethyl carbonate (DMC) was used as electrolyte. The coin cell was assembled in a glove box filled with argon (oxygen content:  3 ppm; moisture content:  3 ppm). We also prepared Si particle slurry coated electrodes with the Si particles, graphite, and polyvinylidene fluoride (PVDF).

2.5. Structural characterization. Transmission electron microscope (TEM) images were obtained using a JEOL 3011. The TEM specimens of Cu NPs were prepared by dripping a single drop of Cu NP solution onto a carbon-coated Cu grid and then allowing the drop to dry in air at ambient temperature. The average diameter of the Cu NPs was calculated by counting Cu NPs from TEM images. The zeta potential and average size of the Cu NPs and Si particles were determined using a Zetasizer instrument (Malvern Instruments, U.K.). Scanning electron microscope (SEM) images were obtained using a FEI Nova 200 SEM. To take the SEM images, the PU/Si LBL film was sputtered with gold at a current of ～20 mA for 30 seconds. However, no sputtering was applied to the PU/Cu and PU/Cu/Si LBL films. UV-Vis absorbance spectra were obtained using an 8453 UV-Vis Chem Station spectrophotometer (Agilent Technologies, USA). The chemical compositions of the LBL films were determined using a PHI 680 Scanning Auger Nanoprobe (Physical Electronics Inc., USA). Prior to data acquisition, the samples were cleaned by Ar+ ion sputtering for 20 9



min (1 nm/min for a 1×1 cm area) to remove any surface contamination. The incident beam current and beam voltage were 10 nA and 10 kV, respectively. Electrical resistance measurements with four-point methods were obtained using a 34401A Digital Multimeter (Agilent Technologies, USA). Stress-strain curves were obtained by testing rectangular strips (3 mm width and 15—25 mm length) of the films using a 100Q1000 mechanical strength tester (TestResources, USA) at a rate of 0.1 mm/s. The force measurement from the load cell was divided by the measured thickness and initial width of the samples to give the nominal stress.

2.6. Electrochemical characterization. The charge-discharge cycle was conducted using a Maccor Series 4000 48-channel battery tester (Maccor, USA) in a voltage range of 0.01—2.0 V (vs. Li+/Li) at various rates between 0.1—2 C (1 C = 4000 mA/g). The specific capacities were calculated based the content of Si in the composite films. The Si content of the [(PU/Si)(PU/Cu)2], [(PU/Si)(PU/Cu)3], [(PU/Si)(PU/Cu)4] films were determined to be 3.13, 2.70, and 2.15 wt%, respectively, using a software—MultiPak V9.4.1 from Physical Electronics, Inc. Electrochemical impedance spectroscopy (EIS) measurements were carried out using a Li foil as counter/reference electrode, and a LBL film as working electrode. Before the EIS measurements, all the electrodes were discharged to 0.01 V and then left on open-circuit condition for one day to obtain equilibrium. The impedance spectra were obtained using a CHI660D electrochemical workstation (CH Instrument Inc., USA) by applying a sine wave with an amplitude of 5 mV over a frequency range from 100 kHz to 0.01 Hz. The impedance spectra were fitted to a proposed equivalent circuit 10


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using a code ZSimpWin (EChem Software Company, USA).

3. RESULTS AND DISCUSSION 3.1. Dispersions of Cu NPs and Si MPs.

To obtain PU/Cu/Si composite films, we

first prepared dispersions of Cu NPs. We focused on aqueous dispersions despite the chemical challenges associated with the propensity of metallic Cu to oxidize in presence of water and oxygen118–120 because water is an intrinsically better solvent than organic ones both from environmental and technological perspectives. Aqueous media also promotes self-assembly of many NPs in superstructures with higher degree of organization, which is essential for gain the simultaneous improvements in both mechanical and transport properties. Chemical reduction of copper salts by hydrazine in presence of mercaptosuccinic acid as surface ligand (see Methods) made possible preparation of Cu nanocolloids as aqueous dispersions.

The obtained Cu NPs were uniform with an average diameter

of 6.90 ± 1.02 nm (Figure 1A and E). High resolution TEM revealed the average lattice spacing of the inorganic core of Cu NPs to be 0.21 nm (Figure 1B), which corresponded well to the spacing of (111) lattice planes in a face-centered cubic (FCC) Cu. After re-dispersion in deionized (DI) water, the average diameter of the Cu NPs increased to 9.45 ± 1.45 nm (Figure 1C and F). The oxidation state of Cu was determined by high-resolution TEM (Figure 1D). Similarly to previously studied CdTe121 and other NPs,122–125 the Cu NPs self-assembles into closely connected to each other, assembling a necklace-like chains 11



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after re-dispersion (Figure 1C-D). SEM images confirmed that the dried NPs exhibited a porous structure (Figure S2A), made from interconnected NP chains (Figure

S2B).

The

distinct

NP-to-NP

bridging

with

lattice-to-lattice

connectivity121,124,126–131 enables electrical conductivity of the NP-assembled films.32,84

We use here the term “lattice-to-lattice” connectivity instead of “oriented

attachment” because the presence of epitaxial match at the NP-NP junctions is uncertain in this system.

Furthermore, specific orientation of the crystal lattices

between the nanoscale grains is not required for metallic conductance in the NP assemblies.84 A parallel can be drawn between the self-assembly of Cu NPs and self-assembly of graphene sheets in other composite materials.23,66,132–135 Both of them result in spontaneous formation of 3D conductive networks in the composites, which is essential for scalability of the process. UV-Vis absorbance spectra were used to characterize the original and re-dispersed NPs (Figure S2C). The peaks between 500—600 nm were assigned to the surface plasmon resonance absorption of Cu NPs, and broadened wavelengths with red shifts indicated the presence of Cu oxides.136–139

At the same concentration

of 0.30 mg/mL, the original NP solution showed a narrow peak at 510 nm, while the re-dispersed NP solution showed a peak at 550 nm with broadening and tailing toward longer wavelengths. These changes indicated that the re-dispersed NPs are Cu NPs with some presence of Cu oxides, which can be further confirmed by EDS (~3.90wt% oxygen, Figure S2D) and observation of the colors of the colloids. At equal concentrations of 0.30 mg/mL, the original NP solution was red, while the 12


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re-dispersed NP dispersion was dark-red (Figure 1G) indicating the formation of dynamic agglomerates with higher light absorption cross-section than individual NPs. When the concentration of the re-dispersed NP dispersion was decreased to 0.03 mg/mL the color of the dispersion changed to yellow. For making LBL composite, we use commercial Si powders that were dispersed by 12 hour sonication of a 1.0 mg/mL slurry in aqueous media buffered at pH 8.5 (Figure S3A). Dynamic light scattering (DLS) revealed an average Si particle size of 270 nm at a concentration of 0.25 mg/mL (Inset of Figure S3A). Electron micrographs showed that the dispersions are comprised of spherical mesoscale particles (MPs) of Si (122.97 ± 70.18 nm in diameter) with wide size distribution from approximately 25 to 400 nm in diameter (Figure S3B-C). The lattice space was about 0.31 nm, which corresponded to (111) lattice planes of Si (Figure S3D).

3.2. LBL assembly of PU/Cu/Si nanocomposites. After preparing the Cu NP and Si MP dispersions, we obtained PU/Cu/Si nanocomposite films by LBL assembly to obtain composites that should hypothetically combine the stretchability of PU, conductivity of Cu, and high lithium storage capacity of Si (Scheme 1). The positively charged PU macromolecules were sequentially deposited with negatively charged Si MPs and Cu NPs (Figure S4), producing multilayers as (PU/Cu)N, (PU/Si)N, and [(PU/Si)(PU/Cu)m]N where N and m are the numbers of the corresponding multilayer sequence.

(PU/Cu)N multilayers served for imparting electron transport to the 13



[(PU/Si)(PU/Cu)m]N composites. The LBL assembly process for the preparation of (PU/Cu)N films was monitored by UV-Vis absorbance spectroscopy (Figure 2A). Optical photographs also verified the rapid formation of (PU/Cu)N LBL films on glass slides (Figure 2B) resulting in semitransparent coatings for yellow-colored (PU/Cu)20 and (PU/Cu)40 bilayer film, but becoming expectedly non-transparent for (PU/Cu)60. Free-standing (PU/Cu)N composites were obtained by etching with 0.5 vol.% hydrofluoric acid (HF) solution. SEM images of (PU/Cu)N were taken without gold or carbon sputtering because of their high conductivity (Figures 2C-H). (PU/Cu)110 displayed a thickness of ～4.3 µm (Figure 2G), with a coarser top than bottom surface (Figure 2C and E). Higher magnification images of both the top and bottom surfaces showed that interconnected Cu NPs chains formed during the LBL deposition from linear particle agglomerates in dispersion121,140–142143 and those formed on the substrate surface84,144,145 lead to spontaneous development of 3D conductive network in this composite (Figure 2D and F, Figure 1C-D and S2A-B). The SEM images of the cross section also showed exposed NPs closely connected to each other (Figure 2H). The NP attaches to each other during LBL deposition,124,146 which resulted in highly conducive material similarly to the previous cases of Au and PbS NPs.84,124,146 The growth of (PU/Si)N LBL films was also monitored by UV-Vis absorbance spectroscopy (Figure 3A). No obvious absorbance peaks were observed in the spectra of PU/Si composites but similar to the PU/Cu film, absorbance increased as N increased for the (PU/Si)N LBL films resulting in yellow-brown (PU/Si)50 composite (Figure 3B). The top surface of the (PU/Si)50 was coarse (Figure 3C), which was due 14


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to the wide size distribution of the Si MPs. A higher magnification SEM image revealed that the Si MPs were tightly packed with interstitial spaces completely filled with PU (Figure 3D) as a result of the strong attraction between Si MPs and PU. The thickness of the (PU/Si)50 films was ～1.0 µm (Figure 3E), which was suitable for battery anodes. Because of the intrinsically low electrical conductivity of Si, we combined the (PU/Si) and (PU/Cu) deposition cycles to obtain [(PU/Si)(PU/Cu)m]N with combined the functionalities essential for LIBs: mechanical toughness, high electron transport, and high Li+ capacity (Scheme 1).

Photographs revealed that the color of the films

became darker when m increased from 1 to 4 (Figure 4A), a result of the increased Cu content. Because of the flexibility and stretchability of PU,84,147 these composites can easily be peeled off from the glass slides, as well as folded and unfolded repeatedly without breaking (Figure 4B). Metal luster was observed on the surface of these films, an indication of electrical conductivity. The top surface of the film was coarser than the back (Figure 4C-F and S5-7). Some pores were formed because of the particle size differences of the Cu NPs and Si MPs. Although these pores may affect the mechanical strength of the composite films, they can effectively sustain the changes in volume of the Si particles during Li insertion/extraction processes. Higher magnification images showed that the Si particles were encapsulated with the Cu NP formed ligaments, which should be highly conducive to the metal conductivity of the composite films. [(PU/Si)(PU/Cu)]100 film showed a thickness of ～7.0 µm. In the earlier experiment period, we also tried to prepare free-standing films by simple mechanical mixing of the three components (PU, Cu NP and Si 15



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MP)-involved solutions and subsequent vacuum filtering, but the obtained composite films were easy to crack into smaller pieces (Fig. S8A-B). This phenomenon was different from our previous report,84 where we prepared free-standing PU/Au composite films with high stretchability and conductivity by the solution mixing and filtering. Therefore, we thought that the agglomeration of the Cu and Si particles in the mixture solution had detrimental impact on the formation of the PU/Cu/Si films. As we have summarized in the introduction section, the LBL assembly technology possesses the merits such as nanometer-level control of the materials architecture and effective suppression of spontaneous phase separations, and thus the free-standing PU/Cu/Si composite films

can

easily

form

by

the

electrostatic

attraction

between

the

positively-charged PU and negatively-charged Cu and Si particles. TEM was also used to investigate the microstructures of the PU/Cu and PU/Cu/Si LBL nanocomposites (Figure 5). The scratched (PU/Cu)N films from glass slides after ultrasonication maintained film structure composed of the interconnected Cu NPs (Figure 5A-B), which was consistent with the date presented in Figure 2C-H. Although the sonication partly destroyed the structure of the PU/Cu/Si composites, all the observed Si MPs are connected Cu NPs (Figure 5C-D, Figure 4C-H). Such film architecture render this composite material mechanically robust and electrically conductive via formation of 3D pathways with efficient charge transport. To prepare anodes for LIBs, we assembled the PU/Cu/Si composites on the Cu foils pretreated with piranha solution (Figure S9-12). Films of PU/Cu/Si composites grown on glass and copper showed the same color and microstructure. We expected that inclusion of Cu NPs forming 3D networks should increase the 16


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electrical conductivity of (PU/Si)N nanocomposites. Despite some oxidation of Cu NPs, they indeed formed extended metallic networks (Figure 1 and S2) in the LBL assembled composites (Figure 2 and 4, and S5-7 and 9-12) via spontaneous assembly into chains with lattice-to-lattice connectivity.121,122,126,129 The conductivities of the (PU/Cu)N and (PU/Cu/Si)N composites were evaluated in great detail (Figure 6). The (PU/Cu)100 film showed a sheet resistance as low as 30 ohm and conductivity of 84 S/cm, respectively.

The addition of Si MPs certainly reduced the conductivity of the

[(PU/Si)(PU/Cu)m]N films. To demonstrate the materials engineering with LBL and improve the conductivity of the resulting nanocomposites, we increased the assembly cycle number m for (PU/Cu)m. When m was increased from 1 to 4, the conductivity of 90 bilayer films increased by an order of magnitude from 0.35 to 12 S/cm. Previous studies demonstrated the unusually high tensile strength and strain of neat PU films and PU-based composite films.84,147,148

The PU films obtained by

vacuum assisted filtration (VAF) showed high ultimate tensile strength of 21.3 MPa and tensile strain of 274.7% (Figure 7). After incorporating the Cu NPs and Si particles into the PU films, the resulting composites are weaker than other biomimetic layered composites from identical components.86,89,92,107

However, (PU/Cu)210

showed ultimate tensile strength of 11.7 MPa and tensile strain of 129.8%, and [(Pu/Si)(Pu/Cu)]90 showed ultimate tensile strength of 10.7 MPa and tensile strain of 110.8%, which is much higher than other coatings and composites for Si anodes. This degree of stretchability should be sufficient to accommodate the volume change of Si component during battery cycling while strongly adhering to both Si and Cu particles. 17



3.3. Battery performance with anodes from [(PU/Si)(PU/Cu)m]N composites. We prepared half coin cells by placing the flexible and conductive [(PU/Si)(PU/Cu)m]N composite membranes on current collectors.

After the first pre-cycling at a rate of

0.10 C (1 C = 4000 mA/g), all cells were further cycled 300 times at a rate of 1.0 C (Figure 8A-B). For the [(PU/Si)(PU/Cu)2]N film electrode, the first discharge capacity was 669 mAh/g, which was almost twice that of graphite (375 mAh/g). After the pre-cycling at 0.10 C for three cycles, the Coulombic efficiency improved to 97.3%. Although the capacity decreased to 333 mAh/g after 300 cycles, the average capacity fading rate was only 0.17% per cycle. Moreover, the average capacity fading rates of the [(PU/Si)(PU/Cu)3]N and [(PU/Si)(PU/Cu)4]N film electrodes were 0.16% and 0.15% per cycle, respectively. This low capacity fade can be attributed to the stretchability of the PU and Cu combined film, which effectively accommodated the Si volume change and maintained good conductivity due to the 3D conductive Cu NP networks. To further investigate the capacity fade, we obtained charge-discharge voltage curves of the [(PU/Si)(PU/Cu)4]N film electrode at 1.0 C rate (Figure 8C). These classical sloping and smooth profiles corresponded to the lithium insertion/extraction in Si.7,118,149–152 The discharge capacities decreased from 1284 mAh/g at the 1st cycle to 1252, 820, 735, and 687 mAh/g at the 2nd, 100th, 200th, and 300th cycles, respectively, giving average capacity fading rates of 0.36%, 0.12%, and 0.07% per cycle at the 1st, 2nd, and 3rd hundred cycles, respectively. The high capacity of 687 mAh/g after 300 cycles at 1.0 C rate was comparable to the 18


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recently-reported values (Table 1).5,11,12,14,43,149–163 The rate of capacity fade also decreased with the increasing cycle numbers, which indicates most likely that a subset of Si MPs in the composite has a smaller propensity to detach from the conductive network of Cu NPs in the course of lithiation/delithiation processes.

Besides, the

capacities greatly improved when increasing the assembly cycle of m in the [(PU/Si)(PU/Cu)m]N composites (Fig. 8A). It should also be noted that, throughout all the 300 cycles, our LBL-designed electrodes displayed single and average Coulombic efficiencies of 96.6—100.0% and 99.0—99.4% (Fig. 8D), respectively, which are higher than or competitive with other Si-based LIBs to date even those made with most sophisticated nanomanufacturing techniques (Table 1).5,11,12,14,43,149–163 We ascribed the high Coulombic efficiencies of the composite anodes to balanced combination of required essential mechanical and charge transport properties of the engineered material.

Additional protection of Si MPs by the PU coating layers (or

matrices) could suppressed excessive formation of solid electrolyte interface (SEI) layers.164 The rate performance, one of the most important properties for LIB applications, was also investigated. Testing the rate capacities of the three composite anodes (Figure 8D) revealed that the [(PU/Si)(PU/Cu)4]N film electrode exhibited simultaneously higher capacity and rate capacity than the battery cells based on [(PU/Si)(PU/Cu)2]N and [(PU/Si)(PU/Cu)3]N. After the rate increased from 0.10 to 2.0 C, the [(PU/Si)(PU/Cu)4]N film retained 43% of its original capacity, which was higher than the capacity retention of the [(PU/Si)(PU/Cu)2]N (29%) and 19



Page 20 of 48

[(PU/Si)(PU/Cu)3]N (37%) film electrodes. However, when the Si MP electrodes were prepared by a slurry coating method with carbon black and PVDF binder, they showed much lower retentions of 2—8% (Figure S13). In contrast, after 50 cycles at a rate of 0.10 C again, the battery cells with [(PU/Si)(PU/Cu)2]N, [(PU/Si)(PU/Cu)3]N, and [(PU/Si)(PU/Cu)4]N anodes showed high capacity retentions of 66%, 78%, and 80%, respectively. Compared to the data presented in the Introduction, this is quite remarkable. We ascribed the greatly improved electrochemical properties (e.g., high capacity, low capacity fading and high Coulombic efficiency) to the LBL composite

film

design

(Scheme

electro-conductivity of the composite

2):

(1)

the

high

stretchability

and

films can effectively accommodate the

volume variation of the Si particles upon lithiation/delithiation processes and facilitate the Li+/e– charge transfer during the charge-discharge cycling; (2) the formed nano/micro-sized pores in the composite films can also relieve the Si volume change and simultaneously offer low-resistance pathway for Li ions; (3) the additional protection of the Si particles by the flexible PU coating layers (or PU matrices) inhibited the excessive side reactions and much more formation of high-resistance SEI films on the Si surface. We used electrochemical impedance spectroscopy (EIS) to investigate charge transfer and Li-ion diffusion in electrodes to better understand the electrochemical performance of the [(PU/Si)(PU/Cu)m]N composite film electrodes (Figure 9A). The electrochemical impedance spectroscopy (EIS) circles in high-frequency and medium-frequency regions and the inclined line in low-frequency region 20


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corresponded to the formation of the SEI film, its contact resistance, the charge-transfer impedance on the electrode/electrolyte interfaces,165,166 and the Li-ion diffusion process within electrodes, respectively. After 60 charge-discharge cycles, the [(PU/Si)(PU/Cu)4]N film showed the smallest semicircle, while the diameter of the [(PU/Si)(PU/Cu)4]N film electrode remained almost identical to that seen after cycle 0; the [(PU/Si)(PU/Cu)4]N film electrode showed more vertical lines than the other two electrodes. These results indicated that, when increasing the m in [(PU/Si)(PU/Cu)m]N, the films exhibited lower charge transfer resistance and Li-ion diffusion resistance. This was because that increasing the m resulted in more Cu NPs forming conductive networks in the films, which led to the enhanced conductivity of the LBL composites (Figure 6). Moreover, the flexible composites accommodated the change in volume of Si during the cycling processes due to the stretchability of PU (Figure 7).84,147,167 Finally, the nano/micro-sized pores formed in the composite films (Figure 4 and S5-12) not only relieved the volume change, but also provided a low-resistance pathway for the Li-ions. These characteristics all synergistically contributed to the high cycle performance. A proposed equivalent circuit for the half cell model (Figure 9B) was used to obtain a quantitative analysis of the charge transport on the LBL composite films.165,166,168 Re is electrolyte resistance, Cf and Rf are capacitance and resistance of the SEI film, respectively, Cdl and Rct are double-layer capacitance and charge-transfer resistance, respectively, and Zw is the Warburg impedance, related to the Li-ion diffusion process. The Rf of the [(PU/Si)(PU/Cu)2] film anode increased from 21



62.6 to 107.5 ohm by 71.7% after 60 cycles at 1.0 C, however, the Rf of the [(PU/Si)(PU/Cu)4] film anode only increased from 28.4 to 33.3 ohm by 17.3% (Table 2). These results furtherly exposed that when increasing m, more robust Cu NP conductive networks formed in the PU/Cu/Si composite films, thus suppressing the Si volume change, film destruction, pore formation and SEI generation during the charge-discharge cycling. To gain insight into the intercalation reaction between the electrode and Li-ions, we calculated exchange current densities according to equation (1):165,168,169 (1) where R is the gas constant, T is absolute temperature (in K), n is the number of transferred electrons, and F is the Faraday constant. Electrochemical kinetic parameters for the film electrodes are summarized in Table 2. The io of the [(PU/Si)(PU/Cu)2] film decreased from 8.2  10-5 A cm-2 to 6.0  10-5 A cm-2 after 60 cycles. However, the io of the [(PU/Si)(PU/Cu)4] anodes decreased only by 1.0  10–5 A cm–2 less than those of the other electrodes after 60 cycles, which was attributed to the enhanced electrical contacts and facilitated migration of the Li-ions within the electrode.

4. CONLUSIONS LBL-made composites from three components, namely Cu NPs, Si MPs, and PU macromolecules, enable multiscale design of materials with a combination of properties defining the utility and competitiveness of the high capacity batteries. Cu 22


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NPs form 3D networks with lattice-to-lattice connectivity during LBL assembly and impart efficient electron transport to the target material.

Si MPs capable of Li

storage impart high battery capacity, whereas PU component with characteristically high stretchability impart high cycle life. Realization of all these properties in the resulting composite would not be possible, however, by simple mixing of the components or other methods of composite preparation associated with poorly controlled phase separation or difficult-to-bridge electron tunneling gaps between the constituent particles. Given high toughness, elasticity and structural uniformity of free-standing [(PU/Si)(PU/Cu)m]N compared to other composite architectures including porous films, vacuum deposited multilayer stacks, and slurry-deposited composites, one can envision that a version of this process can be realized in a roll-to-roll format. Current methods of rapid LBL deposition enable scalable and continuous production of rolled composites.16,89,95,112,113,115

Drastic relaxation of

thermodynamic restrictions of phase separation as well as universality and versatility characteristic of LBL opens the road toward development of a general toolbox for materials engineering of materials with multiple property requirements.

Supporting Information Available: SEM images of the redispersed Cu NPs; UV-Vis absorbance spectra of the original and redispersed Cu NPs; photographs, SEM and TEM images, and size distribution of the Si particles; zeta potentials of the PU, Cu NP, and Si particle solutions; SEM images of the PU/Cu/Si films, LBL assembled on glass slides and Cu foils; rate capabilities of Si particle slurry coated electrodes; AES 23



spectra of the PU/Cu/Si films LBL assembled on Cu foils.

5. ACKNOWLEDGMENTS. We thank the Electron Microscopy and Analysis Laboratory (EMAL) in University of Michigan for its assistance of electron microscopy. We greatly acknowledge Dr. Zhongrui Li at EMAL for the acquisition and data analysis of AES spectra. We also thank the Phoenix Memorial Laboratory in University of Michigan for the assembly and electrochemical tests of the Li+ coin cells. We are also greatly grateful to China Scholarship Council and Northwestern Polytechnical University for the scholarships and China Postdoctoral Science Foundation (203127) to L. Liu. This project is supported by Samsung as a part of GRO program. This material is based upon work partially supported by the Center for Solar and Thermal Energy Conversion, an Energy Frontier Research Center funded by the U.S. Department of Energy, Office of Science, and Office of Basic Energy Sciences under Award Number #DE-SC0000957. The authors would like to than National Science Foundation for NSF the grant #1538180 titled “Layered Composites from Branched Nanofibers for Lithium Ion Batteries”.

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FIGURES AND TABLES:

Scheme 1. Schematic diagram of a layer-by-layer assembly process involving three components: PU, Si MPs, and Cu NPs; the latter undergo self-assembly into three-dimensional networks upon deposition.

35



Scheme 2. Schematic presentations of the morphology/structure changes of (A) commonly-used Si particle anodes and (B) LBL-assembled PU/Cu/Si composite films during the electrochemical cycling process. The Si particles with polyurethane protection are easily to crack into smaller pieces due to the huge volume variation upon lithiation/delithiation and much more SEI formation on the Si particle surface.

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Figure 1. TEM images of (A-B) original and (C-D) re-dispersed Cu NPs. Size histograms of (E) original (size: 6.90 ± 1.02 nm) and (F) re-dispersed (size: 9.45 ± 1.45 nm) Cu NPs obtained by counting NPs in TEM images. (G) Photographs of (a) 0.30 mg/mL original, and (b) 2.0, (c) 0.30, and (d) 0.03 mg/mL re-dispersed Cu NPs.

37



Figure 2. (A) UV-Vis absorbance spectra of the (PU/Cu)N composite assembled on both sides of glass slides for various N. (B) Photograph of the (PU/Cu)N composites with N equal to (a) 60, (b) 40, and (c) 20. (C-D) Top-view, (E-F) back-view, and (G-H) cross-sectional SEM images of the (PU/Cu)110 composite. 38


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Figure 3. (A) UV-Vis absorbance spectra of the (PU/Si)N composites grown on both sides of glass slides with various N. (B) Photograph of a glass slide with (PU/Si)50 grown on both sides. (C-D) top-view and (E-F) cross-sectional SEM images of the (PU/Si)50.

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Figure 4. Photographs of (A) glass slide supported and (B) free-standing [(PU/Si)(PU/Cu)m]N composites with m=1, 2, 3, and 4, whereas N = 28, 27, 27, and 21, for (a), (b), (c), and (d) respectively. (C-D) Top-view, (E-F) back-view, and (G-H) cross-sectional SEM images of the [(PU/Si)(PU/Cu)]100 composites.

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Figure 5. TEM images of the scratched (A-B) (PU/Cu)110 and (C-D) [(PU/Si)(PU/Cu)]100 composite films from glass slides.

41


5000

Sheet resistance (ohm)

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

Page 42 of 48

(PU/Cu) (PU/Si)(PU/Cu)1 (PU/Si)(PU/Cu)2 (PU/Si)(PU/Cu)3 (PU/Si)(PU/Cu)4

-RESIS.- -COND.- ID

4000 3000

1500

200 180 160 140 120 100 80 20 15

1000

10

500

5

0

Conductivity (S/cm)


0 20

40

60

Bilayer number

80

100

Figure 6. The dependences of sheet resistance and conductivity on the bilayer numbers of the (PU/Cu)N and [(PU/Si)(PU/Cu)m]N composites. The total bilayer numbers are N and (1+m)×N for the (PU/Cu)N and [(PU/Si)(PU/Cu)m]N films, respectively. We assumed that each bilayer film exhibited the same thickness based on the linear growth of the LBL composites. The sheet resistance was divided by the film thickness to give the conductivity.

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25 20

Stress (MPa)

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


PU film

15 PU/Cu film 10

PU/Cu/Si flim

5 0

0

50

100

150

Strain (%)

200

250

300

Figure 7. Stress-strain curves of the PU, (PU/Cu)210, and [(PU/Si)(PU/Cu)]90.

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Figure 8. Dependence of (A) cell capacities and (B) Coulombic efficiencies on the number of charge-discharge cycles of the [(PU/Si)(PU/Cu)m]N anodes at 1.0 C rate after 3 charge-discharge cycles at 0.10 C rate. (C) Charge-discharge voltage curves of the (PU/Si)(PU/Cu)4 composite anodes cycled at the 1st, 2nd, 100th, 200th, and 300th between 2.0 and 0.01 V (vs. Li+/Li) at 1.0 C rate. (D) Rate capabilities of the [(PU/Si)(PU/Cu)m]N composite anodes at various charge-discharge rates between 0.10 and 2.0 C.

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Figure 9. (A) Nyquist plots of the [(PU/Si)(PU/Cu)m]N composite film electrodes over a frequency range of 100 K—0.01 Hz after 0 and 60 charge-discharge cycles at 1.0 C rate. (B) Equivalent circuit of the half-cell model.

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Table 1. Performance comparison of the recently-developed Si-based anode materials Cycling performance Morphology/structure

Preparation method

Ref. (Year)

Capacity Coulombic efficiency

Cycles

Current rate

300

1.00 C

This work

(mAh/g) 96.6–100.0% Flexible PU/Cu/Si composite film

Layer-by-layer assembly

(Sin.);

687 99.0-99.4% (Ave.)

Flexible Si NP/CNT film

Rod coating

1711

≥98%

50

0.10 C

155 (2010)

Flexible Si NP/graphene film

Vacuum filtration

2656

~97%

150

0.33 C

156 (2011)

Si ribbons on Cr/Au/Cr/PDMS substrate

Thermal evaporation

3498

92–95%

500

0.25 C

149 (2012)

Flexible Si NP/C nanofiber fabric

Electrospinning

600

100

1.00 C

157 (2013)

Cu-Si core-shell nanotube array

Template method

1506

98.44% (Ave.)

400

0.20 C

43 (2014)

Si nanotubes

Electrospinning and metallothermic reduction

765

~99.5%

280

0.50 C

158 (2015)

Si nanoparticle-graphene composite

Top-down dispersion and bottom-up synthesis

920

≥98.6%

600

0.12 C

159 (2016)

Si NP@void@graphene

Melting self-assembly and CVD

1287

~100%

500

0.12 C

160 (2016)

Si NP/graphene foam

Freeze-drying

1295

180

0.12 C

161 (2016)

Si nanolayer-embedded graphite/C hybrid

CVD

496

≥99.5%

100

0.50 C

162 (2016)

Granadilla-like Si/carbon composite

Template method

1100

~100%

200

0.06 C

163 (2016)

Si/edge-activated graphite composite

Ni-catalyzed hydrogenation and CVD

521

~100%

50

0.50 C

11 (2017)

Si/Ta layer on porous Cu substrate

Physical vapor deposition

700

~100%

200

0.50 C

12 (2017)

Si@amorphous TiO2 coating

Sol-gel

1720

~100%

200

0.10 C

14 (2017)

Porous Si encapsulated by organic C/Si

Magnesiothermic reduction and carbonization

~2800

99.5% (Ave.)

100

0.25 C

153 (2017)

Si NP/graphene composite

Drop-casting and annealing

1450

99.5%

100

0.08 C

154 (2017)

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Table 2. Kinetic parameters of the PU/Cu/Si LBL-engineered composites Electrode type

Cycle number

Rf (ohm)

Rct (ohm)

io (A cm-2)

[(PU/Si)(PU/Cu)2]

0

62.6±2.2

312.5±8.9

8.2  10-5

60

107.5±3.4

426.7±9.7

6.0  10-5

0

41.2±1.8

219.5±4.2

1.2  10-4

60

68.2±3.2

294.4±4.8

8.7  10-5

0

28.4±1.2

158.2±5.1

1.6  10-4

60

33.3±1.5

175.7±2.2

1.5  10-4

[(PU/Si)(PU/Cu)3]

[(PU/Si)(PU/Cu)4]

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Graphic abstract:

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Materials Engineering of High-Performance Anodes as Layered

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