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Self-Relaxant Super-Elastic Matrix Derived from C60 Incorporated Sn Nanoparticles for Ultra-High-Performance Li-Ion Batteries Ryanda Enggar Anugrah Ardhi, Guicheng Liu, Xuan Minh Tran, Chairul Hudaya, Ji Young Kim, Hyunjin Yu, and Joong Kee Lee ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.8b01345 • Publication Date (Web): 04 Jun 2018 Downloaded from http://pubs.acs.org on June 4, 2018

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Self-Relaxant Super-Elastic Matrix Derived from C60 Incorporated Sn Nanoparticles for Ultra-HighPerformance Li-Ion Batteries Ryanda Enggar Anugrah Ardhi,†,‡,┴ Guicheng Liu,*,†,┴ Xuan Minh Tran,†,‡ Chairul Hudaya,†,‡,# Ji Young Kim,†,§ Hyunjin Yu,†,║ and Joong Kee Lee*,†,‡ †

Center for Energy Storage Research, Green City Research Institute, Korea Institute of Science

and Technology (KIST), Hwarang-ro 14-gil 5, Seongbuk-gu, Seoul 02792, Republic of Korea. ‡

Division of Energy and Environment Technology, KIST-School, Korea University of Science

and Technology (UST), Hwarang-ro 14-gil 5, Seongbuk-gu, Seoul 02792, South Korea. §

Department of Chemical and Biomolecular Engineering, Yonsei University, 50 Yonsei-ro,

Sodaemun-gu, Seoul 120-749, Republic of Korea. ║

Department of Material Science and Engineering, Korea University, Seoul 136-701, Republic

of Korea.

┴

These authors contributed equally to this work.

* Corresponding authors: [email protected], [email protected] (Liu); [email protected] (Lee). Tel.: +82 2 958 5252, Fax: +82 2 958 5229. #

Present address: Department of Electrical Engineering, Faculty of Engineering, Universitas

Indonesia, Kampus Baru UI, Depok 16424, Indonesia. ACS Paragon Plus Environment

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ABSTRACT Homogeneously dispersed Sn nanoparticles approximately ≤ 10 nm in a polymerized C60 (PC60) matrix, employed as the anode of an Li-ion battery, are prepared using plasma-assisted thermal evaporation coupled by chemical vapor deposition. The self-relaxant super-elastic characteristics of the PC60 possess the ability to absorb the stress–strain generated by the Sn nanoparticles and can thus alleviate the problem of their extreme volume changes. Meanwhile, well-dispersed dot-like Sn nanoparticles, which are surrounded by a thin SnO2 layer, have suitable inter-particle spacing and multilayer structures for alleviating the aggregation of Sn nanoparticles during repeated cycles. The Ohmic characteristic and the built-in electric field formed in the inter-particle junction play important roles in enhancing the diffusion and transport rate of Li-ions. SPC-50, an Sn-PC60 anode consisting of 50 wt% Sn and 50 wt% PC60, as confirmed by energy-dispersive X-ray spectroscopy analysis, exhibited the highest electrochemical performance. The resulting SPC-50 anode, in a half- cell configuration, exhibited an excellent capacity retention of 97.18%, even after 5000 cycles at a current density of 1000 mA g-1 with a discharge capacity of 834.25 mAh g-1. In addition, the rate-capability performance of this SPC-50 half-cell exhibited a discharge capacity of 544.33 mAh g-1 at a high current density of 10 000 mA g-1, even after the current density was increased 100-fold. Moreover, a very high discharge capacity of 1040.09 mAh g-1 was achieved with a capacity retention of 98.67 % after 50 cycles at a current density of 100 mA g-1. Futhermore, a SPC-50 full-cell containing LiCoO2 cathode exhibited a discharge capacity of 801.04 mAh g-1 and an areal capacity of 1.57 mAh cm-2 with a capacity retention of 95.27 % after 350 cycles at a current density of 1000 mA g-1.

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KEYWORDS: polymerized C60, self-relaxant super-elastic characteristics, fast charge transport, ultra-high performance anode, lithium-ion batteries

The rapidly increasing market for electronic devices has resulted in a greater demand for Liion batteries (LIBs) owing to their higher capacity, better cycling stability, and improved rate capability.1,2 Technological developments have led to the continuous miniaturization of electronic devices into microelectronic-scale devices such as microelectromechanical systems, pocket flash cameras, and implantable smart medical devices. This has led to the miniaturization of LIBs to become small-scale energy-storage systems through the application of thin-film architecture.3–5 A thin-film anode exhibits a higher Li-ion (Li+) diffusivity and electronic (e-) conductivity owing to the shorter Li+ and e- path lengths, compared with thick films or bulk anodes, enabling facile Li+ transport and diffusion to improve the power density. The implementation of a metallic material as an LIB anode, which utilizes an alloy/de-alloy mechanism with Li+ during charge/discharge (lithiation/de-lithiation) processes, is important in substituting for conventional graphite-based anode materials because of the associated increase in the energy consumption of recent electronic devices. Sn has emerged as an alternative anode material because it has an almost three-fold theoretical gravimetric capacity (≈ 994 mAh g-1) compared with a conventional graphite-based anode (≈ 372 mAh g-1).6 Structural instability, such as particle fracture and pulverization, and loss of e- and Li+ pathways have been identified as major drawbacks of Sn-based anodes owing to significant volume expansion/contraction changes (≈ 259%), which are caused by the formation/decomposition of a large Li22Sn5 alloy phase during the charge/discharge step.7 To address the aforementioned drawbacks and to realize improved cycling stability and rate capability, several researchers have proposed a reduction in the Sn particle size to the


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nanoscale.8,9 Although the particle-size reduction has successfully suppressed the volume-change problem for the first several tens of cycles, relative to that for a bulk anode, and has improved the cycle life of an Sn anode, the very small-sized Sn nanoparticles have a tendency to aggregate in subsequent cycling processes to form larger particles, which also leads to pulverization of the Sn anode. Therefore, this strategy cannot further diminish the structural instability of the Sn anode during a cycling test involving several hundreds of cycles.10 Furthermore, the unstable solidelectrolyte interphase (SEI) has also deformed and been broken by the expansion and contraction of the Sn during cycling tests. In consequence, accumulated SEI thickening has still occurred at the freshly exposed Sn crack interface as a result of those volume changes. Such SEI thickening will impede the Li+ transport/diffusion into/from the Sn host. Such a nanoscale Sn-based anode needs to be further structurally and interfacially engineered. The use of an inert metallic material that is not reactive to Li+, such as Cu,11 Ni,12 or Ag,13 as an inactive matrix for creating an intermetallic composite with Sn nanoparticles is another way of alleviating the volume-change problem of an Sn anode. However, its cycle life is improved by sacrificing the capacity of the Sn anode as a result of the non-contribution of those inactive matrix materials to the capacity. The process of dispersing Sn nanoparticles into an active matrix material that is also reactive toward Li+, which is referred to as an “active-active nanocomposite system,” has been extensively investigated as a means of further suppressing the structural instability and SEI thickening without significantly sacrificing the capacity of the Sn anode. The sp2-hybridized carbonaceous materials, including graphene,14,15 graphite,16,17 carbon nanotubes (CNTs),18–20 and carbon nanowires,21 have been widely used as active matrix materials of Sn-based anodes owing to their good physical properties such as high tensile strength and high ductility, in addition to their low reactivity with Sn and high Li+ diffusion coefficient.22 Moreover, the electrochemical


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performance and mechanical properties of the C matrix material can be improved by introducing defect-rich sites, such as point defects or so-called “Stone-Wales defects,” single vacancies (SV), and multiple or double vacancies (DV) to form an amorphous carbon (a-C) structure.23 This defect site enables the creation of a more active site for Li+ adsorption and also enhances the mechanical elasticity. Ideally, an a-C-based matrix material implemented in a metallic host anode such as Sn possesses some of following features: (1) It has a highly sp2-hybridized portion to ensure high electrical/Li+ conductivity and facile charge transport and (2) it has a finely tuned defects ratio, with a proportional sp3-hybridized site, to provides more Li+ active sites and to enhance the elasticity of the C-based matrix. However, realizing a highly uniform dispersion of Sn nanoparticles with a proper inter-particle spacing distribution in a C matrix is also important for avoiding the aggregation of Sn nanoparticles, which again causes particle pulverization, by providing sufficient space to compensate for the volume stress-strain during the cycle process and ensuring high-capacity retention over a very large number of cycles. With respect to Sn-based C nanocomposite anodes prepared by an ex situ process such as powder-phase ball milling,24 liquid-phase synthesis,25 or the hydrothermal method,26 uniform dispersion of Sn nanoparticles with an appropriate interparticle spacing distribution in a C matrix is difficult to achieve and control. To overcome this limitation, a direct in situ synthesis route, such as that using a gas-phase or plasma reaction, is proposed.27–29 It is clear that all of the aforementioned C-based matrices could be used to arrange a regular dispersion of Sn nanoparticles, as well as to buffer the volume change and aggregation of the Sn anode material with low current densities. However, the Sn-C nanocomposite thin-film anode system still needs to be improved to achieve higher-capacity retention at higher current


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rates, enabling it to undergo several thousands of cycles. Hence, the uniform spacing distribution of the Sn nanoparticles in a C matrix requires further improvement. Few reports have mentioned the transport mechanism of Li+ at the interface between two or more types of metallic or semiconductor materials, namely at inter-particle junction, in an Sn-C nanocomposite anode. It is important to understand the behavior of interfacial contact between two different materials with different charge transport properties because of the different mechanisms associated with their charge transport. Owing to the contact between two different materials, at thermal equilibrium, Fermi-level alignment occurs, depending on the work function positions of those materials based on the basic semiconductor theory. Moreover, several contacts can be formed, such as an Ohmic, p-n, referred to as p-n junction, or Schottky contact. Their contact behavior can be used as a charge transporter or charge blocker. In metal/semiconductor contacts, wether p-type or n-type semiconductors, an Ohmic or Schottky contact can be formed. In an Ohmic contact, which is a non-rectifying electrical junction, charges can easily move through the junction in both directions owing to the absence of a barrier charge diffusion region (BCDR), providing low Ohmic resistance and high electrical conductivity. In contrast, owing to the presence of a BCDR near the junction, a Schottky contact allows a charge to flow only in one direction, and it has a high Ohmic resistance. However, a p-type semiconductor/n-type semiconductor contact consists of majority free carriers on their own regions, such as a negative (positive) charge for n-type (p-type) semiconductor. The majority carriers provide a built-in electric field in which the direction is from p-type semiconductor material with a positive charge to n-type semiconductor material with a negative charge. Moreover, these majority free carriers of p-type and n-type semiconductor


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materials also drive external charges, such as Li+ and e- in terms of battery electrodes, to pass through the p-n junction easily. To obtain some of the aforementioned a-C matrix material, we selected a nanostructure fullerene (C60) material that is more reactive toward Li+ and significantly easier to prepare as a thin-film anode because it can be prepared without the addition of a catalyst at a substrate temperature of 250 oC. For the preparation of thin film, CNT and graphene require a catalyst such as Fe or Ni at a substrate temperature range of 600-1000 oC despite the use of high-energy plasma (microwave plasma).30–31 Elimination of the catalyst and the lower substrate temperature make it easier to prepare for the processing. To the best of our knowledge, C60 has not been investigated as an active matrix material for Sn-based anodes. Moreover, C60 is a promising Cbased material that can be easily polymerized under high-pressure or a plasma condition at a certain temperature to form a polymerized C60 (PC60), an a-C material, and it exhibits some promising properties, such as very good electrical and mechanical properties for the fast transport of Li+ and e- into/from the Sn host during charge/discharge cycles and ensuring high mechanical stability, respectively.32–36 Meanwhile, based on our previous results, a PC60 coating material could also diminish the formation of a side reaction to form an SEI layer owing to its natural properties, which further improves the Coulombic efficiency and reduces the irreversible capacity in the first cycle. The self-relaxant super-elastic characteristic of this material also alleviates the volume-change problems.37,38 Inspired by the promising properties of PC60 material and classical semiconductor theories, in the present study, we set out to demonstrate an Sn-PC60 nanocomposite thin-film anode for LIBs prepared by a hybrid in situ thin-film deposition system, a radio-frequency plasma-assisted thermal evaporator coupled with a chemical vapor deposition (RFPATE-CVD) method via a gas-


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phase and plasma reaction, combined with interfacial engineering management of Sn, the thin amorphous SnO2 (a-SnO2) layer, and PC60 materials within the Sn-PC60 nanocomposite system. Ion bombardment from Ar+ plasma was used to introduce physical defect sites, which are disordered amorphous of open-ball structures and dangling bonds sites, on C60 material while forming a PC60 matrix. The Sn-PC60 nanocomposite possesses three main advantages: (1) the PC60 matrix has self-relaxant super-elastic characteristics, and so can buffer the huge cyclic volume change of the Sn host material, while the well-dispersed Sn nanoparticles with suitable inter-particle spacing and a multilayer structure prevent the aggregation of Sn nanoparticles over thousands of charge/discharge cycles; (2) the ternary metallic/n-type semiconductor/p-type semiconductor structure of Sn/a-SnO2/PC60 plays an important role in the creation of the Ohmic/built-in electric field (BEF) effect in their binary junction to enhance and promote the diffusion and transport rate of Li+ into/from the Sn host during charge and discharge processes, and to maintain a high capacity at high currents; and (3) the dot-like Sn nanoparticles dispersed in the PC60 matrix minimize the length of the Li+ diffusion pathway, and also provide a high capacity at a high current density.

RESULTS AND DISCUSSION Figure 1a shows the architectural concept of the Sn-PC60 nanocomposite anode, with a thickness of ≈ 220 nm, consisting of dot-like Sn nanoparticles and the PC60 matrix. The Sn material is the main host material for Li+. The minimal thickness of this anode enables it to develop a diffusion coefficient of Li+ and electronic (e-) conductivity owing to the Li+ and e- path lengths being shorter than those of thick anodes. This also reduces the possibility of a volume change. Meanwhile, the C material, that is, the PC60 matrix (a secondary host), can be used as a


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buffer to suppress any volumetric change during charge/discharge processes owing to its selfrelaxant super-elastic characteristics.37,38 The design goal for the Sn-PC60 anode, which incorporates well-dispersed dot-like Sn nanoparticles with suitable inter-particle spacing and a multilayer structure, was to maintain high structural stability while retaining the cycling performance and rate capability, even under harsh conditions, such as several thousands of cycles and high current densities. Moreover, instead of C60, which is an n-type semiconductor,39,40 the PC60 which is a p-type semiconductor material with positive charges (holes) as the majority carriers,33 was selected together with a thin a-SnO2 layer and Sn material to establish Ohmic/BEF characteristics for enhancing the diffusion rate of Li+ within the Sn/aSnO2/PC60 binary junction. It should be noted that a-SnO2 is an n-type semiconductor material with negative charges (electron) as the majority carriers. To illustrate this more clearly, Figure 1b shows a detailed model with only one Sn nanoparticle surrounded by its environment and includes a schematic diagram of the working mechanism (top panel) and an energy diagram (bottom panel) of the inter-particle interfaces of our Sn-PC60 anode system. The a-SnO2 thin layer shows more clearly the formation of ternary structures of Sn/a-SnO2/PC60. In the initial state, the metal/n-type semiconductor/p-type semiconductor, referred to simply as m/n/p contact within the ternary structure of the Sn/SnO2/PC60 creates a Fermi energy level (EF) pinning. This results in the emergence of band bending, and finally the appearance of an Ohmic contact at the Sn/SnO2 inter-particle junction and a p-n semiconductor contact at the SnO2/PC60 inter-particle junction. During the charge process, the driving external voltage narrowing the space charge region (SCR) in the p-n junction and the built-in electric field at the a-SnO2/PC60 region drives Li+ from the PC60 matrix to the Sn particle owing to electrostatic interactions. In detail, the majority free carrier in the PC60 matrix


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is repelled and forces the Li+ to move from the PC60 matrix to the Sn particle owing to the electrostatic interaction. Moreover, the majority free carrier in the a-SnO2 attracts Li+ from the PC60 matrix to the Sn particle owing also to the electrostatic interaction. Hence, the Li+ can easily move quickly through the junction. After the nanocomposite anode has been fully lithiated, the Sn particle and PC60 matrix are transformed to LixSn and LiyC, respectively, and the a-SnO2 thin layer phase is also transformed to LixSn and Li2O according to Equations 1–3, below.22,25 Sn + Li + e → ← Li Sn

(0 ≤ ≤ 4.4)

(1)

C + Li + e → ← Li C

(y = randomly value)

(2)

SnO + 4Li + 4e → ← Sn + 2Li O

(3)

Because the LixSn particle phase will take substantially more Li+ than the LiyC matrix phase, the work function of the LixSn phase, from the Sn nanoparticle and the a-SnO2 layer, will become relatively lower than that of the LiyC. As a consequence, the majority free carrier in the PC60 matrix is transformed from positive charge to negative charge, while the majority free carrier in a-SnO2 thin layer is transformed from negative charge to positive charge. Now, the direction of the built-in electric field is reversed. During the discharge process, the Li+ in the LixSn particlephase region can easily and rapidly cross through the Sn/a-SnO2/PC60 junction to the matrix side owing to the electrostatic force from the built-in electric field. Therefore, this Sn/a-SnO2/PC60 binary junction, is expected to promote fast Li+ diffusion and transport during the charge/discharge process. The Li+ in the LixSn particle-phase region can also easily and rapidly cross through the LixSn/SnO2/LiyC binary junction to the matrix side. Therefore, the Ohmic/BEF effect in their binary junction enables fast Li+ diffusion and transport during the charge/discharge process.


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Figure 1. (a) Architecture of the Sn-PC60 nanocomposite anode possesses stable structure after many cycles due to the presence of a suitable inter-particle space distribution of Sn nanoparticles and a self-relaxant super-elastic characteristics of PC60 matrix, (b) schematic illustration (top panel) and schematic energy diagram (bottom panel) of metallic Sn (βSn)/n-type semiconductor SnO2/p-type semiconductor (PC60) contact at the initial, lithiation, and de-lithiation state. The majority free carriers of PC60 and SnO2 are the positive charge and the negative charge, respectively. The Ohmic characteristic in Sn/SnO2 junction and the built-in electric field formed in between SnO2 and PC60 is enabling fast charge transfer with a low Ohmic resistance. EVac, EF, Eg, CBM, VBM, and W are vacuum energy level, Fermi energy level, Band-gap energy, conduction-band minimum, valenceband maximum, and space charge region (SCR), respectively.


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To realize the aforementioned anode structure, illustrated in Figures S1a–c, the RFPATECVD method was employed; and some key parameters for anode deposition are summarized in Table S1. For simplification, all of the tested anodes are abbreviated as listed in Table S2. To intuitively exhibit the internal structure and physico-chemical properties of the Sn-PC60 anode, in the present study, the SPC-50 sample was chosen for comparison with pristine Sn and pristine PC60 with their lowest mass loading configuration as shown in Table S3. The deposition process of the SPC-50 anode material using the RFPATE-CVD method is shown in Supporting Information Video S1. To clearly illustrate the elemental distribution and structure of the SnPC60 anode, surface and cross-sectional analyses were carried out using high-angle annular dark field (HAADF) images, elemental maps, and transmission electron microscopy (TEM) images (in the bright-field mode). As shown in Figures 2a–c, the dissemination of Sn nanoparticles, which have suitable inter-particle spacing in a PC60 matrix on the surface detected by TEM image and elemental maps, corresponds to the HAADF images in Figure S2a. From the surface high-resolution (HR)-TEM images shown in Figures 2b–c, the sizes of almost all of the dot-like Sn particles with a highly crystalline phase are ≤ 10 nm. To understand the spatial structure of this anode, HAADF images, elemental maps, and TEM images of the cross-sectional view were obtained, as shown in Figures 2d–f and S2b. In Figures 2d and S2b, the anode appears to have a multilayer structure, i.e., the Sn particles and PC60 matrix are stacked layer-by-layer. However, as shown in Figure 2e, there are many smaller Sn nanoparticles between the main layers, which also confirm the homogeneously dispersed multilayer Sn nanoparticle structure with an appropriate inter-particle space distribution in the SPC-50 anode. In agreement with Figure 2c, almost all the Sn particles are ≤ 10 nm from their top to bottom surfaces, as shown in Figures 2d–f. Moreover, the surface morphologies and


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thicknesses of the pristine PC60, pristine Sn, and SPC-50 were further observed using a fieldemission scanning electron microscopy (FE-SEM) spectrometer. As shown in Figures S3a–c, the grain particle sizes of SPC-50 (≤ 10 nm) and pristine PC60 (≤ 5 nm) are significantly smaller than that of pristine Sn (150–200 nm). The cross-sectional SEM images shown in Figures S3d–f indicate that all three samples prepared using the same RFPATE-CVD method possess similar thicknesses of ≈ 220 nm, which is reconfirmed by the X-ray photoelectron spectroscopy (XPS) depth profile shown in Figure S4. They are also in good agreement with the cross-sectional TEM image shown in Figure 2d. The Sn particles shown in Figures 2c and 2f are surrounded by the amorphous thin layer. In agreement with our proposed ternary structures of Sn/a-SnO2/PC60 shown in Figure 1b, this thin amorphous layer is an a-SnO2 material, as determined by the XPS spectra on the 3d core level in Figure S5a and Fourier transform infrared (FT-IR) spectra (Figure S5c). In detail, the XPS Sn 3d spectra of SPC-50 in Figure S5a shows that the O-Sn-O peaks are located at 486.4 eV and 494.9 eV, whereas the FT-IR spectra of SPC-50 in Figure S5c also detected an O-Sn-O vibrational peak at 600 cm-1. As indicated by the FT-IR spectra in Figure S5b, the O element likely originated from the C60 precursor (powder) at 1100 cm-1 for the C-O vibrational peak and at 1718 cm-1 for the C=O vibrational peak. The formation mechanism of this ternary structure of Sn/a-SnO2/PC60 contains the a-SnO2 thin layer with uniform dispersion, as explained in Figure S6. Briefly, during the deposition of the SPC-50 anode, the Ar+ RF plasma breaks the Sn-CH3 bonds (∆Hd(Sn–CH3)= 227 kJ/mol) in the tetramethyltin Sn (TMT) precursor and creates Sn radical (Sn●). The high-energy electron from the Ar+ plasma source also collides with the TMT precursor, which leads to the formation of Sn4+ through ionization of electrons from the Sn●.


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This Sn● and Sn4+ are deposited on the Cu substrate first owing to the heavy atomic mass of Sn (118.71 g mol-1). Further, this Sn● is deposited at the middle part of the Sn nanoparticle, whereas the Sn4+ part is deposited at the edge of the Sn nanoparticle and creates a positive charge at the outer side of the Sn particle. This positive charge creates a Van der Waals repulsive force between the Sn nanoparticles and separates them at considerable distances. At the same time, the Ar+ plasma also breaks some of C-C (∆Hd(C-C)= 346 kJ/mol) and C=C (∆Hd(C=C)= 602 kJ/mol) bonds in the C60 precursor, which leads to the formation of the Cy-Ox and Cy=Ox components to form an open-ball structure with dangling bonds at the edge of the C component. These components, which also contain radical components, fall down to the substrate and fill-in the spaces among the Sn nanoparticles. Considering that the O element possesses the highest electron affinity compared with Sn and C elements, and the Sn element possesses the lowest electron affinity, the oxygen edge of Cy-Ox and Cy=Ox is reduced and contains a negative charge. This result occurred through acceptance of a free ionization electron from the Sn4+ and the Ar+ plasma source, and the oxygen edge bonds with the Sn4+ on the outer Sn nanoparticle to form SnO2. Therefore, the size of the Sn particles surrounded by the outer thin layer of the a-SnO2 is limited to nanometer size, at ~5 nm. The dangling bond was revealed by FT-IR measurement in the open-ball PC60 structure, as shown in Figure S5c. FT-IR spectrum of pristine PC60 and SPC-50 anode verify the existence of the O-H peak located from 3000 cm-1 to 3600 cm-1 possibly owing to the adsorption of water molecules. While this peak is absence in the FT-IR spectra of C60 powder precursor, as exhibited in Figure S5b.32 Therefore, the water molecules was only adsorbed on surfaces of the pristine PC60 and SPC-50. According to our previous investigations, we also found those peaks of C-C


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and C=C were weakened by the increasing RF plasma power.36 Namely, higher power can break the bonds of C-C and C=C to produce the open-ball structure.

Figure 2. (a–c) Surface and (d–f) Cross-sectional view of TEM image of SPC-50, (g) SAED pattern from the area inside the dashed red box in Figure 2d, (h) XRD pattern of pristine Sn, pristine PC60, SPC-50, and the corresponding XRD reference pattern of β-Sn (tetragonal Sn) from JCPDS database number 04-0673, and (i) Raman spectra of pristine Sn, C60, pristine PC60, and SPC-50. As shown in Figure 2g, seven β-Sn crystalline planes, indexed as (200), (101), (220), (301), (112), (400), and (312), are detected by the selected area electron diffraction (SAED) pattern of


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the area inside the dashed red box in Figure 2d.15 Owing to the presence of a bright-white dotted pattern, the fast Fourier transform (FFT) patterns shown in Figure S7 also clearly verify the β-Sn crystalline planes in the SPC-50. In contrast, because of the absence of the crystalline C60 pattern in the SAED and the existence of a milky blurred FFT ring pattern, the PC60 matrix is identified as being in an amorphous state.15 This amorphous state of PC60 was created by the Ar+ plasma bombardment that occurred during the deposition process. Moreover, only highly tetragonal crystalline planes of β-Sn are observed in the X-ray diffraction (XRD) patterns (Figure 2h),9,15 which is consistent with the SAED and FFT patterns. The Raman spectra shown in Figure 2i further confirm the transformation of C60 to PC60, in that the characteristic peaks of the C60 material, such as Hg (eight fivefold modes of symmetry) and Ag (two non-degenerate modes of symmetry), were completely missing from the PC60 and Sn-PC60 anodes. Instead, two bands commonly observed in amorphous C material, namely, the D- and G-bands, appeared.15,19 To better understand the contact behavior of a Sn/a-SnO2/PC60 junction, the electrical properties of the PC60 material must first be understood. Ultraviolet photoelectron spectroscopy (UPS) was used to reveal the work function of the PC60 and Sn materials. The work function can be determined from Equation 4, as 41 = EHe1 – Ecut-off

(4)

where , EHe1, and Ecut-off are the work function, radiation energy of the UPS source (21.2 eV), and cut-off energy, respectively. As shown in Figure 3a, the cut-off region of the UPS spectra reveals that pristine PC60 material possesses a lower work function (4.08 eV) compared to pristine Sn (4.40 eV). Meanwhile, the valence-band region of the UPS spectra in Figure 3b show that the valence-band maximum (VBM) values for PC60 and the Sn material are located at 0.46 eV below EF and 0.22 eV above EF, respectively, indicating that the PC60 and Sn are p-type


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semiconductor and metallic materials, respectively. Further, the optical band gap (Eg) of the PC60 material was tested by the ultraviolet visible (UV-vis)measurement and calculated using Equation 5, as 42 Eg = (h × c) /λ

(5)

where h, c, and λ are the Planck constant (4.143 × 10-15 eV), the speed of light (3 × 108 m s-1), and the cut-off wavelength, respectively. As shown in Figure 3c, the optical band gap of PC60 is 1.73 eV, which further confirms the p-type semiconductor nature of the PC60 material. Because the work function of Sn (4.40 eV) is lower than that a-SnO2 (theoretically ~5.00 eV) and higher than that PC60 (4.08 eV) is lower than that of Sn (4.40 eV), it results in the creation of an Ohmic/BEF behavior at the Sn/a-SnO2/PC60 junction, which supports our assumption shown in Figure 1b. Figure 3d summarizes the UPS and UV-vis results of PC60 shown in Figures 3a–c.


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Figure 3. UPS analysis results for the pristine Sn and pristine PC60 at (a) cut-off region and (b) valence-band region, (c) UV-vis absorption spectra of pristine PC60, and (d) schematic illustration of energy band diagram of PC60 having p-type semiconductor behavior. To validate our proposed approach to improving battery performance, we further investigated the electrochemical performances of Sn-PC60 with different Sn contents and compared the results with those of pristine Sn and pristine PC60 anodes. The exact Sn content of all these anodes was obtained from the energy-dispersive X-ray spectroscopy (EDS), as depicted and summarized in Figure S8 and Table S4. The half-cell battery performance at varying current


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density values was first tested using the coin-type half-cell LIBs structure depicted in Figure S9a. The mass loading for all the tested anode materials is summarized in Table S3. In the present study, to prevent confusion related to the charge and discharge processes, Figure S9b shows the working concept of a half-cell LIB system with an Sn-PC60 anode. Figure 4a shows the rate-capability performances of pristine Sn, SPC-50, and pristine PC60. The initial discharge capacity of SPC-50 (1054.11 mAh g-1) is slightly higher than that of pristine Sn (990.93 mAh g1

) and pristine PC60 (904.35 mAh g-1) at a current density of 100 mA g-1. This may be attributed

to the capacity contribution of a-SnO2 thin layers surrounding the Sn nanoparticles with a theoretical capacity of 1494 mAh g-1.43 The SPC-50 anode exhibits the highest and most stable rate performance among the three anode materials for every current density regime. Specifically, at 10 000 mA g-1, which is the highest current density tested, the discharge capacity of the SPC50 remains at 544.33 mAh g-1, which is 10.42 and 2.57 times higher than those of pristine Sn (52.27 mAh g-1) and pristine PC60 (211.96 mAh g-1), respectively. Similarly, Figure S10a shows that the SPC-50 anode also demonstrates the best rate performance among the Sn-PC60 anodes with varying Sn content, at every current density regime. The highest rate performance of SPC50 originated at the lowest Li+ transport resistance and the largest Li+ diffusion depth within the grain particles owing to the presence of Ohmic/BEF behavior at the Sn/a-SnO2/PC60 junction, resulting in facile Li+ transport and diffusion at high current densities. Furthermore, when a current density of 100 mA g-1 was reapplied, as shown in Table S5, the capacity retention rate increased along with the increasing PC60 content after 50 cycles. When the PC60 content is ≥ 50 wt%, the capacity retention rate increases to ≥ 98.67% and does not change significantly. It should be noted that, at the 50th cycle, the SPC-50 anode exhibits the highest discharge capacity (1040.09 mAh g-1) of all anodes, as shown in Figure S10b. This result is attributed to a


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combination of the self-relaxant super-elastic characteristics of the PC60 matrix and the homogeneously dispersed dot-like Sn nanoparticles with the suitable inter-particle spacing in the PC60 matrix, thus maintaining the structural stability and Li+ pathway availability.


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Figure 4. (a) Rate capability performance of pristine Sn, SPC-50, and pristine PC60 at a different current density, (b) plot of Coulombic efficiency value from Figure 4a, (c) plot of the average value of capacity retention at a different current density which is extracted from Figure 4a, (d) cycle performance of pristine Sn, SPC-50, and pristine PC60 at 1000 mA g-1 of current density, and (e) cycling ability of SPC-50 for 5000 cycles at 1000 mA g-1 of current density, (f) Cycle performance of SPC-50 anode with different mass loadings at -1

1000 mA g of a current density charge/discharge condition, and (g) plot of areal capacity st

th

of SPC-50 anode with different mass loadings at the 1 and 350 cycle which is extracted from Figure 4f. The gravimetric capacity was calculated based on the mass of active material (Sn for pristine Sn, PC60 and Sn for SPC-50, and for PC60 pristine PC60), while the areal capacity was calculated based on the surface area of active material (1.13 cm2). Results from Figure 4b, extracted from Figures 4a and S11a–c, show that both SPC-50 and pristine PC60 possess higher Coulombic efficiencies compared to pristine Sn at the initial cycle. This is a result of the natural ability of the PC60 material to suppress the formation of a high-level side reaction to form a thick SEI layer owing to the electrolyte decomposition.36–38 Furthermore, in the following two cycles, the Coulombic efficiencies of both the SPC-50 and pristine PC60 are significantly improved, and exhibit much smaller irreversible capacities than that of pristine Sn, thus indicating the ability of the self-relaxant super-elastic characteristics of the PC60 material to sustain a stable anode structure, as well as to suppress the presence of mechanical cracks and the further formation of new SEI layers. Voltage profile and cyclic voltammetry (CV) analyses (Figure S11 and Table S6) revealed, in more detail, the electrochemical activity related to the suppression of a high-level side reaction and mechanical cracking in an SPC-50 anode. As shown in Figures S11a–c, during the first


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charge process, the voltage plateau appeared in the case of the pristine Sn anode, starting between 2.50 V and 0.60 V. However, the voltage plateau of the pristine PC60 and SPC-50 did not appear until around 1.00 V, and between 0.70 V and 0.40 V, respectively. These results are in good agreement with the locations of the A peaks shown in the results of the CV analysis in Figures S11d–e. The emergence of the broad voltage-plateau region of the pristine Sn during the first charge process is believed to be a result of the formation of a thick SEI layer caused by the high-level side reaction which mainly contains LiF components, while this did not occur in the case of the SPC-50 anode.44,45 This indicates the ability of the PC60 material to suppress the high levels of electrolyte decomposition in the initial state, as previously mentioned. On the other hand, the ability of the self-relaxant super-elastic characteristics of the PC60 matrix to sustain a stable structure and suppress mechanical cracking in the SPC-50 anode during cycling is evidenced by the voltage profiles and the CV peak intensities of the Sn alloy (B)/de-alloy (B') regions during the 2nd and 3rd cycles being similar, while pristine Sn does not. Figure 4c shows the average capacity retention at each current density for the pristine Sn, pristine PC60, and SPC-50, as extracted from the rate capability performance shown in Figure 4a. The SPC-50 anode exhibits the highest capacity-retention ratio, which is even higher than that of the pristine PC60, at all the tested current densities. This capacity retention extracted from the rate performance of SPC-50 further highlights the benefit of the Ohmic/BEF behavior within the Sn/a-SnO2/PC60 inter-particle junction in the SPC-50 anode to facilitate facile Li+ movement at high current densities. To compare the cycling stability of the pristine Sn, SPC-50, and pristine PC60 anodes, the cycle performance was tested at a current density of 1000 mA g-1, as depicted in Figure 4d and Table 1. The results show that SPC-50 (858.46 mAh g-1) exhibit the highest discharge capacity


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compared to the pristine Sn (832.00 mAh g-1) and pristine PC60 (557.89 mAh g-1). This is attributed to the combination of the high-level capacities of the Sn and the PC60, and possibly to the capacity contribution of the a-SnO2 thin layers surrounding the Sn nanoparticles as previously mentioned. As expected, the discharge capacity of the pristine Sn completely degrades in the first 150 cycles with a very low capacity retention (7.71%) and Coulombic efficiency (87.52%) as a result of the usual volume change problem. In contrast, as we previously estimated, the SPC-50 anode exhibits a superior capacity retention (97.35%) and Coulombic efficiency (99.54%) after 350 cycles. As shown in Figure 4e, the SPC-50 anode maintains a significantly stable capacity retention (97.18%) and Coulombic efficiency (99.91) with a final discharge capacity of 834.25 mAh g-1, even after 5000 cycles at a current density of 1000 mA g-1. This stable cycling performance of the SPC-50 anode is attributed to the stable mechanical structure that was achieved through the use of a self-relaxant super-elastic PC60 matrix material with a uniform dispersion of Sn nanoparticles in a dot-like architecture. Table 1. Summary of the cycle performance results in Figure 4d, e. Discharge capacity (mAh g-1) Material name

Capacity retention (%)

Coulombic efficiency (%)

1st cycle

150th cycle

350th cycle

5000th cycle

150th cycle

350th cycle

5000th cycle

1st cycle

150th cycle

350th cycle

5000th cycle

Pristine Sn

832.00

64.15

-

-

7.71

-

-

46.60

87.52

-

-

SPC-50

858.46

836.40

835.71

834.25

97.43

97.35

97.18

63.74

99.49

99.54

99.91

Pristine PC60

557.89

554.88

554.49

-

99.46

99.39

-

72.11

99.83

99.87

-

To further verify the super-elastic ability of the PC60 matrix material, we provide a half-cell cycle performance of pristine Sn, SPC-50, and pristine PC60 anodes with higher mass loading. The preparation parameters used to prepare these higher mass loading anodes are summarized in Table S7. As shown in Figures 4f and S12 and summarized in Table S8, by increasing the mass


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loading, the cycle performance of the SPC-50 anode still surpasses those of the pristine Sn and pristine PC60 with similar discharge capacity and Coulombic efficiency compared with those of the SPC-50 anode with the lowest mass loading. The areal capacities of the pristine Sn, SPC-50, and pristine PC60 are extracted from their gravimetric capacities as shown in Figures 4f,g, S12a,b; the results are also summarized is Table S8. The areal capacity of the SPC-50 anode at 1.96 mg cm-2 of mass loading exhibited 1.64 mAh cm-2 at the initial cycle and 1.60 mAh cm-2 after the 350th cycle, which is the highest compared with those of the pristine Sn (1.64 mAh cm-2 at initial cycle and 0 mAh cm-2 after 150 cycles) and pristine PC60 (1.10 mAh cm-2 at initial cycle and 1.08 mAh cm-2 after 150 cycles) at similar mass loading conditions. In order to further clarify that the SPC-50 anode can be used to facilitate facile Li+ transport and movement, the rate performance of the SPC-50 anode has been compared with other rate performances of Sn-based anodes available in literature, as shown in Figure 5a and summarized in Tables S9 and S10. In fact, the rate performance of the SPC-50 anode clearly outperforms the other Sn-based anodes reported in literature, prepared through in situ (upper panel)7,20,28,29,46–51 and ex situ processes (bottom panel).6,15,52–69 Furthermore, as shown in Figure 5b and summarized in Tables S9 and S10, the results of the cycle performance obtained with the SPC50 also surpass the capacity retention of the current Sn-based anodes prepared using in situ (left panel)7,20,27–29,46–51,70–79 and ex situ processes (right panel).6,15,52–69,80


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Figure 5. Comparison of (a) rate capability performance and (b) cycle performance of our SPC-50 vs. Sn-based anode for LIBs application that available in literatures. To further clarify the mechanical stability of the SPC-50, we used SEM to conduct a postmortem analysis of the post 150-cycle test samples of the pristine Sn anode and post 350-cycle test samples of the SPC-50 and pristine PC60 anodes. Figures 6a, b shows that, after undergoing only 150 cycles, the pristine Sn exhibits severe mechanical crumbling and pulverization owing to the very large volume change. The aggregation of Sn nanoparticles is also seen here. In contrast, both SPC-50 (Figures 6c, d) and pristine PC60 (Figures 6e, f) maintain their structures without any mechanical degradation even after undergoing 350 cycles, as a result of the ability of the PC60 matrix to alleviate the volume-change problem during charge/discharge cycles owing to its self-relaxant super-elastic characteristics. The usual mechanical cracking problem of the Snbased anode remains absent from the SPC-50 anode after performing the cycle test for 5000 cycles, as shown in Figure S13. Further supporting the results shown in Figures 6a-f, the atomic force microscopy (AFM) results shown in Figure S14 and Table S11 show that both the SPC-50 and pristine PC60 exhibits a much more stable structure after undergoing 350 cycles, relative to the pristine Sn that only undergoes 150 cycles. This can be determined from the surface


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roughness values. Initially, before the cycle tests, the pristine PC60 (12.16 nm) and SPC-50 (15. 73 nm) have smaller roughness values because their grain sizes are smaller than that of the pristine Sn (22.92 nm). After the cycle tests, the pristine Sn exhibits a much higher surface roughness (115.57 nm) than the pristine PC60 (14.32 nm) and SPC-50 (17.92 nm), due to the presence of mechanical cracking and aggregation of the Sn nanoparticles during charge/discharge cycles. Moreover, the percentage of surface roughness before/after the cycle test of the pristine Sn is also much higher at 40.04 %, compared with others, at 9.54 % for the pristine PC60 and 14.00 % for SPC-50.

Figure 6. Surface SEM image of (a and b) pristine Sn, (c and d) SPC-50, and (e and f) pristine PC60. The SEM analysis on Figure 6a, c, and e, and Figure 6b, d, and f were performed before and after cycle test (150 cycles for pristine Sn, and 350 cycles for SPC-50 and pristine PC60), respectively. The (g) C 1s, (h) O 1s, and (i) F 1s XPS core level spectra


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from surface of PC60, pristine Sn, and SPC-50 with corresponding peak deconvolution. These XPS spectra were extracted from Figure S15b. It is well known that an anode with a high mechanical stability is helpful for the formation of a stable SEI layer without SEI thickening during the cycle test. To reinforce this assumption, we performed another post-mortem analysis, XPS analysis, on the surface of the pre-cycled test sample of the pristine Sn, SPC-50, and pristine PC60 (Figure S15a), and compared it with the post-cycle test sample, the post-150-cycle test sample of the pristine Sn anode, and the post-350cycle test samples of the SPC-50 and pristine PC60 anodes (Figure S15b). The full-range XPS surveys of the post-cycle test samples shown in Figure S15b indicate several F, P, and Li peaks, which are believed to originate from electrolyte decomposition during the cycles to form an SEI layer, which is absent from the surface of the pre-cycled test samples (Figure S15a). To support our estimation, the peak deconvolution of pristine Sn, SPC-50, and pristine PC60 at the C 1s, O1s, and F 1s core-level regions, shown in Figure S15b, was carried out, as shown in Figures 6g–i. Overall, in these core level regions, seven chemical bonds were found, which are attributed to the SEI species, the polyethylene oxide C-O bond from (PEO) oligomers ((-CH2CH2-O-)n) located at 286.5 eV and 533.3 eV, the C=O bond from Li alkyl carbonate (RCH2OCO2Li) at 287.6 eV and 530.6 eV, the O-C=O bond from carboxylate C at 289.1 eV and 532.2 eV, the O-C(=O)-O bond from Li carbonate (Li2CO3) at 290.1 eV and 531.5 eV, the Li-O bond from Li2O at 528.6 eV, the Li-F bond from LiF at 686 eV, and the LiPF6 salt residue at 688 eV.81–83 Interestingly, the deconvoluted peak intensities of the SEI species in the pristine Sn anode are the highest compared to those in the pristine PC60 and SPC-50 anodes. This result is attributed to the SEI layer thickening owing to the mechanical crumbling and pulverization of the pristine Sn anode during the cycles.


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To further investigate the SEI thickening in the pristine Sn, we conducted XPS measurement with different Ar+ sputter depths (with observations at different depths), at the F 1s core level spectra, as shown in Figure 7 and summarized in Table S12. The Ar+ sputter rate was 40 nm min-1 and the sputter step was 0.3 min (≈ 12 nm). After 1 charge/discharge cycle, the SEI thicknesses of the pristine Sn (Figure 7a), SPC-50 (Figure 7b), and the pristine PC60 (Figure 7c) appear to be similar. However, Figure 7d reveals the SEI thickness estimated from Figures 7a–c in more detail: The pristine Sn has a greater SEI thickness (48 nm) compared to SPC-50 (36 nm) and pristine PC60 (12 nm). This further confirms our Coulombic efficiency result, shown in Figure 4b, and the abovementioned results of CV analysis (Figures S11d–f). On the other hand, the estimated SEI thicknesses of the pristine Sn (after 150 cycles), and SPC-50 and the pristine PC60 (after 350 cycles) are clearly different. Further supporting the results shown in Figure 6i, the pristine Sn possesses a higher peak intensity of the F 1s core level spectra at the surface (at a depth of 0 nm) compared to the SPC-50 and pristine PC60. The F 1s spectra of the pristine Sn (Figure 7e) exhibits a hill-like shape, while the SPC-50 (Figure 7f) and pristine PC60 (Figure 7g) do not. Figure 7h clearly shows a large difference between the estimated SEI thickness of the pristine Sn (480 nm) and that of SPC-50 (36 nm) and pristine PC60 (12 nm).


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Figure 7. F1s XPS spectra of (a) pristine Sn, (b) SPC-50, and (c) pristine PC60 after 1 cycle test, (e) pristine Sn after 150 cycles test, and (f) SPC-50 and (g) pristine PC60 after 350 cycles test. Atomic concentration plots of F element with the estimated SEI thickness in pristine Sn, SPC-50, and pristine PC60 were extracted from (d) Figure 7a–c, and (h) Figure 7e–g. Ar+ sputter rate was 40 nm min-1 and the sputter step was 0.3 min (~12 nm). The significant difference in the estimated change in the SEI thickness of the pristine Sn after 1 and 150 cycles indicates that pulverized Sn nanoparticles are produced by the volume change during the charge/discharge cycles. In consequence, the SEI thickening will degrade the capacity of the LIBs. On the other hand, the self-relaxant super-elastic characteristics of the PC60 material allows the SPC-50 and pristine PC60 to exhibit a stable structure, leading to a stable SEI layer thickness during the charge/discharge cycle. The SEI layer thickening by the continuous accumulation of the SEI formed at the freshly exposed Sn crack interface is prevented in the case of the SPC-50 anode. The SEI layer thickening of the pristine Sn can be explained by steps (i) to (v), illustrated in Figure S16.


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The superior mechanical stability of SPC-50 anode over hundreds, even thousands, of cycles is strongly linked to the presence of a self-relaxant super-elastic PC60 matrix. Elastic C-based materials with Li+-active sites can be achieved by tuning their defective level.23,81 In the present study, we set a constant Ar+ plasma power of 300 W to create defect-rich sites in an open-ball C60 structure (PC60). To estimate the defect level of the pristine PC60 and SPC-50 anode materials, Raman analysis was conducted before and after the 350-cycle tests, as shown in Figure 8a and summarized in Table S13. As shown in Figure 2i, the broad band of all of Raman spectra could be deconvoluted into two peaks corresponding to the D- and G-bands.15,19 The defect level in the pristine PC60 and SPC-50 can be estimated by comparing the intensity ratio of the D- and G-bands (ID/IG).15,19,84 Prior to the cycle test, the SPC-50 anode possessed a higher ID/IG value (0.912) than the pristine PC60 (0.790), revealing more defect sites of the PC60 material in the SPC-50 anode,85–87 due to the inclusion of Sn nanoparticles,88 compared to the PC60 material in the pristine PC60 anode. This condition provides additional Li+ active sites and more kinetics of Li+ transport, further facilitates a higher discharge capacity with faster Li+ transport of the SPC-50 anode compared to those of the pristine PC60 anode (Figures 4a, d).23,89,90 After 350 cycles, their ID/IG values, 0.926 for the SPC-50 and 0.793 for the pristine PC60, do not seem to change significantly with the percentage change value of ID/IG before and after the 350-cycle tests exhibiting a difference of only about 1.535% for the SPC-50 and 0.380% for the pristine PC60. This further confirms the stability of the structure of the SPC-50 and pristine PC60 anodes after the cycle tests, and is in good agreement with the results of other post-mortem analyses (SEM and AFM).


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Figure 8. (a) Raman and (b) C 1s XPS core level spectra of pristine PC60 and SPC-50 before and after 350 cycles with corresponding peak deconvolution. It is understood that the elastic properties of a C-based material can also be achieved by tuning its sp3/sp2 ratio.84 If C-based materials have a much greater sp2 content, their electrical conductivity will be excellent owing to the presence of one empty orbital, allowing free electrons to move easily, while mechanically they are very brittle (such as in the case of a single-layer graphene and graphite). However, if C-based materials have a much greater sp3 content, their electrical conductivity will be quite low, although mechanically they will be very strong with a high Young’s modulus, tensile strength, and hardness (as in the case of diamond). The sp3/sp2 ratio of the PC60 material in an SPC-50 anode and the pristine PC60 anode can be estimated by


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using XPS analysis, before and after 350-cycle tests, at different sputter depths (Figure S17). The results are summarized in Table S14. Figure 8b shows the C 1s core level spectra of SPC50 and the pristine PC60 at a sputter depth of about 72 nm. Surprisingly, the PC60 material in SPC-50 and the pristine PC60 anode exhibits a similar sp3/sp2 ratio of about 0.29. The Young’s modulus and tensile strength obtained from this sp3/sp2 value can be estimated to be about ≥ 288.80 GPa and ≥ 28.88 GPa, respectively, representing the elastic behavior of the PC60 material.84 Moreover, the exact values of the Young’s modulus and tensile strength of the PC60 material in SPC-50 and the pristine PC60 anode can be further theoretically obtained using Equations 6 and 7 as 90–93 = 4350.6 + &' =

$.% !" #

(6)

(

(7)

$)

where E,

!" ,

and TS are the Young’s modulus, sp3 fraction obtained from XPS analysis

(Figures 8b and S17, and Table S14), and tensile strength, respectively. As summarized in Table S15, before the cycle test, the Young’s modulus and tensile strength of PC60 materials were calculated to be about 301.06 GPa and 30.11 GPa for pristine PC60 anode, and 298.82 GPa and 29.88 GPa for the SPC-50 anode, respectively. Furthermore, after 350 cycles test, the PC60 materials were still able to maintain similar Young’s modulus and tensile strength valuses, i.e. about 302.77 GPa and 30.28 GPa in pristine PC60 anode, and 296.44 GPa and 29.64 GPa in SPC50 anode, repsectively. Those values are much higher than the Young’s modulus and tensile strength of 50.00 to 52.34 GPa and 5.00 to 5.23 GPa for Sn material, and 24.70 GPa and 2.47 GPa for the Li4.4Sn material (i.e., the fully lithiated phase of the Sn anode), respectively.94–97 Moreover, these Young’s moduli and tensile strength behavior of the PC60 material shows the self-relaxant super-elastic characteristics. Furthermore, the superior rate capability performance


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(Figures 4a), cycle performance (Figures 4d-f), stable structure (Figures 6c, 6d, and S13), and thin and stable SEI layer (Figures 6g, 6h, 6i, and 7b, 7d, 7f, and 7h) of the SPC-50 anode are definitely originated to the presence of this self-relaxant super-elastic PC60 material. To verify our previous evidence related to the facilitation of fast Li+ transport and diffusion, as well as the high mechanical stability of the SPC-50 anode, the electrochemical impedance spectroscopies (EIS) and simulated results (based on the equivalent circuit shown in Figure 9a) of the pristine Sn (at the 1st and 150th charge), and SPC-50 and the pristine PC60 (at the 1st and 350th charge) in the fully lithiated condition are shown in Figure 9b and Table 2. The components of the equivalent circuit include Ohmic resistance (RS, at a very high frequency), which is also known as series resistance, SEI resistance (RSEI, in the high-frequency region), Faradaic charge-transfer resistance (RCT, in the medium-frequency region), and the Warburg diffusion component (WO, in the low-frequency region). Wo is divided into two parts: the Warburg charge-diffusion resistance (WO-R) and the Warburg charge-diffusion time (WO-T). Furthermore, the electrical conductivity (σ) of the anode material, the Warburg coefficient (σW, Figure S18), chemical-diffusion coefficient (Dch) of Li+ through the anode, and the exchange current density (j0) from Li+ inserted into the anode active material can be calculated using Equations 8–11 as 72,98 *=

+

(8)

,-. /01

234 = *5 6 ).% + 78 + 79: ; /

,:

@ 01 ? AB 9CD

(9)

E

(10)

,:

F) = G?,

(11)

-.

where t, Aac, Zre, ω, R, T, Aac, F, CLi, and n are thickness of anode materials (obtained from Figures S3 and S4), the surface area of the anode active material (1.13 cm2), the real impedance


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in the low frequency region, the angular frequency obtained from the low-frequency regime, the ideal gas constant (8.314 J mol-1 K-1), room temperature (298 K), Faraday’s constant (96 485 C mol-1), the initial Li+ molar concentration in the electrolyte (1 mol L-1), and the number of electrons transferred per molecule during the charge condition via the alloying mechanism (n = 1), respectively.99

Figure 9. (a) Equivalent circuit used for fitting impedance spectra, (b) Nyquist plots of pristine Sn, SPC-50, and pristine PC60 anode active material at 1st, and at 150th (pristine Sn)


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and at 350th (SPC-50 and pristine PC60) charge condition in a series of cycles for full-range (top panel) and enlarge (bottom panel) image. Figure 9b shows EIS analysis at the first charge condition, and provides some information regarding the enhancement of the Li+ transport properties in SPC-50 anode due to the presence of the Ohmic/BEF behavior within the Sn/a-SnO2/PC60 junction and the ability of the PC60 material to prevent the formation of a high-level side reaction. The order of the RS value (inset in Figure 9b), from the lowest to the highest, is as follows: SPC-50 (0.88 Ω) < pristine PC60 (1.47 Ω) < pristine Sn (2.47 Ω). This order shows the Ohmic/BEF behavior of the SPC-50 anodes at a Sn/a-SnO2/PC60 binary junction, also further verifies the facile e-/Li+ transport in an SPC-50 anode, which is in agreement with the rate capability performance in Figures 4a, 4c, and 5a. On the other hand, the RSEI of the pristine PC60 is the smallest (10.17 Ω) among the three samples, which further verifies the ability of the PC60 material to prevent a high-level side reaction forming a thick SEI layer, and is in good agreement with the F 1s XPS results shown in Figures 7a–d. Meanwhile, the beneficial effect of the Ohmic/BEF behavior within the Sn/a-SnO2/PC60 junction on the facile Li+ transport is also evidenced by the smallest RCT (20.16 Ω), WO-R (29.31 Ω), and WO-T (6.24 × 10-4 s) values of the SPC-50 compared to those of pristine Sn and pristine PC60. Moreover, the SPC-50 also possesses the highest σ (9.83 × 10-7 S cm-1), Dch (1.46 × 10-12 cm2 s-1), and j0 (1.27 × 10-3 A cm-2) values, further implying the highest Li+ transfer kinetics. Table 2 Summary of the EIS result in Figure 9b. Condition

After 1st Chrg.

Material name

RS (Ω)

RSEI (Ω)

RCT (Ω)

σ (S cm-1)

WO-R (Ω)

WO-T (s)

σW (Ω cm2 s-0.5)

Dch (cm2 s-1)

j0 (A cm-2)

Pristine Sn

2.47 63.14

232.09

8.39 × 10-8 147.70 2.49 × 10-1

1970.60

7.14 × 10-15 1.11 × 10-4

Pristine PC60

1.47 10.17

75.77

2.52 × 10-7

60.75

1.54 × 10-3

378.59

1.93 × 10-13 3.39 × 10-4

SPC-50

0.88 15.93

20.16

9.83 × 10-7

29.31

6.24 × 10-4

137.60

1.46 × 10-12 1.27 × 10-3


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After 150th Chrg.

Pristine Sn

Pristine PC60 After th 350 Chrg. SPC-50

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6.04 360.91 1264.57 1.54 × 10-8 1011.00

25.66

5763.20

8.34 × 10-16 2.03 × 10-5

1.52 10.66

77.17

2.48 × 10-7

59.83

2.25 × 10-3

476.20

1.22 × 10-13 3.33 × 10-4

1.00 16.44

22.54

8.79 × 10-7

30.61

6.96 × 10-4

175.03

9.05 × 10-13 1.14 × 10-3

The EIS analysis at the higher charge condition reveals suppression of the volume change of the SPC-50 anode (at the 350th charge condition) owing to the self-relaxant super-elastic ability of the PC60 material and the suitable Sn inter-particle spacing dispersion and its multilayer structure. In contrast, the pristine Sn (at the 150th charge condition) has the highest RS (6.04 Ω), which indicates the delamination process of the Sn material from the Cu current collector owing to the volume-change process as illustrated in step (iv) in Figure S16. The thickening of the SEI layer of the pristine Sn anode is also evidenced by the increase in RSEI (360.91 Ω) relative to that of the pristine Sn at the first charge condition (63.14 Ω), which agrees with the XPS results shown in Figures 6i, 7a, 7d, 7e, and 7h. Moreover, the highest RCT (1264.57 Ω), WO-R (1011.00 Ω), and WO-T (25.66 s), as well as the lowest σ (1.54 × 10-8 S cm-1), Dch (8.34 × 10-16 cm2 s-1) and j0 (2.03 × 10-5 A cm-2) of the pristine Sn, compared to the other values, further proves the occurrence of the volume-change problem during the charge/discharge cycles. The σ, Dch , and j0 values of the pristine Sn at the 150th charge condition also significantly improved relative to those of the pristine Sn at the first charge condition, in good agreement with the increment of its RSEI value, implying SEI thickening impedes Li+ movement and degrades the capacity of pristine Sn anode during cycles. In contrast, for all of the EIS parameters of the SPC-50 anode, the small difference in EIS parameters between the first and 350th charge conditions further verifies our proposed approach, as illustrated in Figure 1a: The highly stable mechanical structure resulting from the selfrelaxant super-elastic characteristics of the PC60 material successfully buffers the usual large


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volume-change problem of Sn materials during prolonged cycling tests. The highest σ (8.79 × 10-7 S cm-1), Dch (9.05 × 10-13 cm2 s-1), and j0 (1.14 × 10-3 A cm-2) values, and the smallest RCT (22.54 Ω), WO-R (30.61 Ω), and WO-T (6.96 × 10-4 s) values of SPC-50 compared to pristine Sn and pristine PC60 also further proves our proposed approach depicted in Figure 1b, relating to the presence of an Ohmic/BEF characteristic at a Sn/SnO2/PC60 inter-particle junction due to fast Li+ transport and movement, which agrees with rate capability performance of SPC-50 (Figures 4a, 4c, 5a, and S10).

Figure 10. (a) Cycle performance of full-cell configuration of the SPC-50 at a current density of 1000 mA g-1, and (b) voltage profile of the 1st and 2nd cycle which is extracted from Figure 10a. Additionally, the cycle performances of the pristine Sn and SPC-50 anodes were evaluated with a full-cell configuration comprising a commercial LiCoO2 (LCO) cathode. The N/P (negative to positive) ratio of active materials between the anode and the cathode is 1.1. First, the cycle performance of the LCO cathode was tested in the half-cell configuration. As shown in Figure S19, the LCO cathode possesses a discharge capacity of about 173.31 mAh g-1 at the initial cycle. However, the discharge capacity of LCO cathode decreases slightly to 157.63 mAh g-1 after 350 cycles with 90.95 % of the capacity retention. This result is possibly attributed to


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the structural change of the LCO cathode during charge/discharge cycles. As shown in Figures 10 and S20 and Table S16, the cycle performance of the SPC-50 is superior to that of the pristine Sn in the full-cell configuration even after 350 cycles. After 350 cycles, the SPC-50 can still deliver 801.04 mAh g-1 of discharge capacity and 1.57 mAh cm-2 of areal capacity, with 98.5 % of capacity retention. In contrast, the pristine Sn shows only 0.62 mAh g-1 of discharge capacity and 0.001 mAh cm-2, even after only 150 cycles and leads to only 0.12 % of capacity retention. Moreover, the pristine Sn is able to deliver a low capacity at initial cycle of only 512.28 mAh g-1 due to the formation of the thick SEI layer at the initial cycle because the LCO cathode provided the limited Li+ content. We believe that the application of the PC60 material and this thin-film architectural approach will be useful as a benchmark in the design of thin-film anodes of other high-capacity materials that incur similar drawbacks for LIB applications. CONCLUSIONS In summary, by applying the in situ synthesis route using the combination of a gas-phase and plasma reaction via the RFPATE-CVD method, Sn-PC60 anode material was successfully prepared as a thin-film anode for LIBs. This was achieved by dispersing Sn nanoparticles into a defect-rich PC60 matrix with a suitable inter-particle spacing and a multilayer structure. The performance of the Sn-based anode was dramatically improved with the inclusion of the PC60 matrix in terms of enhanced structural and mechanical stabilities, and in Li+ transport and diffusion ability. Three multifunctional advantages were realized from utilizing the PC60 matrix to improve the performance of the Sn host in the SPC-50 anode. (1) The self-relaxant superelastic characteristics of the PC60 matrix and the highly dispersed Sn nanoparticles with suitable inter-particle spacing and its multilayer structure effectively buffer the volumetric change over thousands of charge/discharge cycles. This resulted in high mechanical stability and very high


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capacity retention (97.18%) with a high discharge capacity (834.25 mAh g-1) after 5000 cycles at a current density of 1000 mA g-1. (2) The ternary metallic/n-type semiconductor/p-type semiconductor structure of Sn/a-SnO2/PC60 played an important role in the creation of Ohmic/ BEF characteristics in their binary junction to enhance and promote the diffusion and transport rate of Li+ into/from the Sn host during charge and discharge processes. In addition, it maintained a high capacity at high currents, thereby maintaining a high capacity (544.33 mAh g1

) under a high current-density condition (10 000 mA g-1). (3) The considerably small size of the

Sn nanoparticles (≤ 10 nm) minimized the Li+ diffusion pathway and also provided a high capacity at a high current density. The present study not only points to the importance of the PC60 material and its properties as a promising matrix material for a nanocomposite anode for LIB applications, but also explains the transport mechanism of Li+ within the junction of the nanograin composite. Furthermore, the results of this study provide a guide for optimizing the inter-particle interface in nanocomposite materials to ensure facile Li+ transport in LIBs anodes.

EXPERIMENTAL SECTION Deposition of anode active material. Commercial tetramethyl tin liquid (TMT, Sn(CH3)4, UP Chemical) and C60 powder (99.99% sublimed, Sigma Aldrich, CAS No. 99685-96-8) were used as the Sn and PC60 precursors as received, respectively. As illustrated in Figure S1 and Supporting Information Video S1, Sn-PC60 nanocomposite anode active materials were deposited on pre-cleaned Cu substrates (15-µm-thickness) at different percentage weight ratios of Sn and PC60 (wt% Sn/PC60) based on the deposition parameters listed in Table S1. A tungsten boat containing C60 powder was placed inside the RFPATE chamber and was connected to a direct current (DC) electrical thermal source to heat the C60 powder precursor, while the TMT


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liquid was placed inside the CVD reactor. The RFPATE base pressure was brought to 1 × 10 % Torr (≈ 1.33 × 10-3 Pa) using a turbo pump, and was supported by a rotary vane-pump. The Cu substrate was heated to 250 °C. After introducing 30 sccm of Ar gas into the PATE chamber controlled by a mass flow controller (MFC), the chamber pressure was adjusted to 42 mTorr (working pressure) and the 13.56-MHz RF source generator was set to a constant 300 W. Subsequently, the TMT from the CVD was injected into the chamber at a different bubble pressure supported by 10 sccm of Ar gas. At the same time, the tungsten boat was heated to the elevated temperature, as supplied by a DC bias voltage to evaporate the C60 powder, and the deposition of the Sn-PC60 was started. The heating procedure of the tungsten boat during C60 evaporation to form a multilayer structure of Sn-PC60 is presented in Figure S1c. Pristine PC60 and a pristine Sn thin-film anode were also prepared to enable a comparison with the Sn-PC60, using the same procedure as that described above, but without the C60 evaporation process (for the pristine Sn) or the TMT injection process (for the pristine PC60). Device fabrication. All electrochemical measurements were tested using coin-type LIBs (series CR2032), and were assembled inside a dry room with a well-maintained dew point (100.2 °C) and temperature (20 °C). Prior to assembly into coin-type half-cells, a 53-µm-thick Li metal foil counter electrode, all of the anodes, and a polypropylene separator were cut into circular shapes with diameter of 1.65 cm, 1.2 cm, and 1.8 cm, respectively, whereas the mass loading of all of the deposited anode active materials were carefully weighed using an Analytical Plus microbalance (Ohaus AP250D, readability ≈ 0.01 mg). 1 M of LiPF6 (PANAX ETEC Co., Ltd.), dissolved in ethylene carbonate (EC), dimethyl carbonate (DMC), and ethyl methyl carbonate (EMC) at a vol% ratio of 1:1:1, was used as a non-aqueous electrolyte with no further modification or addition of additives. For the full-cell test, the anode material was paired with


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LiCoO2 (LCO) cathode. The LCO cathode was prepared by mixing 90 wt% of LCO powder, 5 wt% of Denka black, and 5 wt% of polyvinylidene fluoride (PVdF) in an N-methyl-2pyrrolidone (NMP) solution. An N/P ratio of ~1.1 was designed between the anode and cathode materials for the full-cell. Device analysis. The rate-capability performances were tested using a battery tester (Maccor battery tester system, series 4000) at a cut-off voltage range of 0.01–3.00 V and at current density values of 100 mA g-1, 300 mA g-1, 500 mA g-1, 700 mA g-1, 1000 mA g-1, 2000 mA g-1, 3000 mA g-1, 5000 mA g-1, and 10 000 mA g-1, while the cycle performances were measured at a constant-current density of 1000 mA g-1 and the same cut-off voltage range as that used in the rate-capability test. The EIS and CV measurements were carried out using a multichannel potentiostat analyzer (Bio-Logic Science Instruments, VMP3 type). The EIS measurements were conducted at an amplitude perturbation signal of 0.10 mV and at a frequency ranging from 106 to 10-1 Hz under the fully lithiated condition, when the open-circuit potential (OCV) of the LIBs reached 0.01 V (vs. Li/Li+). The non-linear least-square method provided in Zview software (Scribner Associates, Inc.) was used for the fitting process of the Nyquist plot, and corresponded to equivalent-circuit modeling. However, the CV spectra were obtained from the first to third cycles at a cut-off voltage range of 0.01–3.00 V and at a scanning rate of 0.10 mV s-1. The cycle test, rate-capability test, EIS, and CV measurements were conducted at ambient temperature. Material characterization. SEM measurements were performed using a Nova 200 NanoLab at an accelerating voltage of 10 kV (FEI Company). TEM, SAED, HAADF and elemental map, and FFT measurements were conducted using a Talos F200X TEM instrument (FEI Company) at an accelerating voltage of 200 kV. Sample cutting for the preparation of TEM samples was carried out using a focused ion beam (FIB) Nova nanoLab 600 (FEI company). AFM


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measurement was conducted by using an AFM spectrometer, XE-100 model, (Park Systems Company) with silicon nitride (Si3N4) cantilever probes operated in non-contact mode. XRD measurements were carried out by using an Ultima IV X-ray diffractometer (Rigaku company) assisted by monochromatic CuKα radiation (λ = 0.15406 nm) with a scanning 2θ range of 25– 90°at a scan rate of 0.006° s-1. The operating current and voltage of the XRD measurements were maintained at 15 mA and 30 kV, respectively. EDS measurements were carried out using a Teneo volume scope (VS). Raman measurements were performed using micro Raman spectrometer (Renishaw Company) with 532 nm of the laser source wavelength. Pure Gaussian function fitting was used to obtain peak deconvolution of the Raman spectra. Ultraviolet photoelectron spectroscopy (UPS) measurements were performed using an AXIS-NOVA (Kratos) system at a base pressure of 5 × 10−8 Torr with HeIα radiation (hv = 21.22 eV) and 0.585 eV of pass energy resolution. XPS measurements were conducted using a PHI 5000 VersaProbe instrument (Ulvac-PHI company) system equipped with a monochromatized Al Kα radiation (h ν = 1486.6 eV) source, operated at a constant voltage of 15 kV. MultiPak and XPSPEAK software (Ulvac-phi Company) was used for the XPS data processing. The full-range XPS spectra were operated at a wide scan-pass energy of 117.4 eV. The peak subtraction for the core-level peak analysis was performed using a Shirley-type background.100,101 Gauss-Lorentzian function fitting, with an 80%/20% of Gaussian/Lorentzian portion, was used to obtain peak deconvolution of the XPS spectra. For the XPS depth profile, Ar-ion sputtering at 3 kV was used with a sputter rate of 40 nm min-1 on a standard SiO2 sample, and the pass energy of the depth profile was 46.95 eV. FT-IR measurement was carried out by using a Nicolet Almega (Thermo Electron Company). ASSOCIATED CONTENT


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1.1. Supporting Information Available The Supporting Information is available free of charge on the ACS Publications website at DOI: _____________________ Schematic diagram illustration of experimental setup, deposition process of the Sn-PC60 nanocomposite, LIB structure, and working concept of the LIB used in this study; additional characterization data and electrochemical experiments; and comparison with previous results (PDF) AUTHOR INFORMATION 1.2. Corresponding Authors * Email: [email protected], [email protected] (G Liu); [email protected] (JK Lee). 1.3. Author Contribution ┴

R.E.A.A. and G.L. contributed equally to this work.

1.4. Present Address #

Present address: Department of Electrical Engineering, Faculty of Engineering, Universitas

Indonesia, Kampus Baru UI, Depok 16424, Indonesia. 1.5. Notes The authors declare no competing financial interest. ACKNOWLEDGMENT This work was supported by research grants of NRF (NRF-2017R1A2B2002607) funded by the National Research Foundation under the Ministry of Science & ICT, Republic of Korea.


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Authors thank Mr. Joo Man Woo and Dr. Martin Halim for technical discussion during the preparation of this study. REFERENCES (1)

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