Hiroyuki Itoi*a,
Hiroyuki Muramatsu
b and
Michio Inagakic
aDepartment of Applied Chemistry, Aichi Institute of Technology, Yachigusa 1247, Yakusa-cho, Toyota, 470-0392 Japan. E-mail: itoi-hiroyuki@aitech.ac.jp
bFaculty of Engineering, Shinshu University, 4-17-1 Wakasato, Nagano, 380-8553 Japan
cProfessor Emeritus of Hokkaido University, 228-7399 Nakagawa, Hosoe-cho, Kita-ku, Hamamatsu, 431-1304 Japan
First published on 23rd July 2019
Nano-sized pores in carbon materials are recently known to give certain constraints to the encapsulated materials by keeping them inside, accompanied with some changes in their structure, morphology, stability, etc. Consequently, nano-sized pores endow the constrained materials with improved performances in comparison with those prepared by conventional processes. These pores may be called “constraint spaces” in carbon materials. Here, we review the experimental results related to these constraint spaces by classifying as nanochannels in carbon nanotubes, nanopores and nanochannels in various porous carbons, and the spaces created by carbon coating.
These pores and channels of carbon materials have played important roles in their applications, such as electrochemical capacitor electrodes by forming electric double-layers on a large surface area for energy storage, adsorbents of polluting species in air and water for environment remediation, etc. The interlayer spaces in graphite crystals are also used to accommodate active materials, the products being called graphite intercalation compounds, (e.g., by using lithium ions in rechargeable batteries). In these cases, the pores, channels, and interlayer spaces of graphite work as storage media of functional materials. Recently, however, some researches demonstrated experimentally that the pores and channels of carbon materials work to endow the encapsulated functional materials with performances either improved or different from those of the bulk or prepared by conventional processes (e.g., powder, films), as well as to enable the reactions of encapsulated materials, which are impossible outside the pores or channels. In other words, these pores and channels give certain constraints to the materials by keeping them inside, accompanied by some changes in their structure, morphology, stability, properties, etc., and therefore can be called “constraint spaces” in carbon materials. In the present review, we try to summarize the experimental results supposedly supported by these constraint spaces of carbon nanotubes and porous carbons. In addition to the pores and channels mentioned above, carbon coating processes of functional materials are included here, because some experimental results suggest certain constraints from the carbon walls to the coated materials.
The coalescence of C60 inside SWCNTs proceeds at high temperatures. C60@SWCNTs were prepared by reacting SWCNTs with C60 vapor at 400 °C in dry air, and then treated at high temperatures in vacuo (<10−6 Torr).14 Encapsulated C60 molecules are self-assembled to make a chain with nearly uniform center-to-center distances, as shown TEM image in Fig. 1a. C60 molecules remain mostly unchanged in CNTs up to 800 °C and start to coalesce with adjacent ones to form linked beans and/or short nanotube above 800 °C, as shown on 1000 oC-treated one in Fig. 1b. At ∼1200 °C, most C60 molecules coalesce to form tube wall, making this part DWCNTs by retaining the empty single-wall tube partly, as shown in Fig. 1c. After 1200 °C treatment, no C60 molecules remain and most of tubes are partly changed to double-walled, and in some cases, the inner tubes are terminated by caps (Fig. 1c). The thermal treatment of bucky-peapods is the most effective way for the fabrication of catalyst-free high-purity DWCNTs among the methods including arc-discharge and chemical vapor deposition (CVD) methods using catalysts, particularly to form small-diameter inner tubes (0.4–0.7 nm). The diameter of the inner tube synthesized from the peapods of SWCNTs with filling factor of 60–80% depends on the temperature of synthesis, as shown in Fig. 2.15 Almost the same diameter distribution of the inner tubes is achieved by the coalescence of fullerenes up to 1800 °C, but it becomes broad after the heat treatment at 2000 °C, including enlarged diameters due to coalescence of adjacent tubes.
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Fig. 1 TEM images of C60-encapsulated SWCNT: (a) as-prepared, (b) after heat treatment at 1000 °C, and (c) after 1200 °C in vacuo. Adapted from ref. 14 with permission from Elsevier. |
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Fig. 2 Distribution of the diameter of inner tube for DWCNTs synthesized at different temperature and high-resolution TEM images of as-prepared peapods and DWCNTs synthesized at 2000 °C. Adapted from ref. 15 with permission from John Wiley and Sons. |
Coalescence of fullerenes in SWCNTs is induced also by electron irradiation to form DWCNTs, which was followed by in situ TEM observations.16 DWCNTs are mechanically, thermally, and structurally more stable than SWCNTs, probably because of the buffer-like function of the outer tubes.17 The inner tube of DWCNT exhibits unique transport and optical properties such as extremely sharp Raman lines for the radial breathing mode (RBM) of the inner tube, suggesting high crystallinity of the inner tube and highly-unperturbed environment in the interior of the tubes.18 Highly crystalline and uniform TWCNTs were also obtained from the heat treatment of DWCNTs-peapods at 2000 °C in Ar.19 The DWCNTs were prepared by catalytically grown method, and their diameters were enlarged to the enough interior spaces of 1.2–1.6 nm for the encapsulation of C60 through the heat treatment at 2400 °C in Ar atmosphere. Such a high temperature triggers diameter enlargement process via the coalescence between DWCNTs.
Two adjacent DWCNTs (Fig. 3a) coalesce into off-centered coaxial carbon nanotubes (Fig. 3c) by annealing above 2100 °C.20,21 DWCNTs first merge into bi-cable carbon nanotubes by coalescence of the outer tubes (Fig. 3b), and the resulting inner tubes further coalesce inside a large diameter tube (Fig. 3c).21 Eventually, symmetric enlarged DWCNTs appear by more structural rearrangement of encapsulated off-centered inner tube along with the outer tube. The process of coalescence of two DWCNTs and merging into bi-cable CNT were discussed on the bases of detail TEM analyses and molecular dynamic calculations. Coalescence of SWCNTs to multi-walled CNTs was also reported.22
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Fig. 3 TEM images on the gradual changes from (a) adjacent DWCNTs, (b) bi-cable structure and (c) off-centered DWCNT during heat treatment at 2100 °C. Adapted from ref. 21 with permission from Elsevier. |
Encapsulation of Gd-metallofullerene cages into SWCNTs, which were previously heated at 420 °C in dry air for 20 min to open the tube ends, was carried out in a sealed glass ampoule at 500 °C for 24 h.23 The TEM images of fully encapsulated SWCNTs demonstrated that fullerene molecules were oriented randomly in respect to the tube axis, which could be supposed by observing dark spots due to Gd atoms in the fullerene cages.
K-doping into C60@SWCNTs was performed by exposing C60@SWCNTs to potassium vapor at 473 K over 50 h.24 TEM images before and after K-doping are shown in Fig. 4. After doping, some dark spots in the figures, which are reasonably supposed to be individual K atoms, are clearly observed inside CNTs and are allocated at the interfullerene sites. Electrochemical insertion/de-insertion of Li into C60@SWCNT was carried out in 1 M LiClO4/(EC + DEC) over the potential range of 0–3.0 V,25 and the cell gave the reversible capacity of 550–610 mA h g−1, which is slightly larger than the pristine SWCNT (460–490 mA h g−1), although C60@SWCNT and DWCNT showed a large irreversible capacity. The enhancement in the capacity is explained by the following two possibilities, the change in the electronic structure of the SWCNTs and a steric effect from the encapsulated C60. The encapsulated C60 stabilizes the Li ion desolvated at the entrance of the tubes.
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Fig. 4 TEM images of (a) the pristine C60@SWCNTs and (b–d) after K-doped ones. Adapted from ref. 24 with permission from the American Physical Society. |
Zn-diphenylporphyrin (Zn-DPP) is not a spherical but a flat molecule, of which size is about 1.0 nm and somewhat larger than C60 and C70, and possible to be encapsulated into SWCNTs with a diameter distribution from 1.25 to 1.47 nm at 400 °C.8 In the resultant Zn-DPP@SWCNTs, both Zn-DPP and SWCNT were supposed to be deformed from the experimental facts of a redshift of the π → π* transition in Zn-DPP and a drastic decrease of the radial breathing mode Raman spectrum intensity of SWCNTs.
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Fig. 5 G-band shift of SWCNTs as a function of applied potential. Adapted from ref. 27 with permission from the PCCP Owner Societies. |
Crystalline KI nanoparticles were grown in the nanospaces of single-wall carbon nanohorns (SWCNHs) by heating the mixture at 1073 K in a sealed quartz tube.29 The nanospaces of SWCNHs, internal tubular spaces and interparticle micropores, were measured as the volume of 0.50 and 0.11 cm3 g−1, respectively. The structure of KI crystals encapsulated in SWCNHs had the same as that formed under the pressure above 1.9 GPa in the bulk crystal, suggesting highly constrained state of KI crystals in the nanospaces of SWCNHs.
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Fig. 6 Sulfur encapsulation into carbon nanotubes: TEM images of (a) two straight S-chains in SWCNT, (b) zigzag and (c) straight S-chains in DWCNTs, and temperature dependences of interatomic distance d in SWCNT and DWCNT of (d) zigzag and (e) straight S-chains. Melting and boiling points for bulk sulfur (∼393 and ∼718 K, respectively) are indicated by vertical lines for comparison. Adapted from ref. 30 with permission from Nature Publishing Group. |
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Fig. 7 Relations of absorption intensity ratio Idoped/Ipristine to ionization energy and electron affinity of encapsulated molecules. Adapted from ref. 32 with permission from Nature Publishing Group. |
Encapsulation of 9,10-anthraquinone (AQ) and 9,10-phenanthrene quinone (PhQ) molecules into SWCNTs was performed by heating their mixture at 200 °C in an evacuated glass-tube.33,34 After the heat treatment, AQ@SWCNTs and PhQ@SWCNTs were recovered by washing with organic solvents (N,N-dimethylformamide for AQ and acetone for PhQ) to remove the excess AQ and PhQ deposited on the outer surface of the SWCNTs. Self-supported films of AQ@SWCNTs and PhQ@SWCNTs were obtained in this step. Electrochemical measurements for lithium-ion batteries (LIBs) and sodium-ion batteries (SIBs) were performed in 1 M LiClO4/(EC + DEC) and 1 M NaClO4/PC electrolytes, respectively. LIB performances of AQ@SWCNT and PhQ@SWCNT were compared with the mechanical mixtures of AQ and PhQ with acetylene black (Fig. 8). These results demonstrated the reversible Li-ion storage of almost all quinone (AQ and PhQ) molecules encapsulated in the SWCNTs. The capacity fading observed on the mechanical mixtures of quinones with acetylene black, which came from the dissolution of quinone molecules into the electrolyte, was markedly suppressed by the encapsulation.33 The effect of the tube diameter of SWCNT on the Li-ion storage was studied using two SWCNTs with different diameters of 1.5 and 2.5 nm (SWCNT-1.5 and -2.5, respectively), of which encapsulated PhQ amount was determined by thermogravimetric measurement to be 22 and 38 wt%, respectively.34 The charge–discharge profiles and the reversible capacities of PhQ@SWCNTs at room temperature were quite similar for Li and Na ions, associating with two steps at 2.8 and 2.4 V for Li and those at 2.3 and 1.9 V for Na. Because the tube diameters of both SWCNTs are larger than the Li and Na ions and do not discriminate between these two ions. The reversible capacity of LIB using PhQ@SWCNT-1.5 was determined from the charge–discharge curves to be very close to the theoretical capacity of the PhQ molecule (258 mA h g−1), but that using PhQ@SWCNT-2.5 was about a half of the theoretical one, suggesting that only a half of the encapsulated PhQ worked as electrode materials. This is explained by the strong interaction between SWCNTs and PhQ for narrower SWCNTs, forming a current flow path at the interface. The reversible LIB capacities of PhQ@SWCNT-1.5 and PhQ@SWCNT-2.5 at 0 °C were much smaller than those at room temperature probably due to kinetic problems, which was supported by the fact that the capacity measured at a very slow charging rate was as high as the theoretical one. Meanwhile, the capacities in SIBs at 0 °C for both electrodes slightly decreased in comparison with those at room temperature. Particularly, the reversible capacity of PhQ@SWCNT-2.5 at 0 °C was almost the same as that at room temperature, probably because of facile Na ion diffusion in SWCNT-2.5 even at a low temperature.
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Fig. 8 Charge–discharge curves collected at a current density of 100 mA g−1 at room temperature: mechanical mixture of (a) AQ with acetylene black, (b) PhQ with acetylene black; quinone-encapsulated SWCNTs of (c) AQ@SWCNT and (d) PhQ@SWCNT. Adapted from ref. 33 with permission from the PCCP Owner Societies. |
Encapsulation of β-carotene (C40H56) which consists of eight isoprene units having beta-rings at both ends of the molecule, was shown to be very efficient to suppress its degradation under UV irradiation.35 The light degradation of β-carotene is induced by the reaction with radical species (e.g., singlet oxygen) and isomerization; however, a surrounding tube wall prevents β-carotene from such reactions inside the constraint space. Encapsulation of β-carotene, 9,10-dichloroanthracene, and coronene into SWCNTs resulted in the enhancement of LIB capacity.36 Reversible capacity of coronene-encapsulated SWCNTs (7.2 wt%) was calculated to be 793 mA h per tube weight at 100 mA g−1 in 1 M LiClO4/(EC + DEC) electrolyte solution, whereas pristine SWCNTs exhibited reversible capacity of 316 mA h g−1. Although both electrodes showed a high irreversible capacity, but the enhancement in the reversible capacity observed for organic molecule-encapsulated SWCNTs is attributed to the increase in the Li ion storage sites in the tubes.
Encapsulation of metallocene compounds into SWCNTs was studied on ferrocene37–40 and cobaltocene.41 Encapsulation of ferrocene, FeCp2, was performed by contacting purified SWCNTs with ferrocene vapor at 300 °C under vacuum, followed by washing with diethyl ether to remove ferrocene deposited on the surface of tubes.37 The resultant FeCp2@SWCNTs retained the redox activity of FeCp2 itself in 0.1 M Bu4NClO4/acetonitrile electrolyte. FeCp2 confined in SWCNTs could be converted to metallic Fe nanoparticles by heating up to 700 °C under vacuum to form Fe-confined SWCNTs (Fe@SWCNT).38 Fe nanoparticles formed in SWCNTs have two morphologies, as shown TEM images with illustrations in Fig. 9: individual Fe particles with the diameter of 0.5–0.7 nm, suggesting several Fe atom aggregation (Fig. 9a) and somewhat linear alignment of Fe particles (Fig. 9b). According to the TEM observation, the filling yield (ratio of filled SWCNTs to empty SWCNTs) was estimated to be more than 80%. Encapsulation of FeCp2 into DWCNTs (FeCp2@DWCNT) resulted in marked change in electronic properties probably due to the charge transfer between FeCp2 molecules and DWCNTs, and Fe@DWCNT prepared from FeCp2@DWCNT was unipolar n-type semiconductive.39 FeCp2@SWCNT could change to DWCNT by high temperature annealing.40 By annealing at 1150 °C, all Fe atoms were released from the tube, while the annealing at 600 °C led to the formation of metastable iron carbide, which was supposed to act as a reactor by absorbing carbon atoms from one side of the tube and generating an inner tube to another side, forming DWCNT. Once the inner tubes had grown, the iron atoms diffused out of the tubes and aggregated into iron nanoparticles on the outside of DWCNTs, not in the channel. The intensity distributions of RBM Raman spectra for the inner-tube of FeCp2-derived DWCNTs were significantly different from those of C60-peapod-derived DWCNTs. Bis(cyclopentadienyl) cobalt (cobaltocene, CoCp2) could be confined only in the SWCNTs with a tube diameter of ∼1 nm under vacuum at 100 °C, whereas bis(ethylcyclopentadienyl) cobalt (Co(EtCp)2) confinement was possible in a range of tube diameters greater than ∼1 nm.41
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Fig. 9 TEM images and illustrations for two morphologies of Fe nanoparticles in Fe@SWCNT. Adapted from ref. 38 with permission. Copyright (2006) the Japan Society of Applied Physics. |
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Fig. 10 Encapsulated water in SWCNT: (a) XRD profiles at different temperatures, (b) temperature dependence of the intensity of the peaks at 22 and 7.2 nm−1, and (c) a schematic illustration of the ice nanotube inside a SWCNT. Adapted from ref. 45 with permission. Copyright 2002 the Physical Society of Japan. |
Loading and releasing processes of water molecules into SWCNTs were studied by X-ray diffraction analysis and hybrid reverse Monte Carlo simulations.47 Water loading was performed at vapor pressures of 3.1 and 3.8 kPa, corresponding to water filling rates of 50 and 100%, respectively, and then water was released up to a vapor pressure of 2.3 kPa (corresponding to water filling of 50%). In the loading process, water molecules formed nanoclusters which were well stabilized in the channel of SWCNTs, while less stable water layers than water nanoclusters were formed in the releasing process.
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Fig. 11 Electrochemical capacitor behaviors of PANI-confined ACs: (a) cyclic voltammograms collected at 1 mV s−1, (b) gravimetric and (c) volumetric capacitance dependences on the current density. Adapted from ref. 48 with permission from the Royal Society of Chemistry. |
Polypyrene (PPY) and PANI were confined into not only micropores but also mesopores of a commercial AC (MSC30), which had a large SBET of 3160 m2 g−1, micropore volume (Vmicro) of 0.99 cm3 g−1, and mesopore volume (Vmeso) of 0.60 cm3 g−1. The vapor of pyrene (PY) and ANI were adsorbed in the AC at 150 and 25 °C, respectively, followed by electrochemical polymerization in an aqueous 1 M H2SO4.49 The adsorption saturations of PY and ANI occurred at 60.2 and 62.4 wt%, respectively, for this AC. The capacitive performances are shown on the ACs with different amounts of PPY in Fig. 12. Anodic and cathodic peaks at 0.3–0.4 V due to the redox reaction of PPY confined into the AC pores are clearly observed in their cyclic voltammograms (Fig. 12a), as observed on the PANI-confined ACs (Fig. 11a), which gave the pseudocapacitance to the capacitors. By the confinement of 30–50 wt% PPY, gravimetric and volumetric capacitances of the cells are markedly enhanced, as shown in Fig. 12c and d. The maximum volumetric capacitance for the PPY-confined ACs and the PANI-confined ACs reached 314 and 299 F cm−3, respectively. Although the increment in the observed volumetric capacitances are reasonably supposed to be due to the pseudocapacitance of PPY, the rate performances of the cells are also improved, as shown in Fig. 12b, suggesting that the redox reaction of PPY occur more rapidly than the double-layer formation inside the nanopores. These results revealed that a large contact area between conductive polymers and carbon surfaces enables compatibility of high volumetric energy and high power densities for electrochemical capacitor electrodes.
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Fig. 12 Electrochemical capacitor behaviors of PPY-confined ACs: (a) cyclic voltammograms with 1 mV s−1 rate, (b) dependences of gravimetric capacitance and coulombic efficiency on the cycle number, (c) rate performances of the gravimetric capacitance, and (d) those of the volumetric capacitance. Adapted from ref. 49 with permission from the Royal Society of Chemistry. |
2,5-Dichloro-1,4-benzoquinone (DCBQ) was confined into the pores of Ketjenblack (KB; SBET: 1340 m2 g−1, Vmicro: 0.48 cm3 g−1, Vmeso: 1.24 cm3 g−1) at 100 °C (a little higher than the sublimation temperature of DCBQ, 92 °C).50 The weight content of DCBQ in the DCBQ-confined KB was tuned to be 5, 10, 20, and 40 wt%. In addition, KB was saturated using excess amount of DCBQ, with the saturation amount reaching 60.1 wt%. N2 adsorption–desorption curves shown in Fig. 13a suggest that DCBQ was adsorbed into both micropores and mesopores, Vmicro and Vmeso being reduced to 0.17 and 0.71 cm3 g−1, respectively, with a 40 wt% DCBQ content; and giving Vmicro of 0.03 cm3 g−1 and Vmeso of 0.15 cm3 g−1 upon saturation. The DCBQ-confined KBs exhibit clearly the presence of redox reaction of DCBQ by a plateau in galvanostatic charge–discharge curves in an aqueous 1 M H2SO4 electrolyte, as shown in Fig. 13b, although the pristine KB shows almost linear characteristics (i.e., electric double-layer behavior). In Fig. 13c and d, gravimetric and volumetric capacitances are plotted against current density, revealing that excellent rate performance of the pristine KB is retained even in DCBQ-confined KBs. DCBQ has poor electrical conductivity by itself, like most organic compounds, but a large contact area between finely dispersed DCBQ in the pores of KB and conductive carbon surfaces enables fast charge transfer at the interface. Gravimetric capacitance reaches a maximum with a DCBQ content of 40 wt%, but the KB saturated with DCBQ gives a maximum volumetric capacitance. The resulting capacitor was characterized by high capacitances, particularly high volumetric capacitances as 4.7 times higher than that of the pristine KB, along with high rate capability up to 5 A g−1 and excellent cycle lifetimes up to 10000 cycles.
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Fig. 13 DBCQ-confined Ketjenblack: (a) N2 adsorption–desorption curves at 77 K, (b) galvanostatic charge–discharge curves in 1 M H2SO4 electrolyte measured at a current density of 50 mA g−1, (c) rate performances of the gravimetric capacitance, and (d) those of the volumetric capacitance. Adapted with permission from ref. 50. Copyright 2016 American Chemical Society. |
One of 2,2,6,6-tetramethylpiperidine-N-oxyl (TEMPO) derivatives, 4-hydroxy-TEMPO benzoate (HTB), was confined into the nanopores of the AC (MSC30; SBET: 3160 m2 g−1, Vmicro: 0.99 cm3 g−1, Vmeso: 0.60 cm3 g−1) through its vapor at 130 °C, and the capacitive performances of the resultant HTB-confined ACs with different HTB contents (20–50 wt%) were studied in an aqueous 1 M H2SO4 electrolyte.51 A marked reduction in pore parameters with increasing adsorbed amount of HTB suggested the confinement of HTB molecules into nanopores of AC, with the SBET, Vmicro, and Vmeso reducing to 250 m2 g−1, 0.07 cm3 g−1, and 0.05 cm3 g−1, respectively, by adsorption of HTB up to 50 wt%. Anodic and cathodic peaks derived from the redox reaction of HTB were clearly observed at around 0.7 V. The redox potential of HTB is higher than those of many organic redox compounds and conductive polymers, which is advantageous to enhance the energy density of the HTB-confined ACs. The composite with 30 wt% HTB exhibited higher volumetric energy and power densities than the pristine AC: 39.5 W h kg−1 at 16.5 W kg−1 and 14.1 W h kg−1 at 8.2 kW kg−1 for the composite in contrast to 18.0 W h kg−1 at 16.5 W kg−1 and 1.1 W h kg−1 at 8.2 kW kg−1 for the pristine AC. Volumetric comparison shows a further enhancement in the volumetric energy and power densities of the composite with 30 wt% HTB: a 3.1 times enhancement (16.5 W h L−1) at 6.9 W L−1 and a 18 time enhancement (5.9 W h L−1) at 3.5 kW L−1. The high redox potential of HTB and a large contact area between the conductive carbon surface and HTB realized compatibility between the enhancements in both volumetric energy and power densities. The confinement of HTB molecules into the nanopores of AC was effective to prevent dissolution of HTB in an aqueous electrolyte with the aid of a hydrophobic group in HTB and to improve markedly the contact of HTB with a conductive carbon surface, which was supposed to allow for fast redox reactions of HTB, in addition to fine dispersion of minute aggregates of HTB molecules.
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Fig. 14 TEM images of MnO2 nanocrystals confined in the pore (c and d) and deposited on the surface (b) of the porous carbon nanofiber (a). Adapted from ref. 52 with permission from Nature Publishing Group. |
The electrochemical cell consisting of the same KOH-activated porous carbon (SBET of 2181 m2 g−1) worked as a hybrid cell, battery-type behavior at the positive electrode and capacitor-type one at the negative electrode in an aqueous electrolyte containing 2 M MnSO4 and 0.5 M KI, although it worked as a capacitor in 2 M MnSO4 electrolyte.53 This hybrid behavior in alkali metal iodide-based aqueous electrolytes was concluded to be owing to the confinement of polyiodides in the porosity of the positive carbon electrode creating a so-called carbon/iodide interface,54 the confinement of polyiodides being confirmed by the measurement of Raman spectra.53 Effectiveness of the confinement of electrochemically-active materials was reported on Sn55,56 and Ge.57,58
BaTiO3 rod-like crystals were synthesized in the nanopores of single-wall carbon nanohorns (SWCNHs).59 TiO2 was firstly encapsulated into the nanopores of SWCNHs by impregnating an aqueous TiCl4 solution and following annealing at 373 K. The resulting TiO2 was reacted with Ba(OCH2CH3)2 at 400 K in an autoclave of N2 atmosphere, and the resultant BaTiO3 rods having high aspect ratio (2–4 nm in diameter and 40–60 nm length) were obtained. By impregnation of SWCNHs in methanol/2-methoxyethanol mixed solution of Ba(OCH2CH3)2 and Ti(OCH(CH3)2)4 and subsequent annealing at 400 K in an autoclave, spherical BaTiO3 particles with the diameters of 2–4 nm were synthesized in the nanopores of SWCNHs, whereas the BaTiO3 particles with the diameters of 20–40 nm were synthesized via the same process without using SWCNHs.
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Fig. 15 TEM images of Pt-loaded carbon materials: (a) Ketjenblack and (b and c) ZTC. Adapted with permission from ref. 60. Copyright 2017 American Chemical Society. |
Pt nanoparticles were loaded on an Ar+-ion irradiated glassy carbon plate by radio-frequency magnetron sputtering with plasma output of 20 W for 60 s at fluences up to 1.0 × 1016 ions per cm2.61 Pt was supposed to be confined at the defects created by the irradiation of 380 keV Ar+ as minute clusters with an average size of about 5 nm. A strong interfacial interaction between Pt nanoparticles and the glassy carbon substrate was estimated from its Pt 4f7/2 and C 1s XPS spectra.
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Fig. 16 g-C3N4/CMK-3 composite: (a) TEM image, (b) SEM image, and (c) N2 adsorption/desorption isotherm. Adapted with permission from ref. 63. Copyright 2011 American Chemical Society. |
In lithium–sulfur (Li–S) batteries, anchoring lithium polysulfides (Li2Sx) in the cathode materials is strongly demanded to inhibit their shuttling between electrodes during charge/discharge cycles, which would lead to the dissolution of Li2Sx (4 ≤ x ≤ 8) into commonly used electrolytes, resulting in the short cycle lifetimes of the electrodes and lithium anode contamination. Codoping of N and S into mesoporous carbons was demonstrated to be effective to enhance the affinity for Li2Sx, preventing their dissolution into an electrolyte.64 The composite of conductive vanadium nitride (VN) with reduced graphene oxide (rGO) was also reported to deliver good rate and cycling performances, with the initial capacity being 1471 mA h g−1 with only 15% capacity reduction after 100 cycles at 0.2C rate.65 The conductive and porous structure of the VN/rGO composite facilitates the electron transportation and Li ion diffusion. In addition, since the VN has not only high electrical conductivity but also catalytic activity, it shows an anchor effect on the polysulfides and accelerates their redox reactions. g-C3N4 nanodots with the sizes less than 5 nm were confined into a MOF-derived N,S-codoped hollow porous carbon shell, and the resulting composite demonstrated excellent performances in Li–S batteries as cathode.66 As shown in Fig. 17, the composite exhibits much higher capacity and much better rate and cycle performances than neat g-C3N4 and the pristine porous carbon. The composite delivered 1447 mA h g−1 at 0.2C rate and 387 mA h g−1 at 5C, with excellent cycling stability as only 0.048% capacity reduction at 1.0C after 500 cycles. The synergetic effect of hollow carbon structure, N,S-codoping, and encapsulation of g-C3N4 facilitates the electron transfer and endows the composite with strong adsorption affinity for polysulfides to suppress the polysulfide shuttling.
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Fig. 17 Li–S battery performances of g-C3N4-confined porous carbon in comparison with g-C3N4 and pristine porous carbon: (a) rate performances and (b) cycle performances at 1C rate. Adapted from ref. 66 with permission from the Royal Society of Chemistry. |
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Fig. 18 Methane adsorption/desorption isotherms of the activated carbon having different pre-absorbed water content Rw. Adapted from ref. 67 with permission from Nature Publishing Group. |
Particle sizes of Si as anode materials exhibited marked influences on rate and cycle performances of LIBs with 1 M LiPF6/(EC + DEC).69 Si particles with the particle size of ca. 1 μm (micron-sized Si) showed abrupt decrease in capacity with the cycle number even at a small current density of 0.2 A g−1 due to the pulverization of Si particles. The poor cycle lifetimes of the micron-sized Si particles were attributed to the typical bulk like behavior. On the other hand, nano-sized Si particles with the size of 82 nm exhibited the remarkably enhanced coulombic efficiency at the 1st and following cycles. The capacity retention of the nano-sized Si particles was enhanced by carbon coating with a thickness of ca. 10 nm through a pressure-pulsed chemical vapor deposition (P-CVD) method; the carbon coating decreased the inner resistance. The particle morphology of the nano-sized Si particles and the carbon-coated ones changed from spherical particles to nano-sized wrinkled structure after 20 charge/discharge (lithiation/de-lithiation) cycles (Fig. 19a and b).69,70 At the same time, the coated carbon was also deformed together with wrinkled Si, as evidenced by scanning transmission electron microscopy (STEM) imaging and energy dispersive X-ray spectroscopy (EDS) analysis.69 After 100 cycles, the wrinkled structure of both samples further transformed into an aggregated lump due to the repeated volumetric expansion and the subsequent agglomeration of pulverized Si particles (Fig. 19c).70 The capacity fading accompanied by the structural deformation of the wrinkled Si was suppressed by limiting the capacity up to 1500 mA h g−1 during lithiation because of restricting the volumetric expansion to some extent. Similar structural transformation through wrinkled one and improved battery performances were observed on ball-milled Si, which is more inexpensive than Si nanoparticles.70 On the ball-milled Si consisting of weakly aggregated primary particles with the size of about 47 nm, the initial reversible capacity was improved reaching ca. 3000 mA h g−1 by carbon coating, as shown in Fig. 20a. By restricting the discharge capacity up to 1500 mA h g−1, cycling with different current densities of 0.2–5.0 A g−1 was possible on the carbon-coated Si up to at least 100 cycles, whereas ball-milled Si without carbon coating showed a marked capacity fading (Fig. 20b). Flaky Si nanoparticles (thickness of ca. 16 nm and lateral size of 0.2–1 μm) were produced from Si sawdust by beads-milling in isopropyl alcohol. Since the Si sawdust contained 4 wt% graphite, which came from the graphite substrate used for the cutting process of Si ingots, the graphite was homogeneously dispersed over the Si particles after a beads-milling process.71 Although their structure was converted into nanoflake morphology through a milling process, they could be transformed to wrinkled morphology by cycling of lithiation/de-lithiation and demonstrated better performances as an anode of LIBs. The CVD of the nanoflake Si was also examined, and the resulting material exhibited the enhanced rate capability at 5 A g−1 under the limiting capacity of 1500 mA h g−1 during lithiation, reaching almost the same value as that at 0.2 A g−1. The enhanced rate capability was attributed to the large continuous structure formed by carbon coating. The carbon layer on Si particles is supposed to work as a buffer between neighboring Si nanoparticles even during marked morphology change from spherical or flaky to wrinkled, in other words, by confining Si nanoparticles within coating carbon.
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Fig. 19 TEM images and morphology illustrations of nano-sized Si after cycling of lithiation/de-lithiation. Adapted from ref. 70 with permission from Elsevier. |
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Fig. 20 Capacity changes with cycling with different current densities on pristine ball-milled Si and carbon-coated one: (a) without and (b) with the capacity restriction to 1.5 A h g−1. Adapted from ref. 70 with permission from Elsevier. |
Thin Si flakes were synthesized by CVD of silane (SiH4) gas on the cubic crystals of NaCl at 550 °C, followed by dissolution of NaCl. The thickness and lateral size of the Si flakes was controlled to be 50 nm and a micrometer, respectively.72 Carbon coating of the Si flakes was performed by heating in acetylene gas at 900 °C, resulting in the carbon content of 7 wt% and the thickness of about 10 nm. The non-porous structure of the pristine and carbon-coated Si flakes with a low surface area of ca. 10 m2 g−1 suppressed excessive SEI formation and reduced initial irreversible capacity (Fig. 21a), in comparison with that of the commercially available Si nanoparticles. Their initial reversible capacities reached 2943 and 2255 mA h g−1, corresponding to the initial coulombic efficiency of 87.2% and 92.3%, respectively in 1.3 M LiPF6/(EC + DEC). As shown in Fig. 21b, the carbon-coated Si flakes delivered a steady rate performance around 2000 mA h g−1 up to 10C rate, whereas the pristine Si flakes show the inferior rate capability. The full cell coupled with a LiCoO2 cathode shows a steady cycle performance at 0.5C rate for the carbon-coated Si flakes, while the reversible capacity of the pristine Si flakes decayed drastically (Fig. 21c). The carbon layer on Si flakes works not only to improve the electrical conductivity in the anode but also to buffer the large expansion of Si during lithiation.
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Fig. 21 Carbon-coated Si flakes in comparison with the pristine Si flakes: (a) the first charge/discharge curves, (b) rate capability at different C-rates and (c) cycle performance of the cell coupled with a LICoO2 cathode at 0.5C rate. Adapted from ref. 72 with permission from Nature Publishing Group. |
The structural deformation of the carbon-coated Si flake was monitored by in situ TEM under a bias of −3 V and 3 V, respectively, where a flake was mounted on a Pt wire connected to the Li2O/Li probe; the Li metal and the LiO2 layer serve as the counter electrode and the solid electrolyte, respectively.72 The structural deformation of the flake in the first two lithiation/de-lithiation cycles was schematically illustrated in Fig. 22. During the first lithiation, the flake showed anisotropic swelling, with only 9.7% expansion along lateral direction but about 200% in thickness upon full lithiation (Li15Si4), giving ∼270% volumetric swelling. By de-lithiation, the flake rippled along the lateral direction of the flake and the flake itself was significantly buckled. The second lithiation smoothed out the ripples and the de-lithiation gave severe buckling to the flake, while maintaining its 2D morphology without fracture. In contrast, the pristine Si flake without carbon coating swelled into the lateral direction as much as 31.9% after the first lithiation and could not recover its original dimensions during de-lithiation, preserving the swollen morphology and sharp edges generated during lithiation. The fact that lithiation of the pristine Si took longer times under higher bias than those of the carbon-coated Si flake indicates the relatively poor rate capability. The distinct difference in morphology change during the lithiation/de-lithiation cycle between the pristine and carbon-coated flakes is reasonably supposed to be due to the carbon layer coating, suggesting certain mechanical constraint of the carbon layers on the deformation of Si flakes.
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Fig. 22 Schematic illustration of morphology change during lithiation/de-lithiation process. Adapted from ref. 72 with permission from Nature Publishing Group. |
B-doped Si/SiO2 composites were prepared from the mixture of ball-milled SiO2 and B2O3 powder (20:
1 by mole) at 950 °C in Ar, where B-doped crystalline Si nanoparticles with the size of about 15 nm were dispersed in amorphous SiO2 particles (particles size of about 3 μm).73 Carbon coating of the composite by CVD of acetylene at 700 °C resulted in high cyclic performance of a hybrid cell coupled with mesoporous carbon spheres as a cathode using a 1 M LiPF6/(EC + DEC + DMC) electrolyte (lithium-ion capacitor, LIC), which delivered an energy density of 128 W h kg−1 at a power density of 1.23 kW kg−1 and even 89 W h kg−1 at 9.7 kW kg−1 with the voltage window of 2.0–4.5 V. Amorphous SiO2 matrix of the composites was supposed to work as a buffer layer for a large volume change of Si during lithiation/de-lithiation. In addition, boron-doping into Si increases the electrical conductivity of Si and enables fast charge transfer, necessary for high power density. Furthermore, the electrical conductivity and cycle lifetimes of Si were markedly improved by carbon coating. Si-included carbon composites were prepared from wasted glass (Si source) and wasted polyvinyl butyral (carbon source), both of which were derived from windshields with a polyvinyl butyral layer sandwiched by two glasses.74 The wasted glass separated from the polyvinyl butyral was reduced to Si by Mg powder as a reducing reagent at 650 °C for 6 h under Ar atmosphere, and the byproducts (Mg2Si and Mg2SiO4) and the unreacted SiO2 were removed by HNO3 and HF etching, respectively. The carbon coating was performed by mixing the obtained Si nanoparticles and the polyvinyl butyral and their subsequent heat treatment at 500 °C for 4 h. Its LIB performance was measured in 1 M LiPF6/(EC + PC) electrolyte solution. The reversible capacity of the non-carbon-coated Si in the first charge/discharge cycle reached 3336 mA h g−1 with irreversible capacity of 489 mA h g−1, but it decayed rapidly with the number of cycling measured at 420 mA g−1, down to less than 1000 mA h g−1 after 50 cycles. In contrast, the carbon-coated Si with a carbon content of 60 wt% showed the initial reversible capacity of 856 mA h g−1, while its capacity retention reached 84% at the 300th cycle. Although its initial reversible capacity was less than 1000 mA h g−1, the capacity per unit mass of Si is calculated to be over 2000 mA h g−1. The carbon matrix stabilized the Si nanoparticles and enhanced electrical conductivity of the electrode, which was supported by the result of the impedance analysis. The composite of Si/rGO was prepared from a colloidal mixture of wasted Si sludge containing SiC from the wafer-slicing saw with GO by aerosol spray pyrolysis with thermal reduction of GO.75 The colloidal mixture was sprayed into a tubular furnace by ultrasonic nebulizer, and at the same time, unnecessary SiC was separated from the droplet because the Si particles have a lighter density and smaller sizes than the SiC particles. The as-prepared composite has ball-like morphology and the Si particles are located inside the rGO shell. The composite prepared from the colloidal mixture with 0.1 wt% GO delivered the reversible capacity of about 1600 mA h g−1 in the initial charge/discharge cycles by retaining the capacity of about 1000 mA h g−1 after 50 cycles. The composite prepared from the mixture with 0.4 wt% GO gave much better cyclability although the capacity became smaller (ca. 700 mA h g−1).
Carbon-coated Si nanowires were synthesized from the mixture of bis(bis(trimethylsilyl)amino)tin, Sn(HMDS)2, and monophenylsilane (MPS) with different mole ratios (Sn/Si of 1/16 to 1/64).76 The reaction proceeded under supercritical fluid condition in anhydrous toluene solution at 490 °C and 10.3 MPa, followed by the carbonization at 900 °C under N2/H2 gas flow. At first, MPS decomposed to atomic Si and high order phenylsilane through disproportionation reaction, while Sn(HMDS)2 decomposed to Sn nanoparticles which worked as a seed to form eutectic with Si for growing Si nanowires. The high order phenyl silane byproduct polymerized on the surface of the Si nanowires and the resulting polyphenylsilane was converted to carbon shells during the heat treatment at 900 °C. The morphology of the carbon-coated Si nanowires highly depends on the Sn/Si molar ratio, and the nanowires prepared below a Sn/Si molar ratio of 1/32 before the heat treatment were heavily kinked and covered by thick polyphenylsilane shell. The carbon-coated Si nanowires prepared with a Sn/Si molar ratio of 1/22 showed the markedly straighter and thinner shell, and the carbon shell retained the uniform thickness after the heat treatment (Fig. 23). This composite exhibited a high capacity over 2000 mA h g−1, with nearly 96% coulombic efficiency when cycled at a slow rate of 0.1C for 100 cycles. Moreover, a stable capacity of about 1300 mA h g−1 with more than 98% coulombic efficiency was attained after 100 cycles when measured at 1C. In contrast, the nanowires without the heat treatment exhibited no capacity after 20 cycles at 1C rate, although the capacity at the first cycles was about 300 mA h g−1. The in situ TEM observation of the lithiation of the carbon-coated Si nanowires indicated that uniform carbon coating on the Si nanowires prevented full lithiation while nonuniform and incomplete carbon coating allowed the Si nanowire to be fully lithiated, relieving the stress created by expanding Si nanowire.
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Fig. 23 Carbon-coated Si nanowires prepared from the mixture of Sn(HMDS)2 and MPS in Sn/Si mole ratio of 1/22 after carbonization at 900 °C: (a) SEM, (b) TEM and (c) HRTEM images. Adapted with permission from ref. 76. Copyright 2013 American Chemical Society. |
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Fig. 24 Results of galvanostatic charge/discharge analysis of carbon-coated LVP (anode)//AC (cell I) and AC//carbon-coated LVP (cathode) (cell II) in 1 M LiPF6/(EC + DEC) electrolyte: (a and d) charge/discharge curve at 100 mA g−1 rate. (b and e) Cyclic performance at 100 mA g−1 rate, and (c and f) rate performance (1C = 131 mA h g−1). Adapted from ref. 77 with permission from Elsevier. |
Carbon-coated Li2MnSiO4 was synthesized by solid-state reaction of the mixture of LiOH, MnCO3 and SiO2 with adipic acid at 900 °C. Its LIC cell coupled with AC anode in 1 M LiPF6/(EC + DMC) delivered the capacitance of 43.2 F g−1 at 1 mA cm−2 and the energy density of 54 W h kg−1 at 0.15 kW kg−1 with the retention of about 85% after 1000 cycles.80 Carbon-coated Li2FeSiO4 was also synthesized as anode materials for LICs by the same procedure and exhibited almost the same LIC performances as the carbon-coated Li2MnSiO4.81 Adipic acid used in these studies is the carbon source to improve the electrical conductivity between the particles of active materials by carbon coating, thus achieving high-power density cathodes for hybrid super capacitors. Carbon coating of H2Ti12O25, which was prepared from Na2CO3 and TiO2 at a molar ratio of 1:
3, was performed by mixing with beta cyclodextrin (4.5 wt%) as a carbon source and subsequent calcination at 800 °C for 2 h. The thickness of the carbon layer was determined to be 3.09 nm by TEM observation, and the carbon coating was effective to improve LIC performance, giving negligible capacitance decrease after 100 cycles and energy density of 38.8 W h kg−1 at a power density of 0.18 kW kg−1 in 1.5 M LiBF4 (EC + DMC).82 The carbon layer was supposed to not only enhance the electrical conductivity but also suppress the swollen phenomenon, which was caused by the reduction decomposition between H2Ti12O25 and the electrolyte solution, generating C2H4, CO, and CH4 gases.
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Fig. 25 Li–S batteries composed of carbon-coated rGO flakes: (a) charge/discharge curves and (b) rate performances for the cathodes with different S-contents. Adapted from ref. 83 with permission from the Royal Society of Chemistry. |
Carbon nanotubes provide us well-defined spaces with various controllable parameters for matrix carbon, including their diameter and length, layer number of the wall, etc., in addition to the possibility for the nanotubes after confinement of the material to be directly observed by microscopic techniques. However, it has to be pointed out that the nanotube wall is limited to an almost perfect but highly-curved layer of carbon hexagons (we may say “graphene” layer, but strongly curved), in addition to the high cost and limited amount of carbon nanotubes. Zeolite-templated carbons are composed of ordered micropores with walls of defective graphene-like layers, but they are not commercialized yet and their synthesis procedure is complicated. Also the carbons prepared by using either mesoporous silica or block copolymer surfactants as the templates provide us ordered channels with mesopore-sized diameters, but the walls of these mesoporous carbons consist of amorphous carbon because carbon precursors for these templates are almost limited to phenolic resins. It is not certain whether only carbon coating of the materials can create constraint space or not, however, carbon coating techniques give us variable carbon walls on the target materials; dense carbon films with controlled thickness can be deposited on the materials by CVD, and porous carbon films can form the mixture of particles of the materials with various resins (carbon precursors), such as thermosetting and thermoplastic resins. Since carbon materials are endowed with the structural and textural diversities, as pointed out above, there are infinite combinations for synthesizing the composites using versatile porous carbons. Moreover, the electrical conductivity of carbon materials can widen their applications toward electronic devices and electrocatalysts.
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