N-Doping and Defective Nanographitic Domain Coupled Hard Carbon Nanoshells for High Performance Lithium/Sodium Storage

Abstract

Hard carbons (HCs) possess high lithium/sodium storage capacities, which however suffer from low electric conductivity and poor ion diffusion kinetics. An efficient structure design with appropriate heteroatoms doping and optimized graphitic/defective degree is highly desired to tackle these problems. This work reports a new design of N-doped HC nanoshells (N-GCNs) with homogeneous defective nanographite domains, fabricated through the prechelation between Ni2+ and chitosan and subsequent catalyst confined graphitization. The as-prepared N-GCNs deliver a high reversible lithium storage capacity of 1253 mA h g−1, with outstanding rate performance (175 mA h g−1 at a high rate of 20 A g−1) and good cycling stability, which outperforms most state-of-the-art HCs. Meanwhile, a high reversible sodium storage capacity of 325 mA h g−1 is also obtained, which stabilizes at 174 mA h g−1 after 200 cycles. Density functional theory calculations are performed to uncover the coupling effect between heteroatom-doping and the defective nanographitic domains down to the atomic scale. The in situ Raman analysis reveals the “adsorption mechanism” for sodium storage and the “adsorption–intercalation mechanism” for lithium storage of N-GCNs.

Full text

FULL PAPER
www.afm-journal.de
© 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
1706294 (1 of 10)
N-Doping and Defective Nanographitic Domain Coupled
Hard Carbon Nanoshells for High Performance Lithium/
Sodium Storage
Shifei Huang, Zhiping Li, Bo Wang, Jiujun Zhang, Zhangquan Peng, Ruijuan Qi,
Jing Wang, and Yufeng Zhao*
Hard carbons (HCs) possess high lithium/sodium storage capacities, which
however suffer from low electric conductivity and poor ion diffusion kinetics.
An efficient structure design with appropriate heteroatoms doping and opti-
mized graphitic/defective degree is highly desired to tackle these problems.
This work reports a new design of N-doped HC nanoshells (N-GCNs) with
homogeneous defective nanographite domains, fabricated through the pre-
chelation between Ni2+ and chitosan and subsequent catalyst confined gra-
phitization. The as-prepared N-GCNs deliver a high reversible lithium storage
capacity of 1253 mA h g−1, with outstanding rate performance (175 mA h g−1
at a high rate of 20 A g−1) and good cycling stability, which outperforms most
state-of-the-art HCs. Meanwhile, a high reversible sodium storage capacity
of 325 mA h g−1 is also obtained, which stabilizes at 174 mA h g−1 after
200 cycles. Density functional theory calculations are performed to uncover
the coupling effect between heteroatom-doping and the defective nanogra-
phitic domains down to the atomic scale. The in situ Raman analysis reveals
the “adsorption mechanism” for sodium storage and the “adsorption–
intercalation mechanism” for lithium storage of N-GCNs.
DOI: 10.1002/adfm.201706294
Dr. S. Huang, Prof. Z. Li, B. Wang, Prof. J. Wang, Prof. Y. Zhao
Key Laboratory of Applied Chemistry
Yanshan University
Qinhuangdao 066004, China
E-mail: yufengzhao@ysu.edu.cn
Prof. J. Zhang, Prof. Y. Zhao
College of Sciences
Shanghai University
Shanghai 200444, China
Prof. Z. Peng
State Key Laboratory of Electroanalytical Chemistry
Changchun Institute of Applied Chemistry Chinese Academy of Science
Changchun, Jilin 130022, China
Prof. R. Qi
Key Laboratory of Polar Materials and Devices
Ministry of Education
East China Normal University
Shanghai 200241, China
demanded for potential applications
in portable electronics, electrical vehi-
cles, and smart grids, etc.[1,2] Currently,
graphite is widely applied as the anode
material for commercial lithium-ion bat-
teries (LIBs), which however suffers
from limited energy and power perfor-
mance due to its low theoretical capacity
(372 mA h g−1) and sluggish lithium dif-
fusion kinetics.[3–5] On the other hand,
graphite is inapplicable to sodium-ion bat-
teries (SIBs) due to the large ionic radius
of Na+ ions (1.02 vs 0.76 Å),[6,7] and the
sodium–graphite system is thermodynam-
ically instable.[6] Therefore, exploring new
anode materials with favorable voltage,
low cost, desired capacity, and high rate
of charge for LIBs/SIBs is of great impor-
tance for practical applications.
In recent years, hard carbons (HCs)
have been emerging as one of the most
promising anodes for LIBs/SIBs,[8,9] which
can be divided into polymer derived HCs
and biomass derived HCs according to different types of the
carbon precursors.[6,7,10–14] These HCs are composed of plen-
tiful randomly packed single graphene layers as described by
the “house of cards” model,[15] which can provide numerous
nanopores and large surface area, thus offering significantly
enhanced capacity by facilitating both intercalation and adsorp-
tion of Li+ or Na+ ions.[16] Nevertheless, the low graphitization
degree of HCs normally cause low electric conductivity and
electrochemical instability.[1,17,18] Attempts have been made by
compositing HCs with graphitic carbon, which however cannot
overcome the disadvantage of the intrinsically disordered micro-
architecture.[18–21] Creating nanographitic domains in HCs
through catalyst assisted synthesis, demonstrates an efficient
way to improve the electric conductivity.[1,22–24] Such unique
structure possesses plenty of nanopore defects in the graphitic
domains, thereby enable stable charge storage through both
interlayer intercalation (nanographite layers) and interfacial
adsorption (nanopore defects). On the other hand, doping het-
eroatoms such as B, N, S, P, etc. to the carbon architecture, can
introduce more active sites and facilitate the charge transfer,
thus significantly enhance the Li/Na storage capability.[25–31]
Therefore, the synergistic effect between heteroatom-doping
and nanopore defects in the nanographitic domains are obliged
Lithium/Sodium Storage
1. Introduction
Electrochemical energy storage technologies based on recharge-
able batteries (e.g., lithium or sodium-ion batteries) with high
storage capacity and excellent rate performance are critically
Adv. Funct. Mater. 2018, 1706294

www.afm-journal.dewww.advancedsciencenews.com
1706294 (2 of 10) © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
to be considered in further research efforts. An effective design
of hard carbons with optimized nanoarchitecture and corre-
sponding theoretical study are critically demanded for high per-
formance lithium/sodium storage.
In this work, we report a new design of N-doped hard carbon
nanoshells with tailored graphitic/defective degree (N-GCNs),
whereby Ni2+ ions are prechelated with chitosan carbon
source to enable a catalyst confined graphitization process and
then assure the homogeneous distribution of nanographitic
domains. As expected, the as-prepared N-GCNs perform out-
standing charge storage capacities and rate performances for
both LIBs and SIBs. The density functional theory (DFT) cal-
culations evidence the synergistic effect between N-doping and
nanopore defects in the nanographitic domains for Li storage,
which demonstrates that proper amount of defects and het-
eroatom doping will improve the Li binding ability, charge
transfer ability, and increase the density of states (DOS) around
Fermi level. In situ Raman analysis unravels the different
charge storage mechanism of N-GCNs for LIBs (adsorption–
intercalation mechanism) and SIBs (adsorption mechanism).
2. Results and Discussion
2.1. Characterization
The N-GCNs were fabricated through a one-pot catalyst con-
fined graphitization process as illustrated in Figure 1a. Typically,
Artemia cyst shell was used as carbon precursor and soaked
with aqueous solution containing Ni2+ ions and KOH, whereby
the Ni2+ is chelated with chitosan to make sure the uniform
distribution of Ni2+ ions in the carbon precursor.[32,33] The
KOH reacts immediately with Ni2+, forming homogeneously
dispersed Ni(OH)2 nanoparticles (NPs) (decomposition yields
NiO at 230 °C). The carbon precursors then go through a
pyrolysis-carbonization process at the elevated temperature,
where the overamount KOH functions as the activating agent
by reacting with carbon forming noncatalytic potassium com-
pounds (K2O, K2CO3, etc., Figure S1a, Supporting Informa-
tion)[32–34] and the NiO are reduced to metallic Ni NPs and
serve as the catalyst for the subsequent graphitization.[35–37]
Simultaneously, the carbon around Ni is catalyzed at 850 °C to
form a graphitic nanoshell (Figure S1b, Supporting Informa-
tion),[35] while the carbon away from Ni remains amorphous,
demonstrating a catalyst confined graphitization process.
The possible reaction equations are listed in the Supporting
Information.[32–37] The graphitic degree of the HCs is tailored
through adjusting the calcination temperature to 750 and
950 °C, respectively. The as-formed noncatalytic potassium
compounds together with the inorganic impurities from carbon
precursor also affect the catalysis reaction, leading to the dis-
continuous feature of the graphitic stripes. These impurities
are removed with acid washing, and form nanopores between
the microcrystal stripes in the graphitic domains. The scanning
electron microscopic (SEM) images and the corresponding
elemental mapping of the N-GCNs before acid washing
(Figure 1b–e; Figure S1c, Supporting Information) illustrate the
interconnected graphitic carbon nanoshells and the imbedded
Ni NPs with noncatalytic potassium compound impurities,
which agree well with the above explanations.
Raman spectra demonstrate (Figure 2a) three characteristic
peaks including D-band (≈1345 cm−1), G-band (≈1570 cm−1),
and 2D-band (≈2700 cm−1). The D-band reflects the disorder
degree of the crystal structure due to the breathing motion of
sp2 atoms in the amorphous carbon or at edge planes/defects
of the graphitic domains, while G-band represents a scattering
E2g order vibration mode of atoms in rings and chains.[38,39]
2D-band is a two-phonon resonance Raman spectra peak,
reflecting the extent of the stacked graphene layers.[40] The
ID/IG ratio of N-GCNs is calculated as 0.85, indicating the
Adv. Funct. Mater. 2018, 1706294
Figure 1. The synthetic mechanism inference of N-GCNs. a) Schematic illustration of the synthesis process of N-GCNs. b) SEM images of the
unwashed sample and c–e) the corresponding element mapping images.

www.afm-journal.dewww.advancedsciencenews.com
1706294 (3 of 10) © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
highly ordered structure of the as prepared HC. The well-
crystalline structure of N-GCNs is confirmed by the X-ray dif-
fraction (XRD) patterns (Figure S2a, Supporting Information),
which exhibits a sharp peak at 2
θ
–26.4°, corresponding to the
(002) plane of the graphitic domains with expanded interlayer
distance of 3.373 Å,[41] while the broad hump is attributed to
the amorphous structure in the partially graphitized samples.
For comparison, the diffraction patterns (Figure S2a, Sup-
porting Information) and Raman spectra (Figure S2b, Sup-
porting Information) of the control sample prepared at 750 °C
(N-GCNs-750) and 950 °C (N-GCNs-950) were also obtained,
which suggest the enhanced crystallinity with the increasing
heat temperature. The surface chemical composition of the
N-GCNs was examined by X-ray photoelectron spectrum (XPS),
indicating the successful doping of N (1.3 at%) in the N-GCNs
(Figure S2c, Supporting Information). The high-resolution N1s
spectrum (Figure 2b) can be deconvoluted into three peaks
ascribed to the pyridinic N (398.3 eV), pyrrolic N (399.9 eV), and
graphitic N (401.2 eV), with a proportion of 24.5%, 57%, and
18.5%, respectively.[32,42] The doping of N to the carbon archi-
tecture can introduce more active sites and facilitate the charge
transfer, thus significantly enhance the Li/Na storage capa-
bility.[25,26] In addition, the very high content of pyridinic and
pyrollic N in the N-GCNs, are believed more beneficial to the
storage of Li/Na ions.[43] The Fourier-transform infrared (FTIR)
spectra (Figure S2f, Supporting Information) demonstrate
a strong characteristic peak at 3435 cm−1, which is attributed
to the stretching vibration of the OH bond. Small peaks at
2920 and 2840 cm−1 arise from the stretching vibration of the
CH (CH and HCO) bond. The peaks at 1700, 1620,
1380, 1210, and 1090 cm−1, can be assigned to the stretching
vibration of carboxyl groups CO, CC, CO, COC, and
Adv. Funct. Mater. 2018, 1706294
Figure 2. a) Raman spectra, b) deconvoluted XPS spectra (N1s), c) porosity distribution curve by original density functional theory of N-GCNs, d) SEM
image, e,g) amplified SEM, f,h,i) TEM image, and j–l) HRTEM images of N-GCNs (inset: enlarged image).

www.afm-journal.dewww.advancedsciencenews.com
1706294 (4 of 10) © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
CC bond. The peak at 1310 cm−1 is corresponding to the
stretching vibration of CN band.[32] It has been reported that,
proper amount of oxygen functional groups on the carbon
framework can react with lithium/sodium ions, resulting in
improved storage capacity.[44] For example, the CO groups can
be reduced by Li+/Na+ when the voltage versus Li/Na is higher
than 1.5 V (Equation (S7), Supporting Information), thus con-
tributing a certain degree of capacity.[45,46] The type I isotherm
plot from nitrogen adsorption–desorption test (Figure S2g,
Supporting Information) suggests the dominant amount of
micropores,[47] with a total Brunauer–Emmett–Teller (BET) spe-
cific surface area of 1490 m2 g−1 and pore volume of 2.03 cm3
g−1, and the porosity distribution by original density functional
theory (Figure 2c) reveals the hierarchical porous structure of
N-GCNs. The SEM image (Figure 2d) illustrates the unique
3D interconnected carbon nanoshell architecture of N-GCNs
with diameter sizes from 200 to 600 nm. The amplified SEM
images (Figure 2e,g) and the corresponding transmission
electron microscopic (TEM) images (Figure 2f,h) clearly illus-
trate the open hollow characteristic of the nanoshell, with the
wall thickness of ≈20–50 nm (Figure 2i). The high-resolution
TEM (HRTEM) images (Figure 2j–l) demonstrate the tur-
bostratic carbon structure of N-GCNs, with ordered nanogra-
phitic domains (the nanovoids surrounded by some parallel
carbon hexagonal layers). The graphite microcrystal stripes in
the ordered nanodomains are discontinuous, and some gra-
phitic layers are crisscrossed with each other in the ordered
nanodomains, giving rise to defects such as edges and pores,
which will provide sufficient active sites for Li+/Na+ storage.[14]
The HRTEM confirms the expanded interlayer distance of
3.47–4.0 Å, which would favor the intercalation of Li+/Na+ ions.
By contrast, the detailed characterization of N-GCNs-750 and
N-GCNs-950 is also provided (Figures S2c–f and S3, Tables S1
and S2, Supporting Information).
2.2. Electrochemical Performances
The Li+ storage behavior in N-GCNs was evaluated using cyclic
voltammetry (CV) and galvanostatic charge–discharge (GCD)
techniques (Figure 3; Figure S4, Supporting Information).
Figure S4a shows the CV curves for the first five cycles of the
electrode at a scanning rate of 0.1 mV s−1 in the voltage range
of 0.01–3.0 V versus Li+/Li. In the first discharge cycle, the irre-
versible peaks at 0.75 V of LIBs should be attributed to the solid
electrolyte interphase (SEI) film formation and the decomposi-
tion of the electrolyte, and the peak at 1.7 V is caused by the
loss of some irreversible lithium storage sites during the initial
lithiation process.[48,49] While the sharp peak at ≈0.01 V can also
be imputed to the irreversible reaction of electrolyte with sur-
face functional groups and the formation of SEI film.[50] The
three pronounced peaks at ≈0.23 V in the corresponding charge
cycles (Figure S4a, Supporting Information) are attributed to
the delithiation of the graphitic carbon layers,[51] which become
more intense with the increase of cycle count attributed to the
electrochemical activation process.[52] The overlap of the CV
curves in the subsequent four cycles indicates the reversible
electrochemical reactions in the electrode. The N-GCNs display
a high discharge capacity of 1253 mA h g−1 in the second cycle
at 0.1 A g−1 for N-GCNs, which stabilizes at 1236 mA h g−1 after
100 cycles (Figure 3a). The corresponding GCD voltage profiles
are illustrated in Figure 3b. The initial coulombic efficiency is
calculated as 53.4%, the capacity loss should be attributed to the
formation of an SEI layer and the decomposition of the elec-
trolyte on the surfaces of the HC nanoshells.[3] A plateau at
≈0.23 V is detected in all charge voltage profiles, corresponding
to the delithiation of the graphitic carbon layers. Most promis-
ingly, the specific capacity of N-GCNs remains 175 mA h g−1
even at a ultrahigh rate of 20 A g−1 (Figure 3c,d), and when
the current density directly reduced from 20 to 0.1 A g−1,
the capacity quickly recovered to 965 mA h g−1 (increased to
1160 mA h g−1 gradually), demonstrating excellent charge
transfer kinetics within the N-GCNs architecture. Promis-
ingly, the N-GCNs also exhibit impressive electrochemical per-
formance in SIBs (Figure 3e–h), with a reversible discharge
capacity of 325 mA h g−1 at 100 mA g−1, which stabilizes at
174 mA h g−1 after 200 cycles. A capacity of 63 mA h g−1 can still
be maintained at high rate of 5 A g−1. Interestingly, there is no
pronounced peak within the potential range of 0.01–3.0 V (vs
Na+/Na) for SIBs (Figure S4b, Supporting Information), indi-
cating that Na+ cannot be intercalated to the graphitic carbon
layers because of its larger ionic radius. Note that, fluctuation
is observed from the cycling curves for LIBs (Figure 3a), which
does not happen for NIBs (Figure 3e). This maybe because of
that, during the charge–discharge process, the intercalation
of lithium ions would cause the internal structure changes of
carbon materials, and formation of new electrode/electrolyte
interfaces, providing more active sites for lithium storage.
Meanwhile, the uninterrupted reformation of SEI on the new
interface will consume the electrolyte and has a certain nega-
tive influence on the energy storage of the materials. While, the
sodium ions are mainly stored through absorption/desorption
processes, which have little influence on the structure of the
material, resulting in a stable cycling performance. This result
agrees well with the literatures.[1,22,53] For comparison, the
corresponding LIBs/SIBs test for the samples prepared at
750 and 950 °C with varied graphitic degree was also conducted
(Figures S4 and S5, Supporting Information). A comparison
between the electrochemical performance of N-GCNs and
carbon materials previously reported are provided for both LIBs
(Table S3, Supporting Information) and SIBs (Table S4, Sup-
porting Information), indicating the superior lithium/sodium
storage capability of N-GCNs. The outperforming performance
of N-GCNs in terms of reversible capacity, cycling stability, and
rate performance, etc., implies that the unique nanostructure
of the as-prepared HCs plays a critical role in electrochemical
applications, which inspires us to further investigate it with in
situ Raman spectroscopy and theoretical calculations.
2.3. In Situ Raman Spectroscopy Study
The in situ Raman spectroscopy study was conducted during
the charge/discharge process of both LIBs and SIBs, to
moni tor the ion interclation or adsorption in the carbon mate-
rials,[48,54,55] and unravel the critical potential for ion adsorp-
tion on the carbonaceous surface and intercalation into the
graphitic layers.[56,57] Figure 4a–d shows the in situ Raman
Adv. Funct. Mater. 2018, 1706294

www.afm-journal.dewww.advancedsciencenews.com
1706294 (5 of 10) © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
Adv. Funct. Mater. 2018, 1706294
Figure 3. Electrochemical properties. Electrochemical performance of N-GCNs tested against Lithium and Sodium: cycling performance and GCD
profiles of N-GCNs at a current density of 0.1 A g−1 in LIBs. a,b) Cycling performance and GCD profiles of N-GCNs at a current density of 0.1 A g−1 in
LIBs. c,d) GCD profiles of N-GCNs at different rates of 0.1 to 20 A g−1 and rate performance gradually increasing from 0.1 to 20 A g−1 and then back
to 0.1 A g−1 in LIBs. e,f) Cycling performance and GCD profiles of N-GCNs at a current density of 0.1 A g−1 in SIBs. g,h) GCD profiles of N-GCNs at
different rates of 0.1 to 5 A g−1 and rate performance gradually increasing from 0.1 to 5 A g−1 and then back to 0.1 A g−1 in SIBs.

www.afm-journal.dewww.advancedsciencenews.com
1706294 (6 of 10) © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
spectra and the corresponding discharge–charge profiles
at the first cycle of N-GCNs for LIBs (Figure 4a,b) and SIBs
(Figure 4c,d) with a current density of 0.25 A g−1. The intercala-
tion of guest ions into carbonaceous host strongly affects the
position, shape, and intensity of the D and G-band (≈1330 and
≈1600 cm−1 at the open-circuit voltage (OCV, ≈3.0 V vs Li/Li+
or Na/Na+)). The decreased intensity of D-band from OCV to
0.01 V can be attributed to the introduction of Li+/Na+, which
occupy the defective sites, i.e., the nanopores in the nanogra-
phitic domains and the amorphous structure of N-GCNs,
thus the breathing motion of sp2 atoms in the rings at edge
planes is limited by the absorbed Li+/Na+ ions. Meanwhile, the
G-band of N-GCNs shows unconspicuous red-shifts from OCV
to ≈0.2 V in LIBs (Figure 4a) and from OCV to 0.01 V in SIBs
(Figure 4c) during discharge process. This can be accounted
for the charge transfer effects from the adsorption of Li+/Na+
ions and the according formation of SEI layer (below ≈0.8 V
in LIBs (Figure 4b) and ≈0.6 V in SIBs (Figure 4d)).[49] These
Li+/Na+ ions randomly accommodate the nanoporous defective
sites and fully cover the N-GCNs surface at relatively higher
voltage, thus causing the early disappearance of G-band. The
intensity of G-band steadily decreases with potential and
finally fades into the signal noise from ≈0.2 to 0.01 V in LIBs
because of the weaken of resonance caused by Li+ intercalation,
which indicates the lithiation of the nanographitic domains
at ≈0.2 V.[49,58] This result is further confirmed with ex situ
XRD test (Figure S6, Supporting Information). Meanwhile,
a large amount of nanovoids filling is in accordance with the
plateau region close to 0.01 V in the corresponding discharge
curves (Figure 4b). However, when the N-GCNs are explored
as SIB electrode, the discharge curves exhibit smooth slope
from OCV to 0.6 V and 0.6 to 0.01 V, neither evident inten-
sity decrease from G-band, nor corresponding plateau region
in the discharge and charge curves are observed (Figure 4b).
This indicates no obvious ion intercalation happened in the
graphitic domains due to the larger ionic radius of Na+. Thus,
the Na+ can only be adsorbed on the disordered carbon archi-
tectures of N-GCNs, and the nanovoids filling region is also
close to 0.01 V.[50] So the “adsorption–intercalation mechanism”
for Li+ storage and “adsorption mechanism” for Na+ storage
of N-GCNs are verified through the in situ Raman analysis.
The position, shape, and intensity of the D and G-band was
recovered progressively because of the extraction of the Li+/Na+
during the charge process which also indicates that the good
reversibility of the materials.[57]
2.4. The First-Principles Calculations
The first-principles calculations were performed to unveil the
origin of excellent electrochemical performance of N-GCNs,
based on DFT as implemented in the Vienna ab initio simula-
tion package. The defective nanographitic domain of N-GCNs
was modeled using a periodically repeated slab of N-doped
graphene with nanopore defects, denoted as NDGs. By con-
trast, control systems including NGs (N-doped nondefective
graphene sheet), DGs (doping free defective graphene sheets),
and PGs (pristine graphene sheets) were also constructed. The
most stable electronic structures for Li absorbed systems were
proposed and optimized (Figure 5a–d) and the corresponding
binding energy (EB) is calculated by Equation (1) (Table S5,
Supporting Information)
=−EEE
BPR
(1)
where EP is the total energy of the Li absorbed graphene and ER
is the total energy of an isolated Li atom and the substrate. The
results demonstrate that, the most stable site for Li adsorption
is at the center of a carbon hexagon (hollow site) for both PGs
and NGs, demonstrating that the Li atom prefers C rather than
N atoms in the graphene structures. The EB between absorbed
Li atom and substrate NGs/PGs is calculated as −0.49/−0.95 eV
(Table S5, Supporting Information), indicating weaker interac-
tion between Li and NGs. Therefore, simple nitrogen doping
cannot effectively improve the Li atom adsorption capa-
bility of graphene. Interestingly, the most stable sites of DGs
(Figure 5c) and NDGs (Figure 5d) are at the center of nano-
pore defects. This implies that, the electron-accepting tendency
of the defective nanoarchitecture is beneficial to gaining elec-
tron from the Li atom.[59] As compared to NGs and PGs, these
defective structures demonstrate several folds of improvement
on the binding energy, i.e., −3.61 eV for NDGs, and −3.17 eV
Adv. Funct. Mater. 2018, 1706294
Figure 4. In-situ Raman spectroscopy test. In situ Raman spectra and the corresponding discharge-charge profiles at the first cycle of N-GCNs for LIBs
a,b) and SIBs c,d) with the current density of 0.25 A g−1, respectively.

www.afm-journal.dewww.advancedsciencenews.com
1706294 (7 of 10) © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
for DGs, suggesting that Li adsorption capability of nanogra-
phitic domains can be significantly enhanced by introducing
nanopore defects. On the other hand, the more negative EB
of NDGs than DGs provides direct evidence of the synergistic
effect between N-doping and nanopore defects for Li storage.[43]
To get further insight into the electronic structure during
the lithiation process, the charge density difference (yellow/
light blue color represents positive/negative charge differences,
respectively) between the charge densities of the adsorp-
tion system and those of the separated substrate/adatom (at
optimized positions) (Equation (2))[60] was also calculated
(Figure 5e–h; Figure S7a,b, Supporting Information)
ρρ ρρ
∆= +−−[Lisubstrate][Li][substrate]
(2)
Figure 5e,f illustrates the large degree of charge transfer
between Li and the adjacent C atoms in PGs, whereby Li is
obviously oxidized by transferring the valence electrons of 2s to
carbon.[61] Whereas the NGs exhibits an electron rich character
which leads to the depressed adsorption ability for Li atoms
and attenuated charge transfer between Li and the adjacent C
atoms (Figure S7a,b, Supporting Information). Interestingly,
the introduction of defects could cause electron deficiency
and endow the graphene with an electron accepting tendency.
Therefore, the charge transfer ability between the lithium
and the carbon atoms are significantly enhanced for NDGs
(Figure 5g,h; Figure S7c,d, Supporting Information), which
would favor improved charge transfer kinetics when applied as
battery electrode.[31]
The electronic DOS for the above systems before and after
lithiation were also calculated (Figure 5; Figure S7, Supporting
Information). The PGs exhibits semiconductor characteristic
with the Dirac point at Fermi level (where the DOS is zero)
clearly visible before lithiation (Figure 5i). The Fermi level is
increased and shifted to the right after lithiation (Figure 5j),
which should be attributed to that the absorbed Li atom trans-
fers part of electrons to the conduction band of the PGs, and
the DOS near Fermi level is improved but bring little effect.
Hence, the pure graphene may not meet the requirements of
the faster electron transport kinetics. As compared to PGs,
the NDGs demonstrate prominent increment of electronic
states near the Fermi level both before and after lithiation
(Figure 5k,l). This suggests the enhanced electron mobility
within NDGs, which can greatly enhance the rate performance
of the electrode material during charge–discharge process.[31,61]
Thus, the DFT calculations evidenced the coupling effect
between N-doping and defects in the nanographitic domains of
N-GCNs for Li storage, which demonstrates improved binding
energy and charge transfer, as well as increased DOS around
Fermi level.
Adv. Funct. Mater. 2018, 1706294
Figure 5. The first-principles calculations. Models a–d) of the PGs, NGs, DGs, and NDGs, respectively. Charge density difference e,f) of an Li atom
adsorbed on the pristine graphene sheets, charge density difference g,h) of an Li atom adsorbed on the N-doped defective graphene sheets (isosur-
faces level = 0.0015e bohr−3; yellow and light blue areas represent positive and negative charge differences). i,j) Corresponding DOS illustrations of the
pristine graphene sheets before and after lithiation, respectively. k,l) Corresponding DOS illustrations of the N-doped defects graphene sheets before
and after lithiation, respectively. The dashed line at 0 eV indicates the Fermi level.

www.afm-journal.dewww.advancedsciencenews.com
1706294 (8 of 10) © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
Adv. Funct. Mater. 2018, 1706294
2.5. Charge Storage Mechanism of N-GCNs
Based on the above analysis, the as-prepared N-GCNs offer
obvious structural advantages in improving the Li/Na storage
capability (Figure 6): (i) The interconnected graphitic nanoshells
(diameter sizes from 200 to 600 nm) with small wall thickness of
20–50 nm can not only make certain fast and successive trans-
portation of electrons in N-GCNs, but also greatly reduce the
solid-state transport pathway for Li+/Na+ diffusion (Figure 6a)
because the electrolyte ions can be stored in both inside and out-
side surface of the ultrathin nanoshell, and the unique nanoshell
structure provides additional open hollow space in sub-micro-
scale, which can buffer volume expansion well during Li+/Na+
insertion and extraction (especially at high rates), assuring the
long cycle stability. (ii) The graphitic domains in the nanoshells
such as intermittent graphitic stripes (with short segments of
graphene layers) (Figure 6b) and crisscrossed graphene layers
(with different directions of ion access points) (Figure 6c) can
greatly increase the electronic transfer and ion diffusion coef-
ficient because of its high conductivity and short ion diffusion
pathway. (iii) The nanographitic domains (Figure 6d) can give
rise to defects such as edges and pores, which will provide suf-
ficient active sites to absorb Li+/Na+ storage, and the expanded
interlayer distance of 3.47–4.0 Å and the amorphous carbon with
nanopores (Figure 6e) can also favor the intercalation/interfacial
storage of Li+/Na+ ions, enabling high reversible capacity of the
N-GCNs, which is evidenced by in situ Raman study. (iv) The
coupled N-doping and nanopore defects in the nanographitic
domains can significantly improve the binding ability for
Li+/Na+, charge transfer ability, increase the DOS around Fermi
level before/after lithiation, thus enhance the charge transport
kinetics of the HCs with improved rate capability. Besides, the
surface oxygen groups introduced by the acidic treatment would
also contribute capacity through Faradaic reactions.
3. Conclusions
In conclusion, N-GCNs with homogeneous defective nanogra-
phitic domains are fabricated through the prechelation
between Ni2+ and chitosan with simulta-
neous N-doping and KOH-activation. The
“adsorption mechanism” of N-GCNs for Na+
storage and the “adsorption–intercalation
mechanism” of N-GCNs for Li+ storage are
verified through in situ Raman analysis.
DFT calculations are indicates the coupled
effect of heteroatom-doping and nanopore
defects down to the atomic scale in the nano-
graphitic domains of the N-GCNs. The as-
prepared N-GCNs deliver a high reversible
lithium storage capacity of 1253 mA h g−1,
and outstanding rate performance which out-
performs most state-of-the-art hard carbons.
Interestingly, a high reversible capacity and
long cycle stability for sodium storage is also
obtained. This excellent long cycling perfor-
mance combined with the high capacity and
rate capability indicates that the N-GCNs are
promising anodes materials for application in rechargeable
LIBs and SIBs.
4. Experimental Section
Materials Preparation: The Artemia cyst shells with the main
component of chitosan were pretreated as described in the previous
paper.[32] Specifically 2.4 g as pretreated precursor was dispersed
into 200 mL deionized water and kept stirring for 6 h. 100 mL
0.025 m Ni(AC)2·4H2O, was then added dropwise and stirred overnight.
Subsequently, 18 mL 6 m KOH was added dropwise into the above
mixture and heated at 100 °C until the whole mixture turned into a
paste. The as-obtained materials were then put into a tube furnace and
heated to 850 °C with the ramping speed of 2 °C min−1 and held there for
4 h. The as-obtained product was treated with concentrated 8 m HNO3,
1 m HCl solution and flushed with deionized water, respectively. After
being cleaned and dried, the final product was obtained and denoted
as N-GCNs. For comparison, the control samples calcined at 750 and
950 °C were also prepared following the same procedure, denoted as
N-GCNs-750 and N-GCNs-950, respectively.
Characterization: Raman spectra were recorded on a Renishaw in Via
Raman microscope with an Ar ion laser at the excitation wavelength
of 532 nm, and in situ Raman spectra (Witech Alpha 300R, 532 nm
wavelength) were tested at the excitation wavelength of 633 nm. Powder
XRD patterns between 10 and 50° (≈2
θ
) were collected by Rigaku D/
MAX-2500 powder diffractometer with Cu-K
α
radiation (
λ
= 0.154 nm)
operated at 40 kV, 200 mA. Hitachi-XPS was measured by a VG ESCALAB
MKIIX-ray photoelectron spectrometer using Mg-K
α
as the exciting
source (1253.6 eV). Carl Zeiss SUPRA 55 SAPPHIRE field emission
scanning electron microscope (Germany, 15 kV), Hitachi-7650 TEM
(Japan, 80 kV), and HRTEM (JEOL JEM-3000F) were used to investigate
the morphology and microstructure.
Electrochemical Measurements: The working electrodes were prepared
by mixing 80 wt% synthesized active materials, 10 wt% acetylene carbon
black, 10 wt% polyvinylidene fluoride (PVDF) binder in N-methyl-2-
pyrrolidone onto a Cu foil and then dried in a vacuum furnace at 120 °C
for 12 h. The 25 µm microporous monolayer membrane (PP, Celgard
2400) and Li/Na metal were used as separator and counter electrodes,
respectively, when assembled into a stainless-steel coin cell (2032) in an
Ar-filled glovebox (both O2 and H2O levels below 0.1 ppm). The charge
storage performance was tested in 1 m LiPF6 solution in a mixture (1:1:1,
vol%) of dimethyl carbonate, ethylene carbonate (EC), and ethylmethyl
carbonate for LIBs, and 1 m NaClO4 in a mixture (1:1, vol%) of EC and
diethyl carbonate solution for SIBs. CV curves were measured at the
Figure 6. Schematic illustration of Li (Na) storage mechanism in N-GCNs. a–e) Schematic
illustration of Li (Na) storage mechanism in N-GCNs.

www.afm-journal.dewww.advancedsciencenews.com
1706294 (9 of 10) © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
Adv. Funct. Mater. 2018, 1706294
scanning rate of 0.1 mV s−1. Electrochemical impedance spectroscopy
(EIS) spectra were recorded at the frequency range of 0.01 to 105 Hz.
Both of EIS and CV were test on CHI650e electrochemical workstation
(Chenhua, China). The constant current charge–discharge tests were
performed with a computer controlled cycling equipment (Land
CT2001A, China) in the potential rage of 0.01–3 V.
Supporting Information
Supporting Information is available from the Wiley Online Library.
Acknowledgements
The authors thank the financial supports from the National Natural
Science Foundation of China (51774251), Hebei Natural Science
Foundation for Distinguished Young Scholars (B2017203313 and
B2015203096), Hundred Excellent Innovative Talents Support Program
in Hebei Province (SLRC2017057), State Key Project of Research and
Development of China (2016YFA0200102), NSFC-RGC Joint Research
Scheme (51361165201), Scientific Research Foundation for the Returned
Overseas Chinese Scholars (CG2014003002), and the open funding from
State Key Laboratory of Advanced Technology for Materials Synthesis
and Processing (Wuhan University of Technology, 2017-KF-14).
Conflict of Interest
The authors declare no conflict of interest.
Keywords
hard carbon nanoshells, lithium/sodium storage, N-doping,
nanographitic domains
Received: October 30, 2017
Revised: December 1, 2017
Published online:
[1] Z. L. Yu, S. Xin, Y. You, L. Yu, Y. Lin, D. Xu, C. Qiao, Z. Huang,
N. Yang, S. Yu, J. B. Goodenough, J. Am. Chem. Soc. 2016, 138,
14915.
[2] J. B. Goodenough, K. S. Park, J. Am. Chem. Soc. 2013, 135, 1167.
[3] Z. Jian, Z. Xing, C. Bommier, Z. Li, X. Ji, Adv. Energy Mater. 2016, 6,
1501874.
[4] M. N. Tahir, B. Oschmann, D. Buchholz, X. Dou, I. Lieberwirth,
M. Panthofer, W. Tremel, R. Zentel, S. Passerini, Adv. Energy Mater.
2016, 6, 1501489.
[5] Z. Jian, Z. Xing, C. Bommier, Z. Li, X. Ji, Adv. Energy Mater. 2016, 6,
1501874.
[6] F. A. Soto, P. Yan, M. H. Engelhard, A. Marzouk, C. Wang, G. Xu,
Z. Chen, K. Amine, J. Liu, V. L. Sprenkle, Adv. Mater. 2017, 29,
1606860.
[7] L. Xiao, Y. Cao, W. A. Henderson, M. L. Sushko, Y. Shao, J. Xiao,
W. Wang, M. H. Engelhard, Z. Nie, J. Liu, Nano Energy 2016, 19,
279.
[8] Y. Cao, L. Xiao, M. L. Sushko, W. Wang, B. Schwenzer, J. Xiao,
Z. Nie, L. V. Saraf, Z. Yang, J. Liu, Nano Lett. 2012, 12, 3783.
[9] S. Qiu, L. Xiao, L. M. L. Sushko, K. S. Han, Y. Shao, M. Yan,
X. Liang, L. Mai, J. Feng, Y. Cao, X. Ai, H. Yang, J. LiuAdv. Energy
Mater. 2017, 7, 1700403.
[10] S. Wang, C. Xiao, Y. Xing, H. Xu, S. Zhang, J. Mater. Chem. A 2015,
3, 6742.
[11] A. Piotrowska, K. Kierzek, P. Rutkowski, J. Machnikowski, J. Anal.
Appl. Pyrolysis 2013, 102, 1.
[12] F. Bonino, S. Brutti, P. Reale, B. Scrosati, L. Gherghel, J. Wu,
K. Müllen, Adv. Mater. 2005, 17, 743.
[13] K. Hong, Q. Long, R. Zeng, Z. Yi, W. Zhang, D. Wang, W. Yin,
C. Wu, Q. Fan, W. Zhang, J. Mater. Chem. A 2014, 2, 12733.
[14] Y. Li, Y. Hu, H. Li, L. Chen, X. Huang, J. Mater. Chem. A 2016, 4, 96.
[15] D. A. Stevens, J. R. Dahn, J. Electrochem. Soc. 2000, 147, 1271.
[16] Y. Xu, Z. Lin, X. Zhong, B. Papandrea, Y. Huang, X. Duan, Angew.
Chem. 2015, 54, 5345.
[17] Y. Fang, Y. Lv, R. Che, H. Wu, X. Zhang, D. Gu, G. Zheng, D. Zhao,
J. Am. Chem. Soc. 2013, 135, 1524.
[18] X. Huang, H. Yu, J. Chen, Z. Lu, R. Yazami, H. H. Hng, Adv. Mater.
2014, 26, 1296.
[19] Y. Fang, Y. Lv, R. Che, H. Wu, X. Zhang, D. Gu, G. Zheng, D. Zhao,
J. Am. Chem. Soc. 2013, 135, 1524.
[20] X. Yang, L. Zhang, F. Zhang, Y. Huang, Y. Chen, ACS Nano 2014, 8,
5208.
[21] H. Zhou, S. Zhu, M. Hibino, I. Honma, M. Ichihara,
J. Cheminformatics 2004, 35, 2107.
[22] Z. Lyu, L. Yang, D. Xu, J. Zhao, H. Lai, Y. Jiang, Q. Wu, Y. Li,
X. Wang, Z. Hu, Nano Res. 2015, 8, 3535.
[23] L. Chen, Z. Wang, C. He, N. Zhao, C. Shi, E. Liu, J. Li, ACS Appl.
Mater. Interfaces 2013, 5, 9537.
[24] F. D. Han, Y. J. Bai, R. Liu, B. Yao, Y. X. Qi, N. Lun, J. X. Zhang, Adv.
Energy Mater. 2011, 1, 798.
[25] X. Liu, J. Zhang, S. Guo, N. Pinna, J. Mater. Chem. A 2016, 4, 1423.
[26] D. Xie, X. Xia, Y. Zhong, Y. Wang, D. Wang, X. Wang, J. Tu, Adv.
Energy Mater. 2017, 7, 1601804.
[27] A. L. M. Reddy, A. Srivastava, S. R. Gowda, H. Gullapalli, M. Dubey,
P. M. Ajayan, ACS Nano 2010, 4, 6337.
[28] J. Dong, Y. Xue, C. Zhang, Q. Weng, P. Dai, Y. Yang, M. Zhou, C. Li,
Q. Cui, X. Kang, Adv. Mater. 2017, 29, 1603692.
[29] Z. Pei, H. Li, Y. Huang, Q. Xue, Y. Huang, M. Zhu, Z. Wang, C. Zhi,
Energy Environ. Sci. 2017, 10, 742.
[30] J. Xu, M. Wang, N. P. Wickramaratne, M. Jaroniec, S. Dou, L. Dai,
Adv. Mater. 2015, 27, 2042.
[31] Y. Yang, D. M. Tang, C. Zhang, Y. Zhang, Q. Liang, S. Chen,
Q. Weng, M. Zhou, Y. Xue, J. Liu, Energy Environ. Sci. 2017, 10,
979.
[32] Y. Zhao, S. Huang, M. Xia, S. Rehman, S. Mu, Z. Kou, Z. Zhang,
Z. Chen, F. Gao, Y. Hou, Nano Energy 2016, 28, 346.
[33] Z. Zhang, S. Mu, B. Zhang, L. Tao, S. Huang, Y. Huang, F. Gao,
Y. Zhao, J. Mater. Chem. A 2016, 4, 2137.
[34] J. Wang, S. Kaskel, J. Mater. Chem. 2012, 22, 23710.
[35] A. Ōya, H. Marsh, J. Mater. Sci. 1982, 17, 309.
[36] J. Gillot, B. Lux, P. Cournault, F. du Chaffaut, Ber Deut. Keram. Ges.
1968, 45, 224.
[37] A. Oberlin, M. Endo, T. Koyama, J. Cryst. Growth 1976, 32, 335.
[38] R. P. Vidano, D. B. Fischbach, L. J. Willis, T. M. Loehr, Solid State
Commun. 1981, 39, 341.
[39] A. C. Ferrari, J. C. Meyer, V. Scardaci, C. Casiraghi, M. Lazzeri,
F. Mauri, S. Piscanec, D. Jiang, K. S. Novoselov, S. Roth, Phys. Rev.
Lett. 2006, 97, 187401.
[40] Y. Hao, Y. Wang, L. Wang, Z. Ni, Z. Wang, R. Wang, C. K. Koo,
Z. Shen, J. T. Thong, Small 2010, 6, 195.
[41] W. Yang, H. Kai, Y. Zhu, F. Han, Y. Xu, I. Matsuda, Y. Ishii,
J. Cumings, C. Wang, Nat. Commun. 2014, 5, 4033.
[42] U. Sim, J. Moon, J. An, J. H. Kang, S. E. Jerng, J. Moon, S. P. Cho,
B. H. Hong, K. T. Nam, Energy Environ. Sci. 2015, 8, 1626.
[43] C. Ma, X. Shao, D. Cao, J. Mater. Chem. 2012, 22, 8911.
[44] M. Lu, W. Yu, J. Shi, W. Liu, S. Chen, X. Wang, H. Wang, Electrochim.
Acta 2017, 251, 369.

www.afm-journal.dewww.advancedsciencenews.com
1706294 (10 of 10) © 2018 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
Adv. Funct. Mater. 2018, 1706294
[45] S. W. Lee, N. Yabuuchi, B. M. Gallant, S. Chen, B. S. Kim,
P. T. Hammond, S.-H. Yang, Nat. Nanotechnol. 2010, 5, 531.
[46] Y. Shao, J. Xiao, W. Wang, M. Engelhard, X. Chen, Z. Nie, M. Gu,
V. S. Laxmikant, E. Gregory, J. Zhang, L. Jun, Nano Lett. 2013, 13,
3909.
[47] J. Niu, R. Shao, J. Liang, M. Dou, Z. Li, Y. Huang, F. Wang, Nano
Energy 2017, 36, 322.
[48] J. Zhang, H. Zhu, P. Wu, C. Ge, D. Sun, L. Xu, Y. Tang, Y. Zhou,
Nanoscale 2015, 7, 18211.
[49] W. Tang, B. M. Goh, M. Y. Hu, C. Wan, B. Tian, X. Deng, C. Peng,
M. Lin, J. Z. Hu, K. P. Loh, J. Phys. Chem. C 2016, 120, 2600.
[50] Y. Li, Y. S. Hu, M. M. Titirici, L. Chen, X. Huang, Adv. Energy Mater.
2016, 6, 1600659.
[51] H. Song, G. Yang, C. Wang, ACS Appl. Mater. Interfaces 2014, 6,
21661.
[52] Y. Zheng, Y. Wang, Y. Lu, Y. S. Hu, J. Li, Nano Energy 2017, 39,
489.
[53] X. Liu, D. Chao, D. Su, S. Liu, L. Chen, C. Chi, J. Lin, Z. Shen,
J. Zhao, L. Mai, Y. Li, Nano Energy 2017, 37, 108.
[54] L. J. Hardwick, H. Buqa, P. Novák, Solid State Ionics 2006, 177,
2801.
[55] L. J. Hardwick, Ph.D. Thesis, Eidgenössische Technische Hochschule
ETH Zürich, 2007.
[56] L. J. Hardwick, M. Hahn, P. Ruch, M. Holzapfel, W. Scheifele,
H. Buqa, F. Krumeich, P. Novák, R. Kötz, Electrochem. Acta 2006,
52, 680.
[57] L. J. Hardwick, P. W. Ruch, M. Hahn, W. Scheifele, R. Kötz,
P. Novák, J. Phys. Chem. Solids 2008, 69, 1232.
[58] C. Sole, N. E. Drewett, L. J. Hardwick, Faraday Discuss. 2014, 172,
223.
[59] G. Mpourmpakis, G. Froudakis, J. Chem. Phys. 2006, 125, 56.
[60] F. Valencia, A. H. Romero, F. Ancilotto, P. L. Silvestrelli, J. Phys.
Chem. B 2006, 110, 14832.
[61] K. Shiota, T. Kawai, Jpn. J. Appl. Phys. 2017, 56, 06GE11.