Fabrication of three-dimensionally interconnected nanoparticle superlattices and their lithium-ion storage properties

W1967728946 · 수집 2026-10-04 18:12:18

T10321 Quantum Dots Synthesis And Properties

CoT 라벨링

라벨링 전

PDF 파일

자동 다운로드한 PDF PDF 보기

기본 정보

DOI 제목 발행년도 저널 URL 페이지 열기 PDF URL PDF 열기 발행기관 피인용수 라이선스

추가 정보 초록

Abstract Three-dimensional superlattices consisting of nanoparticles represent a new class of condensed materials with collective properties arising from coupling interactions between close-packed nanoparticles. Despite recent advances in self-assembly of nanoparticle superlattices, the constituent materials have been limited to those that are attainable as monodisperse nanoparticles. In addition, self-assembled nanoparticle superlattices are generally weakly coupled due to the surface-coating ligands. Here we report the fabrication of three-dimensionally interconnected nanoparticle superlattices with face-centered cubic symmetry without the presynthesis of the constituent nanoparticles. We show that mesoporous carbon frameworks derived from self-assembled supercrystals can be used as a robust matrix for the growth of nanoparticle superlattices with diverse compositions. The resulting interconnected nanoparticle superlattices embedded in a carbon matrix are particularly suitable for energy storage applications. We demonstrate this by incorporating tin oxide nanoparticle superlattices as anode materials for lithium-ion batteries, and the resulting electrochemical performance is attributable to their unique architectures.

본문

### ARTICLE

Received 13 Nov 2014 | Accepted 27 Jan 2015 | Published 3 Mar 2015 DOI: 10.1038/ncomms7420 OPEN
## Fabrication of three-dimensionally interconnected nanoparticle superlattices and their lithium-ion storage properties

Yucong Jiao 1,2, Dandan Han 2, Yi Ding 1, Xianfeng Zhang 1,2, Guannan Guo 1, Jianhua Hu 1,
Dong Yang 1 & Angang Dong 2

Three-dimensional superlattices consisting of nanoparticles represent a new class of

condensed materials with collective properties arising from coupling interactions between

close-packed nanoparticles. Despite recent advances in self-assembly of nanoparticle

superlattices, the constituent materials have been limited to those that are attainable as

monodisperse nanoparticles. In addition, self-assembled nanoparticle superlattices are

generally weakly coupled due to the surface-coating ligands. Here we report the fabrication of

three-dimensionally interconnected nanoparticle superlattices with face-centered cubic

symmetry without the presynthesis of the constituent nanoparticles. We show that meso
porous carbon frameworks derived from self-assembled supercrystals can be used as a

robust matrix for the growth of nanoparticle superlattices with diverse compositions.

The resulting interconnected nanoparticle superlattices embedded in a carbon matrix are

particularly suitable for energy storage applications. We demonstrate this by incorporating tin

oxide nanoparticle superlattices as anode materials for lithium-ion batteries, and the resulting

electrochemical performance is attributable to their unique architectures.

1 State Key Laboratory of Molecular Engineering of Polymers and Department of Macromolecular Science, Fudan University, Shanghai 200433, China.
2 Department of Chemistry and Collaborative Innovation Center of Chemistry for Energy Materials, Fudan University, Shanghai 200433, China.
Correspondence and requests for materials should be addressed to D.Y. (email: [yangdong@fudan.edu.cn)](mailto:yangdong@fudan.edu.cn) or to A.D. (email: [agdong@fudan.edu.cn).](mailto:agdong@fudan.edu.cn)

hree-dimensional (3D) superlattices consisting of nanoparticles (NPs) are emerging as a new and important class
of nanostructured materials 1, the properties of which can
# T
be rationally tuned by manipulating the size, shape and
composition of the constituent NPs. In particular, interparticle
interactions in NP superlattices can lead to collective properties,
which are significantly different from the isolated NPs 1–3 .
Existing methods of growing NP superlattices rely on the selfassembly of monodisperse colloidal NPs induced by solvent
evaporation or antisolvent destabilization 1 . Recent progress on
the self-assembly of NPs has been witnessed by the growth of a
rich array of single- and multicomponent NP superlattices 4–10,
which have found wide applications in electronic and
optoelectronic devices, catalysis and energy storage.

The self-assembly of colloidal NPs is a complicated process,
involving various driving forces and interactions (for example,
van der Waals, Coulombic, dipolar) 1 . As a consequence,
successful assembly of long-range ordered NP superlattices
requires exquisite control over many parameters, such as NP
size distribution, surface-coating ligands and solvent evaporation
kineticsand so on. In addition, self-assembled NP superlattices
typically suffer from poor electrical conductivity due to the large
interparticle spacing maintained by the capping ligands 1,11 .
Therefore, post-surface treatment such as ligand exchange is
required to enhance electronic coupling, which unfortunately can
lead to severe structural defects such as cracks 12,13 . Another
major bottleneck hampering prospects of NP superlattices is the
limited choice of monodisperse NP building blocks, despite recent
progress in colloidal synthesis 1,14 .

In this work we report an approach that can overcome the
aforementioned limitations associated with self-assembly methods, enabling the growth of 3D interconnected, strongly coupled
NP superlattices without the presynthesis of the constituent NPs.
The resulting connected NP superlattices embedded in 3D
continuous carbon frameworks represent a new class of superlattice materials, which are demonstrated to be remarkably
promising for energy storage applications. SnO2 NP superlattices
are chosen as a model system to study Li-ion storage properties,
which exhibit exceptional cycling stability and rate capability
when used as Li-ion battery (LIB) anodes.

Results
Fabrication procedure. Figure 1a schematically illustrates the
overall fabrication procedure. Briefly, starting with self-assembled
Fe3O4 NP supercrystals, we obtain 3D continuous, highly ordered
mesoporous carbon frameworks on the carbonization of the
surface-coating oleic acid (OA) ligands followed by the removal
of Fe3O4 NPs. The subsequent impregnation of carbon frameworks with desired precursors leads to 3D NP superlattices on
hydrolysis and/or thermal annealing. The constituent NP composition can be modulated by selecting appropriate precursors,
while the NP size can be tailored by controlling that of the original Fe3O4 NPs. The reason to select Fe3O4 NPs as the starting
material is that the scalable production of Fe3O4 NP superlattices
can be routinely achieved by self-assembly 14,15 . In addition,
Fe3O4 is an inexpensive and thermally stable material, which can
be readily removed by acid etching.

It is worth to mention that although the structure of our
mesoporous carbon frameworks appears to be similar to that of
previous mesoporous carbons prepared from hard template (that
is, silica opals or mesoporous silica) 16–22, we should emphasize
that this is the first preparation of ordered mesoporous carbons
from inorganic nanocrystals. There are several unique features
associated with this preparation strategy, leading to mesoporous
carbons with distinctive structural and textural properties. First,

the surface-coating OA ligands, which are a necessary component
for the growth and self-assembly of Fe3O4 NPs, concurrently
serve as the carbon source, yielding mesoporous carbons having
ultrathin pore walls (B2 nm) and the same topological structure
as the Fe3O4 NPs supercrystal template. This is in sharp contrast
with previous hard-templating approaches where additional
carbon precursors (that is, polymers or prepolymers) have to be
applied to fill the voids of the template and the resulting
mesoporous carbons usually possess an inverse structure of the
template 16–22 . Furthermore, unlike previous approaches where
silica opals or mesoporous silica only act as template 16–22, the
Fe3O4 NPs used here also serve as a graphitization catalyst (to be
discussed below), allowing for partially graphitic frameworks at
low temperatures (500 �C). More importantly, the unique
structure of our carbon frameworks enables interconnected,
strongly coupled NP superlattices, which could not be readily
accessible by conventional self-assembly methods.

Mesoporous carbon frameworks. Monodisperse Fe3O4 NPs
(B11 nm) stabilized by OA are synthesized by a literature
method 14 (Supplementary Fig. 1a), and Fe3O4 NP superlattices
are grown by a conventional drying-mediated self-assembly
process (Supplementary Fig. 1b) 4 . The content of the surfacecoating OA ligands determined from thermogravimetric analysis
(TGA) isB15.9 wt% (Supplementary Fig. 1c), corresponding to
the area ligand coverage ofB3.83 nm � 2 (that is,B1455 OA
molecules per NP, Supplementary Note 1). The prior addition of
squalane (B1 wt%) into the NP solution is found to favour the
gram-scale growth of micrometer-sized NP supercrystals with
well-developed facets 15, although squalane is not critical for the
subsequent formation of mesoporous carbon frameworks
(Supplementary Fig. 1d). The as-assembled NP supercrystals are
then heated at 500 �C in an argon atmosphere, producing grey
powders arising from the carbonization of OA (Supplementary
Fig. 2a). The carbon species constituteB11 wt% of the carbonized
NP supercrystals, determined from TGA (Supplementary
Fig. 2b). Scanning electron microscopy (SEM) and highresolution SEM (HRSEM) establish that both the faceted
morphology and the long-range NP ordering are well retained
in carbonized Fe3O4 NP supercrystals (Fig. 1b,c and
Supplementary Fig. 2c–f). These results strongly suggest that
cracks and other structural defects, which have been observed in
many thermally treated NP superlattices reported previously 23,24,
are largely prevented in our experiments, probably due to the 3D
crystal-like morphology as well as the long-range structural
ordering of our Fe3O4 NP supercrystals. The small-angle X-ray
scattering (SAXS) pattern of carbonized Fe3O4 NP supercrystals
exhibits at least four scattering peaks, which can be assigned to
the 111, 220, 311, 420 reflections of a well-crystallized facecentered cubic (fcc) structure (Fig. 1d, black curve), consistent
with the highly ordered superlattice structure observed in
HRSEM (Fig. 1c and Supplementary Fig. 2f). The unit cell
parameter calculated from the SAXS data is 22.8 nm.

HNO3 or HCl etching is employed to treat carbonized NP
supercrystals. The complete removal of Fe3O4 NPs yields black
carbon powders after washing and drying (Supplementary
Fig. 3a), which display a highly ordered porous structure as
revealed by SEM (Supplementary Fig. 3b) and transmission
electron microscopy (TEM, Fig. 1e,f and Supplementary
Fig. 3c,d). SAXS (Fig. 1d, red curve) and SEM (Fig. 1e, inset)
indicate that the porous carbon frameworks possess the same fcc
structure and faceted morphology inherited from Fe3O4 NP
supercrystals, while the pore size (B10 nm) is slightly smaller
than the diameter of Fe3O4 NPs, probably caused by the
framework shrinkage during etching and/or post drying

500 °C Acid Precursor

Fe3O4 NP supercrystals

(111)

(200)

(311)

Fe3O4 NP supercrystals

Carbon frameworks

(420)

Carbonized NP Mesoporous carbon

Supercrystals frameworks

Mesoporous carbon

Supercrystals

Interconnected NP

superlattices

1,500

1,000

500

(111)

(220)

(220)
(311)

0.5 1.0 1.5 2.0 2.5
2 _�_ (°)

(011)

(211)

Pore diameter (nm)

0 5 10 15 20 25 30
7
6
5
4
3
2
1
0

0.0 0.2

0.4 0.6 0.8 1.0
P/P0

Figure 1 | Fabrication of three-dimensionally interconnected NP superlattices from mesoporous carbon frameworks. (a) Schematic illustration of the
fabrication procedure (cross-sectional view). (b,c) SEM and HRSEM images of carbonized Fe3O4 NP supercrystals, respectively. Scale bars, 1 mm and
200 nm, respectively. (d) SAXS patterns of carbonized Fe3O4 NP supercrystals and mesoporous carbon frameworks, respectively. (e,f) TEM images of
mesoporous carbon frameworks with different lattice projections. Scale bars, 50 and 20 nm, respectively. The inset in (e) is a low-magnification SEM image
of mesoporous carbon frameworks. Scale bar, 1 mm. The red arrows in (f) indicate the interconnected windows. (g) N2 adsorption–desorption isotherms and
the corresponding pore size distribution (inset) of mesoporous carbon frameworks. The red arrow indicates the small pores corresponding to the
interconnected windows observed in (f).

processes. Interestingly, high-resolution TEM (HRTEM) and
Raman spectroscopy suggest that the pore walls of carbon
frameworks are partially graphitic (Supplementary Fig. 4), which
is remarkable considering the low carbonization temperature
(500 �C). We attribute the formation of partially graphitic
frameworks at such a low temperature to the use of Fe3O4
NPs, as transition metals such as Fe and Co and the corresponding metal oxides have been widely used as graphitization
catalysts 25–27 . We surmise that the carbonaceous species arising
from the thermal decomposition of OA molecules leads to the
partial reduction of Fe3O4 NPs (most likely surface atoms) into
metallic iron, which concurrently promotes the graphitization of
the surrounding carbonaceous layers into partially graphitic pore
walls. Despite the thin wall thickness (B2 nm, Supplementary
Fig. 4b), the graphitization degree of carbon frameworks can be
further increased by heat treatment at 1,200 �C in argon,
converting the pore walls into few-layer graphene while
retaining structural ordering (Supplementary Fig. 4c). More
interestingly, tilted experiments along the [011] zone axis in TEM
clearly shows that the adjacent pores are connected through
small windows with dimensions ofB3 nm (Fig. 1f, indicated by
red arrows), leading to a 3D continuous porous structure.
Presumably, the formation of the interconnected windows is
attributed to the slight sintering of neighbouring Fe3O4 NPs
occurring during the carbonization process.

The porous structure of carbon frameworks is further
characterized by N2 adsorption–desorption isotherms, which
show a type-IV curve with a sharp capillary condensation
step occurring in the relative pressure (P/P0) range of 0.8–0.85
(Fig. 1g), a typical feature of mesoporous materials 28 . The
Brunauer–Emmett–Teller surface area and the pore volume are
determined to be B1,500 m 2 g � 1 and B2.5 cm 3 g � 1, respectively.
The pore size distribution curve determined by using the Barrett–
Joyner–Halenda model suggests a bimodal porous structure
(Fig. 1g, inset). The large pores at B10 nm correspond to the
removed Fe3O4 NPs, while the small pores in the range of 2–4 nm
are ascribed to the interconnected windows observed in TEM.

Interconnected NP superlattices. The 3D continuous porosity
and high surface area of carbon frameworks are expected to
facilitate the diffusion of precursors, which is crucial for the
subsequent growth of NP superlattices. As a proof-of-concept
demonstration, tetraethoxysilane is chosen as the precursor for
SiO2 NP superlattices, which are obtained by repeated precursor
infiltration followed by hydrolysis. Remarkably, detailed structural features such as facets and surface terraces originated from
Fe3O4 NP supercrystals are also observed in the product (Fig. 2a,b
and Supplementary Fig. 5), indicating a high degree of 3D NP
ordering. The flexibility offered by the approach and the
robustness of carbon frameworks allow the ready tuning of NP
compositions by simply selecting appropriate precursors. For
example, pure carbon NP superlattices can be realized by
impregnating carbon frameworks with an aqueous solution of
sucrose followed by drying and calcination (Fig. 2c,d). Energydispersive X-ray spectroscopy (EDS) and elemental mapping
confirm that the resulting superlattices are composed of carbon
NPs (Supplementary Fig. 6). To our knowledge, this is the first
preparation of 3D carbon NP superlattices, as monodisperse,
ligand-capped carbon NPs are not yet accessible for selfassembly 29 .

To further illustrate the generality of the approach, we fabricate
superlattices of crystalline metal oxide NPs such as SnO2 by using
SnCl2 as the precursor. Thermal annealing at 350 �C in argon is
carried out to crystallize the embedded NPs. TEM suggests that
the crystallization process does not disrupt the ordered structure

of SnO2 NP superlattices (Fig. 3a and Supplementary Fig. 7),
while SAXS confirms the fcc geometry with a lattice constant of
21.2 nm (Fig. 3b). Powder X-ray diffraction (XRD, Fig. 3c) and
HRTEM (Fig. 3d) establish that the constituent SnO2 NPs possess
a high crystallinity with a tetragonal crystal structure. In addition
to SnO2, other types of metal oxide nanocrystal superlattices can
be prepared in a similar way using metal alkoxides or anhydrous
inorganic salts as precursors. For instance, 3D superlattices
consisting of anatase TiO2 nanocrystals are formed by impregnating carbon frameworks with titanium isopropoxide (TIP)
followed by hydrolysis and thermal annealing (Supplementary
Fig. 8). Moreover, NP superlattices of multicomponent phases,
such as mixed oxides and metal phosphates, are also accessible by
using a pre-mixed precursor (Supplementary Fig. 9). Apparently,
the NP composition can be tailored by changing the ratio of the
two precursors. For example, Ti0.3Sn0.7O2 NP superlattices are
obtained when a mixture of SnCl2 and TIP is employed as the
precursor, and the uniform distribution of Ti and Sn as revealed
by elemental mapping indicates that the homogenous filling of
Ti0.3Sn0.7O2 NPs within the carbon framework (Supplementary
Fig. 9c–e). The ability to fabricate such multiphase NP superlattices is particularly important, as they are generally hard to be
prepared by self-assembly due to the challenge to obtain the
corresponding monodisperse NPs.

Li-ion storage properties. One unique structural feature of NP
superlattices produced by our approach is that the neighbouring
NPs are connected with each other through interconnected
windows within carbon frameworks, as shown in Fig. 3d where
SnO2 NP connections can be clearly observed (indicated by red
arrows). Such connection configuration could substantially
enhance electronic coupling between close-packed NPs 30,31,
which is highly desirable for applications requiring high
electrical conductivity. In comparison to self-assembled NP
superlattices, another important feature of our NP superlattices
is that the constituent NPs are embedded in a continuous and
partially graphitic carbon matrix. Such carbon coating could
further facilitate electron transport and accommodate volume
variations of NPs, which is particularly beneficial for energy
storage devices 32–34 . To demonstrate this, SnO2 NP superlattices
are chosen as anode materials for LIBs, which exhibit excellent
electrochemical performance in terms of cycling stability and rate
capability. The reason to select SnO2 NP superlattices is that SnO2
has attracted increasingly greater attention for Li-ion storage due
to its high theoretical capacity (780 mAh g � 1 ) and environmental
benignity 35–44 .

To investigate their electrochemical performance, LIB anodes
based on SnO2 NP superlattices are cycled on the basis of the
half-cell configuration. The carbon content of SnO2 NP superlattices is B28 wt%, determined from TGA (Supplementary
Fig. 10a). Electron microscopies indicate that SnO2 NP superlattices are uniformly distributed in the electrode with wellretained structural ordering before cycling (Supplementary
Fig. 10b–d). Fig. 4a shows the cyclic voltammograms of SnO2
NP superlattices in the potential range from 3.0 to 0.005 V (versus
Li/Li þ ) at a scan rate of 0.5 mV s � 1 . The irreversible peak at
0.75 V in the first lithiation process is attributed to the reduction
of SnO2 into Sn as well as the formation of a solid electrolyte
interphase (SEI) 39 . The two peaks at 0.7 and 1.3 V in the first
delithiation process correspond to the de-alloying of LixSn and
the partially reversible reaction of SnO2 with Li þ, respectively,
consistent with previous results 37,38 . The cycling stability of SnO2
NP superlattices is studied by galvanostatic charge/discharge. The
first discharge process leads to a capacity of 1,570 mAh g � 1 at a
relatively high current density of 600 mA g � 1, whereas the

Figure 2 | Representative electron microscopy images of various NP superlattices. (a) SEM image of SiO2 NP superlattices, showing the faceted
morphology. Scale bar, 500 nm. (b) HRSEM image of SiO2 NP superlattices, showing the surface terraces and long-range NP ordering. Scale bar, 100 nm.
(c,d) SEM and TEM images of 3D carbon NP superlattices, respectively. Scale bars, 100 and 50 nm, respectively.

(111)

(200)

(220) (311)

1.0 1.5 2.0 2.5

2� (°)

(110)

(101)

(301)

(310)

(200)

(211)

(220)

(210) (002)

(321)
(202)

20 30 40 50 60 70
2 _�_ (°)

Figure 3 | Structural characterization of SnO2 NP superlattices. (a) TEM image of SnO2 NP superlattices. Scale bar, 20 nm. (b,c) SAXS pattern and XRD
pattern of SnO2 NP superlattices, respectively. (d) HRTEM image of SnO2 NP superlattices, showing the high crystallinity of the embedded SnO2 NPs. Scale
bar, 5 nm. The red arrows indicate NP interconnections.

subsequent charge process delivers a capacity of 676 mAh g � 1
(Supplementary Fig. 11). This irreversibility is probably caused by
the SEI formation as well as the lithium insertion into carbon

frameworks 36 . Despite the first-cycle capacity loss commonly
observed for SnO2 anodes 36,37, our SnO2 NP superlattices exhibit
excellent cycling stability, as manifested by the retention of a

1,600

1,400

1,200

1,000

800

600

400

200

Current density = 600 mA g –1

SnO2 NP superlattices

0.5

0.0

–0.5

–1.0

–1.5

–2.0

0.7 V
1.3 V

1st
0.75 V 2nd

3rd

Carbon-coated 13-nm SnO2 NPs

Uncoated 13-nm SnO2 NPs

Carbon frameworks

Discharge
Charge

100

0.0 0.5 1.0 1.5 2.0 2.5 3.0
Voltage (V versus Li/Li + )

1,800
Discharge

1,600 Charge

1,400

1,200

0 50 100 150 200
Cycle number

1,000

800

600

400

200

120 mA g –1

300 mA g –1 300 mA g –1

600 mA g –1

1,200 mA g –1

3,000 mA g –1

0 0 10 20 30 40 50 60
Cycle number

Figure 4 | Electrochemical characterization of SnO2 NP superlattices. (a) Representative cyclic voltammograms at a scan rate of 0.5 mVs � 1 . (b) Cycling
performance at a current density of 600 mA g � 1 and the corresponding Coulombic efficiency. The cycling performance of 13-nm SnO2 NPs with and
without carbon coating as well as carbon frameworks tested under the same conditions are also included for comparison. (c) Rate-capability test at current
densities ranging from 120 to 3,000 mA g � 1 . (d) TEM image and (e,f) the corresponding EDS elemental mapping of SnO2 NP superlattices after 200
cycles, showing the preservation of ordered structure without NP aggregation. Scale bar, 50 nm.

specific capacity of 640 mAh g � 1 after 200 cycles and a stabilized
Coulombic efficiency over 98% from the 10th cycle (Fig. 4b). LIB
anodes based on mesoporous carbon frameworks only are also
tested under the same conditions, which exhibit a stable capacity
at 185 mAh g � 1 after 200 cycles (Fig. 4b, pink curve), indicating
that the high capacity of our SnO2 NP superlattices is primarily
attributed to the embedded SnO2 NPs. To better evaluate the
charge/discharge performance of our SnO2 NP superlattices, we
synthesize colloidal SnO2 NPs with a similar diameter (B 13 nm,
Supplementary Fig. 12) 45, which are cycled under the same
conditions in control experiments. As expected, the uncoated
SnO2 NPs exhibit fast capacity decay within 50 cycles (Fig. 4b,
olive curve), presumably caused by the pulverization and/or
aggregation of SnO2 NPs in the absence of carbon protection. In
comparison, the carbon-coated SnO2 NPs (carbon content:B26
wt%) exhibit better cycling stability, but their capacity rapidly
fades below 100 mAh g � 1 after 50 cycles (Fig. 4b, blue curve).
The cycling performance of our SnO2 NP superlattices is further
evaluated at different current densities. As shown in Fig. 4c, the
electrode is able to deliver a capacity of 300 mAh g � 1 even
at a current density of 3,000 mA g � 1, and a capacity above
850 mAh g � 1 is recovered when the current density is
switched back to 120 mA g � 1 . Such cycling stability and rate
capability outperform those of most SnO2 NP anodes reported
previously 38–44 .

Discussion
The superior battery performance of SnO2 NP superlattices is
attributed to their unique structural characteristics. First, our NP
superlattices exist as micrometer-sized secondary particles
inherited from the original Fe3O4 NP supercrystals, which are
believed to be the ideal architectures for LIBs due to the reduced

interfacial areas between active materials and electrolyte 46,47 .
Second, and perhaps most importantly, the 3D graphitic carbon
frameworks combined with interconnected NPs provide a
continuous electron pathway, facilitating electronic connectivity
within the components of the electrode. Moreover, the inherent
flexibility of carbon frameworks can buffer large volume
expansions of the embedded SnO2 NPs during cycling 35,36,
reducing the strain of the entire electrode.

To investigate the structural evolution of our SnO2 NP
superlattices during cycling, ex situ XRD, SEM, TEM, and EDS
elemental mapping studies are performed after 200 cycles. In
accordance with previous results for SnO2-based anodes 35,43,
XRD confirms the conversion of SnO2 NPs into Sn NPs after
cycling (Supplementary Fig. 13a), while the broad diffraction
peaks imply that the resulting Sn NPs are small in size without
agglomeration. SEM reveals that the secondary particle
morphology is largely preserved after cycling (Supplementary
Fig. 13b,c), while TEM confirms that the Sn-based NPs are
homogeneously confined within carbon frameworks with good
structural ordering without aggregation (Fig. 4d–f and
Supplementary Fig. 13d). These results strongly suggest that our
carbon frameworks effectively suppress the agglomeration and
pulverization of Sn-based NPs, which could explain the superior
cycling stability of our SnO2 NP superlattices. It should also be
noted that the LIB performance could be further improved by
enhancing the electrical conductivity of carbon frameworks
before precursor infiltration, which can be realized by
increasing the graphitization degree via heat treatment as
mentioned above.

In summary, we have developed a new approach to fabricate
three-dimensionally interconnected, strongly coupled NP superlattices with fcc packing symmetry. The generality of the
approach is illustrated by the growth of NP superlattices with

various compositions including oxides, carbon, mixed oxides, and
phosphates. We also demonstrate the successful use of SnO2 NP
superlattices as anode materials for LIBs, and the excellent cycling
stability and rate capability are attributed to their unique
architectures (that is, 3D continuous and graphitic carbon
coating, NP interconnections, and supercrystal morphology).
Given the fact that the fabrication procedure does not require the
presynthesis of monodisperse NP building blocks, we anticipate
that in the future a wider range of materials can be prepared as
interconnected NP superlattices, which can find various applications in electronics, catalysis, and energy storage.

Methods
Materials. Oleic acid (OA, 90%), squalane, and 1-octadecene (ODE, 90%) were
purchased from Aldrich. Sodium oleate was obtained from TCI. Iron chloride
hexahydrate (FeCl3.6H2O), titanium isopropoxide (TIP), anhydrous tin chloride
(SnCl2), anhydrous tin tetrachloride (SnCl4) tetraethoxysilane (TEOS), anhydrous
zirconium tetrachloride (ZrCl4), Cetyl trimethyl ammonium bromide (CTAB), and
triethyl phosphate were purchased from Aladdin. All chemicals were used as
received without further purification.

Synthesis and self-assembly of Fe3O4 NPs. Monodisperse, OA-stabilized Fe3O4
NPs with a diameter ofB11 nm were synthesized according to the literature
method 14 . In a typical synthesis, 72 g of iron oleate and 11.4 g of OA were dissolved
in 400 g of ODE in a three-neck flask, and the resulting solution was heated at
320 �C under a N2 atmosphere for 1 h. After cooling down to room temperature,
ethanol and isopropanol were added to precipitate Fe3O4 NPs, and the precipitates
were dispersed in hexane to form a stable colloidal solution with a concentration
ofB10 mg ml � 1 . To self-assemble Fe3O4 NP supercrystals, squalane (B 1 wt%)
was added to the NP solution and the solvent was then allowed to evaporate under
ambient conditions. The complete solvent evaporation yielded faceted Fe3O4 NP
supercrystals with dimensions on the micrometer scale.

Fabrication of three-dimensionally interconnected NP superlattices. The asassembled Fe3O4 NP supercrystals were heated in a quartz tube furnace at 500 �C
for 2 h in argon, converting the surface-coating ligands into carbon. The carbonized NP supercrystals were then refluxed in a HNO3 or HCl solution to remove
the embedded Fe3O4 NPs. The resulting mesoporous carbon frameworks retrieved
by centrifugation were washed with deionized water, yielding black powders after
drying. To grow interconnected NP superlattices, the dried porous carbon frameworks were infiltrated with desired precursors via wet impregnation followed by
hydrolysis and/or thermal annealing.

Preparation of SiO2 NP superlattices. 5 mg of dried mesoporous carbon powder
was dissolved in 1 ml of TEOS. After stirring for 6 h, the precipitated powder
collected by centrifugation was washed with ethanol to remove the excess TEOS.
The hydrolysis of TEOS was induced by the addition of ammonia hydroxide,
leading to the conversion of TEOS into SiO2. This impregnation/washing cycle was
repeated twice in order to completely fill carbon frameworks with SiO2 NPs.

Preparation of carbon NP superlattices. 3D carbon NP superlattices were synthesized using sucrose as the precursor. Briefly, 0.5 g of sucrose and 0.1 g of concentrated H2SO4 were first dissolved in 1 ml of H2O to form a mixture, into which
5 mg of dried mesoporous carbon powder was added with constant stirring. After
2 h, the precipitated powder collected by centrifugation was washed with H2O to
remove the excess sucrose and H2SO4. After that, the product was placed in an
oven at 100 �C for 2 h and then 160 �C for another 2 h. The black powder was then
impregnated in a diluted sucrose solution (0.05 g of sucrose, 0.02 g of H2SO4, and
1 ml of H2O) for 2 h. After washing and drying, the product was heated at 600 �C in
argon for 2 h to carbonize sucrose, yielding 3D carbon NP superlattices.

Preparation of SnO2 NP superlattices. 50 mg of mesoporous carbon powder was
mixed with a SnCl2 solution dissolved in ethanol (0.1 g ml � 1 ) and the resulting
suspension was stirred for 2 h. After that, the black powder was collected by
centrifugation and the precipitated powder was washed with ethanol to remove the
excess SnCl2. The precursor hydrolysis was initiated by the addition of ammonium
hydroxide. This impregnation/washing cycle was repeated twice in order to completely load carbon frameworks with SnO2. Thermal annealing at 350 �C in argon
for 1 h was carried out to crystallize the embedded SnO2 NPs.

Preparation of TiO2 NP superlattices. TIP was used as the precursor for the
fabrication of TiO2 NP superlattices. Briefly, 50 mg of dried mesoporous carbon
powder was mixed with a TIP solution dissolved in isopropanol (0.5 g ml � 1 ). After
stirring for 6 h, the black powder was collected by centrifugation and the

precipitated powder was washed with isopropanol to remove the excess TIP. The
subsequent hydrolysis of TIP was initiated by exposure to air, leading to the
conversion of TIP into TiO2. This impregnation/washing cycle was repeated twice
in order to completely fill carbon frameworks with TiO2. After drying, the product
was heated at 350 �C in argon for 1 h to crystallize the embedded TiO2 NPs.

Preparation of TixSn1 � xO2 NP superlattices. A homogeneous mixture of TIP
and SnCl2 with different molar ratios dissolved in ethanol was used as the precursor for the growth of TixSn1 � xO2 NP superlattices. In a typical synthesis of
Ti0.3Sn0.7O2 NP superlattices, 0.4 mmol of TIP and 0.6 mmol of SnCl2 were dissolved in 1 ml of ethanol to form a homogeneous solution, into which 5 mg of dried
mesoporous carbon powder was added with stirring. After 24 h, the product was
centrifuged and washed with ethanol to remove the excess precursor. The precursor hydrolysis was initiated by the addition of ammonium hydroxide. This
impregnation/washing cycle was repeated twice in order to completely load carbon
frameworks with Ti0.3Sn0.7O2 NPs. After drying, the product was heated at 350 �C
in argon for 1 h to crystallize the embedded Ti0.3Sn0.7O2 NPs.

Preparation of zirconium phosphate NP superlattices. Zirconium phosphate
(ZrP) NP superlattices were prepared by using a homogeneous mixture of ZrCl4
and triethyl phosphate as the precursor. Briefly, 0.45 mmol of ZrCl4 and 0.45 mmol
of triethyl phosphate were mixed in 2 ml of ethanol to form a homogeneous
solution, into which 5 mg of dried mesoporous carbon powder was added with
vigorous stirring for 2 h. After that, the product was centrifuged and washed with
ethanol to remove the excess precursor. The precipitated powder was then put in
an oven at 110 �C for 2 h to induce the formation of ZrP NP superlattices.

Synthesis of 13-nm SnO2 NPs in control experiments. In control experiments,
colloidal SnO2 NPs with a similar size as our SnO2 NP superlattices were synthesized by a literature method 45 . In a typical synthesis, 10 ml of cationic surfactant
(CTAB) solution (0.08 mol l � 1 ) was mixed with 10 ml of NH3.H2O (25% aqueous
solution) to form a homogeneous solution, into which 4.65 g of SnCl4 was added
under vigorous stirring. After 4 h, the product was filtered and washed with
distilled water for several times. The product was then heated at 400 �C in air for
2 h to increase the crystallinity, leading to SnO2 NPs with a mean diameter of
13 nm. Sucrose was used as the carbon source for the formation of carbon-coated
SnO2 nanocomposite. In a typical preparation of SnO2/C nanocomposite with a
carbon content ofB26 wt%, 0.35 g of SnO2 NPs was mixed with a sucrose solution
(0.4 g in 5 ml H2O), and the resulting mixture was pre-heated at 180 �C in an oven
for 3 h. The dried powder was then heated in a quartz furnace at 500 �C for 5 h in
argon, resulting into carbon-coated SnO2 nanocomposite.

Instrumentation. TEM images, HRTEM images, scanning TEM images, elemental
mapping, and EDS spectra were obtained using a Tecnai G2 20 TWIN microscope
operated at 200 kV. SEM images and EDS spectra were recorded using a Zeiss Ultra55 microscope operated at 5 and 10 kV, respectively. XRD was carried out on a
Bruker D4 X-ray diffractometer, while SAXS was performed on a Nanostar U smallangle X-ray scattering system using Cu Ka radiation (40 kV, 35 mA). Nitrogen
adsorption–desorption isotherms were recorded on a Tristar 3000 instrument.
Before measurements, the samples were degassed at 300 �C for 5 h. Raman spectra
were collected at room temperature on an XploRA Raman system. TGA measurements were carried out on a Perkin–Elmer Pyris 1 thermogravimetric analyzer.

Electrochemical measurements. The battery performance was evaluated by
galvanostatic cycling of 2025-type coin cells assembled in an argon-filled glove box,
with SnO2 NP superlattices as the working electrode and lithium foil as the counter
and reference electrode. The electrolyte was a 1.0 M LiPF6 solution in a mixture of
ethylene carbonate, dimethyl carbonate and ethyl methyl carbonate (1:1:1 in
volume), and a polypropylene film (Celgard-2300) was used as the separator. The
working electrodes were prepared by a slurry-coating procedure. The slurry consisted of SnO2 NP superlattices, acetylene black (Super P) and polyvinylidene
fluoride binder with a mass ratio of 7:2:1 dissolved in N-methyl-2-pyrrolidinone.
This slurry was spread on a copper foil, which acted as a current collector. The
electrodes were dried at 90 �C for 4 h in air, and then at 90 �C in vacuum for
another 12 h. Cyclic voltammetry was carried out on an electrochemical workstation (Autolab 204 N), while galvanostatic measurements were performed on a
Neware cell test instrument, which was cycled between 0.005 and 3.00 V (versus Li/
Li þ ) at various current densities.

References
1. Talapin, D. V., Lee, J.-S., Kovalenko, M. V. & Shevchenko, E. V. Prospects of
colloidal nanocrystals for electronic and optoelectronic applications. Chem.
Rev. 110, 389–458 (2010).
2. Panthani, M. G. & Korgel, B. A. Nanocrystals for electronics. Annu. Rev. Chem.
Biomol. Eng. 3, 287–311 (2012).
3. Pileni, M. P. Self-assembly of inorganic nanocrystals: fabrication and collective
intrinsic properties. Acc. Chem. Res. 40, 685–693 (2007).

4. Murray, C. B., Kagan, C. R. & Bawendi, M. G. Synthesis and characterization of
monodisperse nanocrystals and close-packed nanocrystal assemblies. Annu.
Rev. Mater. Sci. 30, 545–610 (2000).
5. Shevchenko, E. V., Talapin, D. V., Kotov, N. A., O’Brien, S. & Murray, C. B.
Structural diversity in binary nanoparticle superlattices. Nature 439, 55–59
(2006).
6. Kalsin, A. M. et al. Electrostatic self-assembly of binary nanoparticle crystals
with a diamond-like lattice. Science 312, 420–424 (2006).
7. Mueggenburg, K. E., Lin, X. M., Goldsmith, R. H. & Jaeger, H. M. Elastic
membranes of close-packed nanoparticle arrays. Nat. Mater. 6, 656–660 (2007).
8. Dong, A. G., Chen, J., Patrick, M. V., Kikkawa, J. M. & Murray, C. B. Binary
nanocrystal superlattice membranes self-assembled at the liquid-air interface.
Nature 466, 474–477 (2010).
9. Wang, T. et al. Self-assembled colloidal superparticles from nanorods. Science
338, 358–363 (2012).
10. Singh, G. et al. Self-assembly of magnetite nanocubes into helical
superstructures. Science 345, 1149–1153 (2014).
11. Dong, A. G., Jiao, Y. C. & Milliron, D. J. Electronically coupled nanocrystal
superlattice films by in situ ligand exchange at the liquid-air interface. ACS
Nano 7, 10978–10984 (2013).
12. Talapin, D. V. & Murray, C. B. PbSe nanocrystal solids for n- and p-channel
thin film field-effect transistors. Science 310, 86–89 (2005).
13. Luther, J. M. et al. Structural, optical and electrical properties of self-assembled
films of PbSe nanocrystals treated with 1,2-ethanedithiol. ACS Nano 2, 271–280
(2008).
14. Park, J. et al. Ultra-large-scale syntheses of monodisperse nanocrystals. Nat.
Mater. 3, 891–895 (2004).
15. Nagaoka, Y., Chen, O., Wang, Z. W. & Cao, Y. C. Structural control of
nanocrystal superlattices using organic guest molecules. J. Am. Chem. Soc. 134,
2868–2871 (2012).
16. Cychosz, K. A. et al. Characterization of the pore structure of threedimensionally ordered mesoporous carbons using high resolution gas sorption.
Langmuir 28, 12647–12654 (2012).
17. Gierszal, K. P. & Jaroniec, M. Carbons with extremely large volume of uniform
mesopores synthesized by carbonization of phenolic resin film formed on
colloidal silica template. J. Am. Chem. Soc. 128, 10026–10027 (2006).
18. Lee, H. I. et al. Spontaneous phase separation mediated synthesis of 3D
mesoporous carbon with controllable cage and window size. Adv. Mater. 23,
2357–2361 (2011).
19. Vu, A. et al. Three-dimensionally ordered mesoporous (3DOm) carbon
materials as electrodes for electrochemical double-layer capacitors with ionic
liquid electrolytes. Chem. Mater. 25, 4137–4148 (2013).
20. Fang, B., Kim, M., Kim, J. H., Lim, S. & Yu, J. Ordered multimodal porous
carbon with hierarchical nanostructure for high Li storage capacity and good
cycling performance. J. Mater. Chem. 20, 10253–10259 (2010).
21. Jun, S. et al. Synthesis of new, nanoporous carbon with hexagonally ordered
mesostructure. J. Am. Chem. Soc. 122, 10712–10713 (2000).
22. Kim, T. W., Park, I. S. & Ryoo, R. A synthetic route to ordered mesoporous
carbon materials with graphitic pore walls. Angew. Chem. Int. Ed. 42,
4375–4379 (2003).
23. Chen, J. et al. Bistable magnetoresistance switching in exchange-coupled
CoFe2O4–Fe3O4 binary nanocrystal superlattices by self-assembly and thermal
annealing. ACS Nano 7, 1478–1486 (2013).
24. Law, M. et al. Structural, optical, and electrical properties of PbSe nanocrystal
solids treated thermally or with simple amines. J. Am. Chem. Soc. 130,
5974–5985 (2008).
25. Zhang, L. S., Li, W., Cui, Z. M. & Song, W. G. Synthesis of porous and graphitic
carbon for electrochemical detection. J. Phys. Chem. C 113, 20594–20598 (2009).
26. Lee, K. T., Ji, X., Rault, M. & Nazar, L. F. Simple synthesis of graphitic ordered
mesoporous carbon materials by a solid-state method using metal
phthalocyanines. Angew. Chem. Int. Ed. 48, 5661–5665 (2009).
27. Gao, W., Wan, Y., Dou, Y. & Zhao, D. Synthesis of partially graphitic ordered
mesoporous carbons with high surface areas. Adv. Energy Mater. 1, 115–123
(2011).
28. Wan, Y. & Zhao, D. Y. On the controllable soft-templating approach to
mesoporous silicates. Chem. Rev. 107, 2821–2860 (2007).
29. Sun, Y. P. et al. Quantum-sized carbon dots for bright and colorful
photoluminescence. J. Am. Chem. Soc. 128, 7756–7757 (2006).
30. Zhang, H. et al. Surfactant ligand removal and rational fabrication of
inorganically connected quantum dots. Nano Lett. 11, 5356–5361 (2011).
31. Boneschanscher, M. P. et al. Long-range orientation and atomic attachment of
nanocrystals in 2D honeycomb superlattices. Science 344, 1377–1380 (2014).
32. Liu, N. et al. A pomegranate-inspired nanoscale design for large-volumechange lithium battery anodes. Nat. Nanotech. 9, 187–192 (2014).
33. Chockla, A. M. et al. Silicon nanowire fabric as a lithium ion battery electrode
material. J. Am. Chem. Soc. 133, 20914–20921 (2011).

34. Jang, B. et al. Direct synthesis of self-assembled ferrite/carbon hybrid
nanosheets for high performance lithium-ion battery anodes. J. Am. Chem. Soc.
134, 15010–15015 (2012).
35. Han, F., Li, W., Li, M. & Lu, A. Fabrication of superior-performance
SnO2@C composites for lithium-ion anodes using tubular mesoporous carbon
with thin carbon walls and high pore volume. J. Mater. Chem. 22, 9645–9651
(2012).
36. Jahel, A., Ghimbeu, C. M., Monconduit, L. & Vix-Guterl, C. Confined
ultrasmall SnO2 particles in micro/mesoporous carbon as an extremely long
cycle-life anode material for Li-ion batteries. Adv. Energy Mater. 4, 1400025
(2014).
37. Liang, J. et al. Bowl-like SnO2@carbon hollow particles as an advanced anode
material for lithium-ion batteries. Angew. Chem. Int. Ed. 53, 1–6 (2014).
38. Lin, J. et al. Graphene nanoribbon and nanostructured SnO2 composite anodes
for lithium Ion batteries. ACS Nano 7, 6001–6006 (2013).
39. Ding, S. J. et al. Formation of SnO2 hollow nanospheres inside mesoporous
silica nanoreactors. J. Am. Chem. Soc. 133, 21–23 (2011).
40. Qiao, H. et al. Sonochemical synthesis of ordered SnO2/CMK-3
nanocomposites and their lithium storage properties. ACS Appl. Mater.
Interfaces 3, 3704–3708 (2011).
41. Lee, K. T., Lytle, J. C., Ergang, N. S., Oh, S. M. & Stein, A. Synthesis and rate
performance of monolithic macroporous carbon electrodes for lithium-ion
secondary batteries. Adv. Funct. Mater. 15, 547–556 (2005).
42. Kim, C., Noh, M., Choi, M., Cho, J. & Park, B. Critical size of a nano SnO2
electrode for Li-secondary battery. Chem. Mater. 17, 3297–3301 (2005).
43. Li, X. et al. Tin oxide with controlled morphology and crystallinity by atomic
layer deposition onto graphene nanosheets for enhanced lithium storage. Adv.
Funct. Mater. 22, 1647–1654 (2012).
44. Zhang, X., Jiang, B., Guo, J., Xie, Y. & Tang, L. Large and stable reversible
lithium-ion storages from mesoporous SnO2 nanosheets with ultralong lifespan
over 1000 cycles. J. Power Sources 268, 365–371 (2014).
45. Wang, Y. D., Ma, C. L., Sun, X. D. & Li, H. D. Preparation and characterization
of SnO2 nanoparticles with a surfactant-mediated method. Nanotechnology 13,
565–569 (2002).
46. Lee, S. H. et al. Self-assembled Fe3O4 nanoparticle clusters as high-performance
anodes for lithium ion batteries via geometric confinement. Nano Lett. 13,
4249–4256 (2013).
47. Oszajca, M. F., Bodnarchuk, M. I. & Kovalenko, M. V. Precisely engineered
colloidal nanoparticles and nanocrystals for Li-ion and Na-ion batteries: model
systems or practical solutions? Chem. Mater. 26, 5422–5432 (2014).

Acknowledgements
D.Y. is grateful for financial support from Natural National Science Foundation of
China (51103026, 51373035, and 51373040), the Shanghai Scientific and Technological
Innovation Project (11JC1400600 and 124119a2400), and International Science and
Technology Cooperation Program of China (2014DFE40130). A.D. acknowledges the
financial support from Fudan University, the ‘1000 Youth Talents’ Plan, Natural National
Science Foundation of China (21373052), and National Basic Research Program of China
(973 program: 2014CB845602).

Author contributions
Y.J. and A.D. conceived and designed the experiments. Y.J. and D.H. performed the
experiments. Y.D., X.Z., G.G. and J.H. collected and analysed the XRD and TEM data.
Y.J., D.Y. and A.D. wrote the paper. All authors discussed the results and commented on
the manuscript.

Additional information
Supplementary Information accompanies this paper at [http://www.nature.com/](http://www.nature.com/naturecommunications)
[naturecommunications](http://www.nature.com/naturecommunications)

Competing financial interests: The authors declare no competing financial interests.

Reprints and permission information is available online at [http://npg.nature.com/](http://npg.nature.com/reprintsandpermissions/)
[reprintsandpermissions/](http://npg.nature.com/reprintsandpermissions/)

How to cite this article: Jiao, Y. et al. Fabrication of three-dimensionally interconnected
nanoparticle superlattices and their lithium-ion storage properties. Nat. Commun. 6:6420
doi: 10.1038/ncomms7420 (2015).

This work is licensed under a Creative Commons Attribution 4.0
International License. The images or other third party material in this
article are included in the article’s Creative Commons license, unless indicated otherwise
in the credit line; if the material is not included under the Creative Commons license,
users will need to obtain permission from the license holder to reproduce the material.
To view a copy of this license, visit [http://creativecommons.org/licenses/by/4.0/](http://creativecommons.org/licenses/by/4.0/)

본문 출처 PDF: PDF 보기 · 다른 PDF는 위 [PDF 업로드]로 올릴 수 있습니다.

활용가능성

이력 0건

아직 편집 이력이 없습니다. 저장할 때마다 이전 내용이 여기에 백업됩니다.