DOI: 10.1002/chem.201801099

Communication

& Energy storage | Hot Paper |

Ordered Mesoporous Titania/Carbon Hybrid Monoliths for
Lithium-ion Battery Anodes with High Areal and Volumetric
Capacity
Tobias S. Dçrr,[a, b] Simon Fleischmann,[a, b] Marco Zeiger,[a, b] Ingrid Grobelsek,[a]
Peter W. de Oliveira,[a] and Volker Presser*[a, b]
Abstract: Free-standing, binder-free, and conductive additive-free mesoporous titanium dioxide/carbon hybrid electrodes were prepared from co-assembly of a poly(isoprene)-block-poly(styrene)-block-poly(ethylene oxide)
block copolymer and a titanium alkoxide. By tailoring an
optimized morphology, we prepared macroscopic mechanically stable 300 mm thick monoliths that were directly employed as lithium-ion battery electrodes. High areal
mass loading of up to 26.4 mg cm@2 and a high bulk density of 0.88 g cm@3 were obtained. This resulted in a highly
increased volumetric capacity of 155 mAh cm@3, compared
to cast thin film electrodes. Further, the areal capacity of
4.5 mAh cm@2 represented a 9-fold increase compared to
conventionally cast electrodes. These attractive performance metrics are related to the superior electrolyte
transport and shortened diffusion lengths provided by the
interconnected mesoporous nature of the monolith material, assuring superior rate handling, even at high cycling
rates.

In the growing field of long-lasting and stable performing
mobile electronics and electrified transportation, lithium-ion
batteries (LIBs) find widespread use in rechargeable energy
storage systems.[1] LIBs commonly employ a graphite anode
and a lithium metal oxide cathode.[2] Energy is stored by Faradaic intercalation reactions of lithium ions into the layered
structure of the graphite anode during the charging step. The
relatively slow lithium ion diffusion into the host material limits
the power performance of the device.[3] However, to meet
future demands, a combination of high overall energy density
with fast charge/discharge rates is required. To that end, transition metal oxides like TiO2,[4] Li4Ti5O12,[5] and Nb2O5[6] emerged
[a] T. S. Dçrr, S. Fleischmann, Dr. M. Zeiger, Dr. I. Grobelsek,
Dr. P. W. de Oliveira, Prof. Dr. V. Presser
INM–Leibniz Institute for New Materials
Campus D2 2, 66123 Saarbrecken (Germany)
E-mail: volker.presser@leibniz-inm.de
[b] T. S. Dçrr, S. Fleischmann, Dr. M. Zeiger, Prof. Dr. V. Presser
Saarland University, Campus D2 2, 66123 Saarbrecken (Germany)
Supporting information for the manuscript, and the ORCID identification
number(s) for the author(s) of this article can be found under https://
doi.org/10.1002/chem.201801099.
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as attractive alternatives to graphite anode materials, because
of their superior lithium intercalation kinetics and safer operation potential, above the lowest unoccupied molecular orbital
of the carbonate electrolytes.[7] The insufficient electrical conductivity of such oxides is typically addressed by admixing of
5–20 mass % carbon conductive additive, forming an electrode
with 5–10 mass % polymer binder to achieve mechanical stability. The high total amount of electrochemically inactive material, up to 30 mass %, drastically decreases the specific capacity
per mass of the final electrode, which is why binder-free and
free-standing electrodes can offer advantages.[8] This is even
more pronounced for the volumetric and areal capacities since
the mixing process with a conductive additive and polymer
binder yields comparably low electrode densities. High electrode density and areal loading of active material are considered crucial for a transfer to practical electrochemical energy
storage applications.[9] While many studies report outstanding
gravimetric performance values for thin film electrodes with
loadings often below 1 mg cm@2, the mass of current collectors,
separator, electrolyte, and housing are often neglected. Considering the added mass of these device components, thin and/or
hollow electrodes become far less attractive since only areal
mass loadings above 10 mg cm@2 are attractive for actual applications.[9, 10] However, increasing the thickness and mass loading of electrodes often substantially decreases the electrode
performance because of limitations posed by electrolyte transport and increased overpotentials, if no well-designed 3D electrode architecture is applied.[10, 11]
Ordered mesoporous metal oxides have been shown to
offer attractive rate handling characteristics due to the beneficial electrolyte transport shortened diffusion paths in their
mesopore network.[12] In a study of Jiao et al., mesoporous lithium manganese oxide was synthesized by a hard-templating
approach.[13] The obtained material possessed a uniform pore
size distribution around 5 nm with a low wall thickness of 7–
15 nm, leading to attractive rate capability as LIB cathodes
when admixed with a binder and a conductive additive.[13] Lee
et al. demonstrated an approach of direct hybridization of
mesoporous titanium oxide with carbon by co-assembly of a
diblock copolymer structure directing agent for uniform mesopores up to 7 nm and higher specific surface areas
(> 100 m2 g@1).[14] They further demonstrate that preservation
of the polymeric template as a thin carbon shell around the
electrode core material framework (e.g., TiO2, Li4Ti5O12) increases the conductivity and maintains the 3D network at the same

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time.[3, 15] Thin films of this material, cast on a current collector
using a polymer binder, showed good rate handling as a LIB
anode material.
Though the described materials showed attractive gravimetric performances when cast as thin film electrodes, the resulting low electrode density and limited mass loading yield low
volumetric and areal capacities. In this study, we report for the
first time on free-standing, binder and conductive additive-free
mesoporous TiO2/carbon hybrid monoliths (TiO2/C) that form
electrodes without further preparation processes. By creating
an optimized pore structure, we obtained mechanically stable
300 mm thick electrodes with unprecedentedly high areal mass
loadings of up to 26.4 mg cm@2. Capitalizing on the high density, mass loading, and additive-free character of these hybrid
electrodes, improved areal and volumetric capacities were obtained, compared to electrodes prepared using conventional
mixing of mesoporous TiO2 (gyroidal TiO2), conductive additive,
and binder. The synthesis can easily be adapted to other Faradaic materials, underlining the high suitability of our one-pot
approach to obtain well-designed 3D electrode structures feasible for practical applications.
Our approach for the synthesis of mesoporous structured
TiO2/carbon hybrids and gyroidal TiO2 by macromolecular co-

assembly of a poly(isoprene)-block-poly(styrene)-block-poly(ethylene oxide) (ISO) and a titanium alkoxide is illustrated in
Figure 1 A. Further details on the synthesis (including materials
and methods) are found in the Supporting Information. The
block copolymer (BCP) template was synthesized by living
anionic polymerization and comprised 30, 60, and 10 vol %
polyisoprene (PI), polystyrene (PS), and polyethylene oxide
(PEO), respectively, with a narrow polydispersity of 1.09 and an
overall molecular weight of 61,000 g mol@1. Through evaporation-induced self-assembly (EISA), microphase separation
occurs due to the internal repulsion of the unique polymer
blocks into pure phase domains; of those the PEO is exclusively swelled by the hydrophilic metal oxide precursor sol.[16]
By changing the conditions during the pyrolysis, we could
selectively choose to 1) preserve the organic polymer template
as a carbon shell around the TiO2 core material to obtain monoliths (500 8C/700 8C, argon atmosphere; samples TiO2/C-500
and TiO2/C-700) or 2) completely remove the BCP (400 8C, air,
sample gyroidal TiO2), resulting in mesoporous TiO2 particles.
Although the ISO was initially designed to preferably direct
into gyroidal morphology (as it was indeed done for the gyroidal TiO2), we found that for free-standing electrodes comprising a conductive carbon shell, a mixture of lamellar and gyroi-

Figure 1. (A) Schematic of the TiO2/C and gyroidal TiO2 preparation via evaporation induced self-assembly (EISA) of the ISO BCP and Ti(iProp)4. Variation of the
pyrolysis atmosphere can selectively preserve a conductive carbon shell. Photographs of the freestanding TiO2/C-500 and TiO2/C-700 electrodes are shown.
(B–D) Scanning electron micrographs and (E-G) transmission electron micrographs of TiO2/C-500, TiO2/C-700, and gyroidal TiO2.
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dal material provides much higher macroscopic mechanical
stability. Earlier studies have highlighted the transition and the
coexistence of a lamellar and a gyroidal phase, designated as
hexagonally perforated lamellar (HPL).[17] We assume that the
complexity and internal stress of pure gyroidal morphology,
with its high amount of grain boundaries, are adverse for a
free-standing electrode, although the carbon shell was expected to support the networked TiO2. Therefore, we have adjusted
the swelling of the PEO phase by adding adequate amounts of
the metal organic precursor, to yield a structure combining the
stability of the less complex lamellar morphology with partially
mesoporous character in the individual lamella. Such a structure was not possible to synthesize without the supporting
carbon shell since the carbon condensates in the deeper
region of the lamellae and thus allows them to be covering
the particle surface. In the case of complete thermal removal
of the BCP, the TiO2 lamellae simply collapse, ultimately leading
to a powder sample. Accordingly, we used a mesoporous alternating gyroidal TiO2 to compare the electrochemical characterization of TiO2 with or without hybridized carbon.
Electron micrographs of the pyrolyzed material show the primary lamellae order of the free-standing TiO2/C hybrids with
partial additional secondary interlamellar porosity; the latter is
very similar for TiO2/C-500 and TiO2/C-700 (Figure 1 B–G). The
average interlamellar distance is about 40 nm, as seen from
transmission electron micrographs. This value is identical for
TiO2/C-500 and TiO2/C-700, thereby suggesting no significant
difference in the carbon contents. All samples are composed
of randomly oriented nanometer-sized TiO2 domains, covered
by a shell of disordered carbon in case of samples pyrolyzed in
argon. We assume that the combination of the polycrystalline

character of TiO2 in combination with the high structural stress
of the gyroidal morphology causes insufficient macroscopic
stability of pure gyroidal samples, independent of the presence
of an additional supporting carbon shell. The elemental analysis by energy dispersive X-ray spectroscopy (EDX) is found in
Supporting Information, Table S1.
Results of the small-angle X-ray scattering (SAXS) characterization for the ISO/precursor, TiO2/C hybrids, and gyroidal TiO2
are depicted in Figure 2 A. From the first diffraction (q*), allowed q/q* ratios for
= 1, 2, 3, …) and alternating
pffiffilamellar
ffi pffiffiffi p(q/q*
ffiffiffi
gyroidal (q/q* = 1, 3, 4, 5,…) morphology are calculated
and given above the corresponding pattern.[18] Based on our
strategy, sufficient macroscopic stability is achieved for a lamellar ordered sample, preserving partial porosity inside the
unique lamellae. The result is an ISO/precursor (TiO2/C) hybrid
comprising lamellar and alternating gyroidal scattering peaks
in first and second order well persevered during pyrolysis.[17a]
The periodicity (dspacing) decreases during the thermal treatment
from 44.8 nm of the ISO/precursor (TiO2/C) hybrid to 33.4 and
32.2 nm, for TiO2/C-500 and TiO2/C-700, respectively, evidenced
by a shift of q* towards higher q-values. Very similar behavior
is observed for just gyroidal TiO2, where the ISO/precursor
hybrid exhibits a unique scattering, further identified as typical
gyroidal diffractions. As expected, a smaller dspacing of 38.7 nm
for the ISO/precursor (gyroidal TiO2) hybrid was found, and we
observed a further decrease to 29.8 nm due to the thermal
shrinking during calcination.[17a]
The mesoporous morphology is additionally studied by nitrogen gas sorption (Supporting Information, Figure S1 A). Both
TiO2/C hybrids show identical hysteresis shape at high relative
pressures, indicating a similar level of mesoporosity in good

Figure 2. (A) Small-angle X-ray scattering pattern of the polymer/precursor, the pyrolyzed hybrid materials TiO2/C-500 and TiO2/C-700, and gyroidal TiO2, including observable ideal peak positions for lamellar (LAM) and alternating gyroidal (GA) morphology. (B) X-ray diffraction pattern (wide angle) and (C) Raman
spectra of TiO2/C-500, TiO2/C-700, and gyroidal TiO2. Inset: Typical B1G, A1/G, and EG bands of crystalline anatase TiO2 from the gyroidal TiO2 sample.
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agreement with SAXS, SEM, and HRTEM. We determined a
total pore volume of 0.12 cm3 g@1 for TiO2/C-500 and of
0.16 cm3 g@1 for TiO2/C-700. Increasing pyrolysis temperature
favors the formation of micropores, observed by the curve progression at low pressure for TiO2/C-700. The BET specific surface area (SSA) was very similar for all samples: 73 m2 g@1 for
TiO2/C-500, 107 m2 g@1 for TiO2/C-700, and 93 m2 g@1 for gyroidal TiO2. The latter has a distinct higher volume of mesopores
(total pore volume: 0.39 cm3 g@1) due to the lack of remaining
carbon and the ordered nature of the pores of the gyroid morphology.
The formation of crystalline titania is evident from X-ray diffraction (Figure 2 B), identified to be pure nano-crystalline anatase for the TiO2/C-500 and gyroidal TiO2. Higher process temperatures favor the formation of rutile in TiO2/C-700, calculated
to be 12 % from Rietveld refinement (Supporting Information
Table S2 and Figure S2). The lack of graphite-related reflection
peaks shows the disordered nature of the carbon phase for
TiO2/C-500 and TiO2/C-700. For such an amorphous carbon, the
typical D-band (prohibited in perfectly crystalline graphite) and
G-band (in-plane bond-stretching of sp2-hybridised carbon
rings) are observed in Raman spectra at 1370 and 1590 cm@1,
respectively (Figure 2 C).[19] An intense and sharp signal at
144 cm@1, assigned to the E1g mode of anatase, evidences the
presence of the TiO2 core material in TiO2/C samples. Strong
fluorescence masks further modes of titania phases, such as
B1g (394 cm@1), Eg (636 cm@1), and the combined B1g/A1g (centered at 514 cm@1). For the gyroidal TiO2, without such a
carbon shell, several anatase modes are clearly observable, as
can be seen in the inset of Figure 2 C.
The total amount of amorphous carbon in the TiO2/C samples was determined by thermogravimetric analysis (TGA) conducted in synthetic air (Supporting Information, Figure S1 B).
Above 300 8C, oxidation takes place and is almost completed
at 500 8C. Thereby, we can identify a total amount of carbon of
ca. 15 mass % for TiO2/C-500 and TiO2/C-700. There was no
noteworthy mass loss for the reference gyroidal TiO2, since the
polymer template was already completely removed during calcination in air at 400 8C as part of the material synthesis.
The electrochemical performance of TiO2/C-500 and TiO2/C700 free-standing monolith electrodes will be evaluated by
comparison with gyroidal TiO2 particles admixed with conductive additive and binder to form conventional electrodes. To
allow for an investigation on the monolith structure on the
performance, conventional electrodes will be prepared from
gyroidal TiO2 as thin films (60 mm) and in the same thickness as
the monolith electrodes (300 mm). The electrodes’ bulk densities and areal mass loadings are given in Supporting Information, Table S3.
Cyclic voltammograms of the samples are recorded at a scan
rate of 0.1 mV s@1 (Figure 3 A) for a qualitative assessment of
the electrode performances versus a lithium counter electrode
in 1 m LiClO4 ethylene carbonate/dimethyl carbonate electrolyte. The onset of the main lithiation peaks can be observed
for all electrodes at about 1.7 V versus Li + /Li, whereas delithiation occurs starting at around 1.9 V versus Li + /Li. This reaction
is associated with the lithiation of anatase with about 0.55 Li
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per TiO2.[20] Peak-stretching can be observed for the monolith
electrodes, related to mass transport limitations in the thick
electrodes.[11] Besides the main intercalation peaks, two small
pairs of peaks can be observed around 1.5 V versus Li + /Li,
which are much more significant for the two monolith electrodes than for the gyroidal TiO2 samples. These intercalation
peaks represent further lithiation/delithiation of the material
up to 1 Li per TiO2, which has been previously reported for
anatase TiO2 with nanoscopic domain sizes.[20] This behavior is
only found for anatase TiO2, which is why the peak duplet is
more pronounced for TiO2/C-500 than for rutile-containing
TiO2/C-700.
Quantitative analysis of the gravimetric and volumetric electrochemical performance and rate handling was conducted by
galvanostatic cycling at varying specific currents (Figure 3 B–C),
the charge/discharge profiles including the first cycle are
found in Supporting Information, Figure S3 and Figure S4 A–D.
At a low rate, TiO2/C-500 electrode showed the highest capacity of 155 mAh cm@3, whereas both cast electrodes exhibited
around 90 mAh cm@3. This is due to the higher density of the
monolith materials, since it can be expected that few limitations are posed by diffusion or mass transport at very low
rates. At increasing rates, up to 1 A g@1, the 300 mm monolith
hybrid electrodes still exhibit a comparable capacity to the
60 mm gyroidal TiO2, with a retention of about 60 mA h cm@3.
For comparison, the cast 300 mm gyroidal TiO2 electrode dropped significantly to around 15 mAh cm@3 at this rate. The much
higher capacity retention of the monolith electrodes is a consequence of the improved electrolyte mass transport within
the interconnected mesoporous network with the carbon shell
as electron pathways in the TiO2/C. In the particle-based gyroidal TiO2 electrode, depletion of lithium ions leads to the capacity drop; only a small part of the electrode volume partakes
in the intercalation reaction. The stability of the electrodes is
tested by prolonged cycling at a low rate of 0.025 A g@1 (Figure 3 C). It can be observed that the hybrid monolith electrodes exhibit a very stable cycling behavior over 100 charge/discharge cycles, whereas the gyroidal TiO2 thin film electrode
shows a slight drop of 9 % over 100 cycles. High stability of
TiO2/C hybrid samples can be related to the interconnected
network, where titania domains are engulfed in a carbon shell,
effectively preventing disintegration during cycling.
The advantages of the hybrid monolith electrodes become
even more apparent when considering the areal performance
(Figure 4 A). The maximum areal capacity of around
4.5 mAh cm@2 of the TiO2/C-500 electrode represents a 9-fold
increase compared to the 60 mm gyroidal TiO2 electrode, and
an almost two-fold increase compared to the corresponding
300 mm gyroidal TiO2 electrode with the same thickness. Even
at a high rate of 2 A g@1, the TiO2/C-700 electrode still shows
twice the area-normalized performance compared to the thin
film electrode.
For better comparability with literature, galvanostatic cycling
was carried out by normalizing to the cycling current to the
electrode masses. However, when comparing different electrodes that are optimized for density and mass loading, normalization to the electrode area can be useful (Figure 4 B). In a re-

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Figure 3. (A) Cyclic voltammograms of TiO2/C-500, TiO2/C-700, and gyroidal TiO2 electrodes at 0.1 mV s@1 in a voltage window of 1.0–3.0 V vs. Li + /Li. (B) Gravimetric discharge capacity with variable rates from 0.025 to 5 A g@1 and (C) volumetric discharge capacities including Coulombic efficiency values over
100 cycles.

Figure 4. Areal capacity at (A) different discharge rates and (B) at different areal current densities for TiO2/C-500, TiO2/C-700, and gyroidal TiO2 casted electrodes with 60 mm and 300 mm thickness. Data for Sun et al. are from Ref. [10].

cently reported study by Sun et al.,[10] it was demonstrated
that a highly interconnected network of holey graphene, hybridized with niobia showed remarkable performance at high
areal mass loadings (11 mg cm@2). When comparing to the data
of this material (replotted; green symbols), the TiO2/C monolith
electrodes exhibit similar performance. This shows that our
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first report on a one-pot synthesis for free-standing electrodes
with high mass loading shows attractive performance, comparable to the state-of-the-art in thick electrodes.
In conclusion, mesoporous, free-standing, binder-free, and
conductive additive-free TiO2/C hybrid monolith electrodes
were prepared for the first time by a one-pot synthesis

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method and studied with a focus on their areal and volumetric
electrochemical capacity. It was found that the carbon shell
supported, lamellar TiO2 microstructure is required for sufficient macroscopic stability, concurrently ensuring sufficient
conductivity. The use of TiO2/C hybrid monolith electrodes as
LIB anodes allowed the operation of ultrathick (300 mm) and
dense (0.88 g cm@3) electrode discs at high rates. This mass
loading is at least an order of magnitude above commonly investigated LIB electrode materials. The increased electrode
density and mass loading show up to 60 % higher volumetric
and areal capacities than cast, mesoporous TiO2 particle-based
electrodes, while still exhibiting high rate handling performance. We believe that our results are highly interesting for
industrial applications, where dense materials are required for
downsizing, and high mass loadings are desired to compensate for the additional mass of cell components. The model
system comprising TiO2 as intercalation material can easily be
exchanged in future studies to synthesize other metal oxidebased high-performance electrode materials.

Acknowledgements
The authors thank Prof. Eduard Arzt (INM) for his continuing
support. This project was supported by the INM FOCUS project
funding (POLION). We thank Benjamin Krener for gas sorption
measurements, Ha Rimbach for GPC measurements, Peng
Zhang for small-angle X-ray scattering data, and Eunho Lim for
scientific discussions (all at INM).

Conflict of interest
The authors declare no conflict of interest.
Keywords: energy storage · hybrid materials · lithium-ion
battery · titanium dioxide
[1] a) G. E. Blomgren, J. Electrochem. Soc. 2017, 164, A5019 – A5025; b) B.
Scrosati, J. Hassoun, Y. K. Sun, Energy Environmental Sci. 2011, 4, 3287 –
3295; c) M. S. Whittingham, Chem. Rev. 2004, 104, 4271 – 4302.

Chem. Eur. J. 2018, 24, 6358 – 6363

www.chemeurj.org

[2] a) K. Mizushima, P. C. Jones, P. J. Wiseman, J. B. Goodenough, Mater. Res.
Bull. 1980, 15, 783 – 789; b) J. B. Goodenough, K. S. Park, J. Am. Chem.
Soc. 2013, 135, 1167 – 1176.
[3] E. Kang, Y. S. Jung, G.-H. Kim, J. Chun, U. Wiesner, A. C. Dillon, J. K. Kim,
J. Lee, Adv. Funct. Mater. 2011, 21, 4349 – 4357.
[4] S. Y. Huang, L. Kavan, I. Exnar, M. Gratzel, J. Electrochem. Soc. 1995, 142,
L142 – L144.
[5] M. Widmaier, N. J-ckel, M. Zeiger, M. Abuzarli, C. Engel, L. Bommer, V.
Presser, Electrochim. Acta 2017, 247, 1006 – 1018.
[6] A. Tolosa, B. Krener, S. Fleischmann, N. J-ckel, M. Zeiger, M. Aslan, I. Grobelsek, V. Presser, J. Mater. Chem. A 2016, 4, 16003 – 16016.
[7] A. Tolosa, S. Fleischmann, I. Grobelsek, A. Quade, E. Lim, V. Presser,
ChemSusChem 2018, 11, 159 – 170.
[8] a) L. David, G. Singh, J. Phys. Chem. C 2014, 118, 28401 – 28408; b) L.
David, R. Bhandavat, U. Barrera, G. Singh, Nat. Commun. 2016, 7, 10998;
c) X. Zhao, C. M. Hayner, M. C. Kung, H. H. Kung, ACS Nano 2011, 5,
8739 – 8749.
[9] Y. Gogotsi, P. Simon, Science 2011, 334, 917 – 918.
[10] H. Sun, L. Mei, J. Liang, Z. Zhao, C. Lee, H. Fei, M. Ding, J. Lau, M. Li, C.
Wang, X. Xu, G. Hao, B. Papandrea, I. Shakir, B. Dunn, Y. Huang, X. Duan,
Science 2017, 356, 599.
[11] M. Singh, J. Kaiser, H. Hahn, J. Electrochem. Soc. 2015, 162, A1196 –
A1201.
[12] a) M. G. Fischer, X. Hua, B. D. Wilts, I. Gunkel, T. M. Bennett, U. Steiner,
ACS Appl. Mater. Interfaces 2017, 9, 22388 – 22397; b) N. Li, G. Liu, C.
Zhen, F. Li, L. Zhang, H.-M. Cheng, Adv. Funct. Mater. 2011, 21, 1717 –
1722; c) S. Fleischmann, D. Leistenschneider, V. Lemkova, B. Krener, M.
Zeiger, L. Borchardt, V. Presser, Chem. Mater. 2017, 29, 8653 – 8662; d) E.
Lim, C. Jo, J. Lee, Nanoscale 2016, 8, 7827 – 7833.
[13] F. Jiao, J. L. Bao, A. H. Hill, P. G. Bruce, Angew. Chem. Int. Ed. 2008, 47,
9711 – 9716; Angew. Chem. 2008, 120, 9857 – 9862.
[14] M. Stefik, J. Lee, U. Wiesner, Chem. Commun. 2009, 2532 – 2534.
[15] J. Lee, Y. S. Jung, S. C. Warren, M. Kamperman, S. M. Oh, F. J. DiSalvo, U.
Wiesner, Macromol. Chem. Phys. 2011, 212, 383 – 390.
[16] M. Stefik, S. Guldin, S. Vignolini, U. Wiesner, U. Steiner, Chem. Soc. Rev.
2015, 44, 5076 – 5091.
[17] a) I. W. Hamley, V. Castelletto, O. O. Mykhaylyk, Z. Yang, R. P. May, K. S.
Lyakhova, G. J. A. Sevink, A. V. Zvelindovsky, Langmuir 2004, 20, 10785 –
10790; b) J. H. Ahn, W. C. Zin, Macromol. Res. 2003, 11, 152 – 156.
[18] T. H. Epps, E. W. Cochran, C. M. Hardy, T. S. Bailey, R. S. Waletzko, F. S.
Bates, Macromolecules 2004, 37, 7085 – 7088.
[19] J. G. Werner, T. N. Hoheisel, U. Wiesner, ACS Nano 2014, 8, 731 – 743.
[20] M. Wagemaker, W. J. H. Borghols, F. M. Mulder, J. Am. Chem. Soc. 2007,
129, 4323 – 4327.

Manuscript received: March 2, 2018
Accepted manuscript online: March 6, 2018
Version of record online: April 5, 2018

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