Intrinsic and extrinsic proton conductivity in metal-organic frameworks

Article abstract of DOI:10.1039/c4ra11473f

Metal-organic frameworks (MOFs), a new class of solid-state materials, have recently been investigated as proton conductors, but little is known about their mechanisms. Since most of the conductivities reported so far were measured using powder samples, there is uncertainty as to whether they exhibit intrinsic proton transport through frameworks and/or micropores, or extrinsic transport through interparticle phases. Herein, we re-visit ferrous oxalate dihydrate [Fe(ox)(H₂O)₂] (ox = oxalate anion), which is a dense MOF and recognized as a model system for MOF-based proton conductors. By single-crystal measurements using microelectrodes, we show that protons do not transport through the crystals (<10^-9 S cm^-1), but that the conductivity observed in powder samples originates from interparticle phases. This result raises a question as to how general is this phenomenon? We have comprehensively surveyed the literature on solid-state proton conductors and found that large numbers of MOFs, including [Fe(ox)(H₂O)₂], have a similar activation energy to those of gels and interparticle conductors in classical solid-state materials. This indicates a considerable contribution from interparticle phases towards proton conductivity in MOFs, and single crystal analysis or special methods for powder analysis are clearly necessary to confirm intrinsic conductivity. This journal is

Full text of DOI:10.1039/c4ra11473f

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Intrinsic and extrinsic proton conductivity in
metal-organic frameworks†
Cite this: RSC Adv., 2014, 4, 54382
S. Tominaka*ab and A. K. Cheetham*a
Metal-organic frameworks (MOFs), a new class of solid-state materials, have recently been investigated as
proton conductors, but little is known about their mechanisms. Since most of the conductivities reported so
far were measured using powder samples, there is uncertainty as to whether they exhibit intrinsic proton
transport through frameworks and/or micropores, or extrinsic transport through interparticle phases.
Herein, we re-visit ferrous oxalate dihydrate [Fe(ox)(H
2
O)
2
] (ox ¼ oxalate anion), which is a dense MOF
and recognized as a model system for MOF-based proton conductors. By single-crystal measurements
ꢀ
9
ꢀ1
using microelectrodes, we show that protons do not transport through the crystals (<10 S cm ), but
that the conductivity observed in powder samples originates from interparticle phases. This result raises
a question as to how general is this phenomenon? We have comprehensively surveyed the literature on
2 2
solid-state proton conductors and found that large numbers of MOFs, including [Fe(ox)(H O) ], have a
Received 29th September 2014
Accepted 15th October 2014
similar activation energy to those of gels and interparticle conductors in classical solid-state materials.
This indicates a considerable contribution from interparticle phases towards proton conductivity in
MOFs, and single crystal analysis or special methods for powder analysis are clearly necessary to confirm
intrinsic conductivity.
DOI: 10.1039/c4ra11473f
www.rsc.org/advances
Solid-state ion conductors are of great importance for the
development of fuel cells, batteries and supercapacitors.
The most extensively studied polymer electrolyte, Naon,
was invented more than 50 years ago, but its proton conduction
1
A variety of solid-state materials, such as polymers, inorganic mechanism is still unclear, mainly due to its complicated ionic
10,11
materials, and, recently, metal-organic frameworks (MOFs), channel structure.
Compared with that situation, the
have been explored for this application. MOFs are a new class mechanism in MOFs should be more straightforward, because
of solid-state materials, comprising metal ions and organic their detailed crystal structures can be determined by diffrac-
ligands that are linked to form one-dimensional to three- tion methods. However, even though a number of proton
2
,3
dimensional structures.
Porous MOFs have attracted conductive MOFs have been reported, they need further exper-
considerable attention in view of their potential applications imental and theoretical analysis in order to establish a better
2
in gas sorption, carbon dioxide trapping and catalysis. Since understanding of their conduction mechanisms. In particular,
the pores can trap proton carrier molecules as well as single-crystal conductivity measurements are necessary to
accommodating modications of the walls, these porous clarify whether protons are transported through the lattice and/
4
–6
MOFs have also been investigated as proton conductors,
or micropores or through interparticle phases. Such interpar-
7
,8
9
especially for anhydrous conduction. Dense MOFs, which ticle conduction is oen observed in inorganic proton
do not have pores that can trap proton carriers, are more conductors and can show conductivities of >10 S cm at
analogous to the classical inorganic proton conductors, and room temperature (RT). Thus, it is reasonable to consider
protons are expected to be transported through the crystal whether it contributes to, or even dominates, proton conduc-
ꢀ
3
ꢀ1
12
lattice via defects.
tion in some of the MOFs.
In the present work we revisit a well-known, proton
conductive dense MOF: ferrous oxalate dihydrate [Fe(ox)(H
2 2
O) ]
a
Department of Materials Science and Metallurgy, University of Cambridge, Charles
(
ox ¼ oxalate anion). This dense MOF shows an excellent
Babbage Road, Cambridge CB3 0FS, UK. E-mail: akc30@cam.ac.uk; Fax: +44 (0)
ꢀ3
ꢀ1
ꢂ
conductivity of 1.3 ꢁ 10 S cm (at 25 C and 98% relative
1
223 334567; Tel: +44 (0)1223 767061
13
humidity (RH)), and is regarded as an outstanding example of
b
International Center for Materials Nanoarchitectonics (WPI-MANA), National
Institute for Materials Science (NIMS), Ibaraki 305-0044, Japan. E-mail: TOMINAKA.
Satoshi@nims.go.jp; Tel: +81 (0)29 860 4594
4
MOF-based proton conductors. However, it is unclear whether
protons are transported through the framework or not because
13
†
Electronic supplementary information (ESI) available: Experimental details, the data were obtained from a polycrystalline pellet. Here we
crystallographic analysis, cif les, thermogravimetric analysis, infrared spectra.
CCDC 1010845–1010846 and 1020095–1020096. For ESI and crystallographic
data in CIF or other electronic format see DOI: 10.1039/c4ra11472f
2 2
investigate the conductivity of [Fe(ox)(H O) ] by single crystal
14
methods using microelectrodes, in order to acquire a better
54382 | RSC Adv., 2014, 4, 54382–54387
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understanding of MOF-based proton conductors and this constants obtained by Pawley tting of the powder X-ray data
system in particular.
agree well with the reported values (Fig. S3 and Table S1†).
Fig. 2a shows the single-crystal impedance set-up and
demonstrates its performance for an electrically conductive
Results and discussion
14
tetrathiafulvalene nitrate crystal (Fig. S4†). A tiny deep red
crystal was placed on two gold microelectrodes having a 80 mm
gap (Fig. 2a, inset), and a semicircle, whose diameter can be
assigned to resistance of the material (1.1 kU cm), was obtained.
2 2
We synthesized two types of [Fe(ox)(H O) ] crystals; one
comprises high quality single-crystalline prisms (1) while the
other comprises prisms with a signicant level of stacking faults
1
5,16
(
2) (which is typical for [Fe(ox)(H
2
O)
2
]).
1 was synthesized by
ꢂ
annealing out the stacking faults in 2 at 90 C. The crystal
structures were determined by single crystal X-ray diffraction,
and phase purity was conrmed by powder diffraction (Fig. 1a).
2 2
Both have [Fe(ox)(H O) ] chains along the long axis of the
prisms (i.e., along the b axis, determined by diffraction, Fig. 1b),
though 2 has stacking faults along the c* axis according to the
streaks observed in the reciprocal lattice (inset in Fig. 1a, and
S1†). The different domains are shied by 50% along the a or
b axes (Fig. 1c and d, and S2†). The unit cell is similar in
both cases, except for a 1.2% expansion of the a axis in 2; 1
˚
˚
˚
[
(
(
a ¼ 12.065 (12) A, b ¼ 5.559 (2) A, c ¼ 9.817(9) A, b ¼ 128.03
ꢂ
˚
˚
15) at 293 K] and 2 [a ¼ 12.209 (4) A, b ¼ 5.5526 (7) A, c ¼ 9.865
˚
ꢂ
3) A, b ¼ 128.20 (6) at 297 K]. Compared with the reported data
˚
˚
˚
ꢂ
[
a ¼ 12.011(11) A, b ¼ 5.557(5) A, c ¼ 9.920(9) A, b ¼ 128.53(3)
16
at 293 K], 1 has a 0.4% longer a axis and a 1.0% shorter c axis
Table S1†). These differences are attributable to the low quality
of the single crystal data of 1 due to twinning; the lattice
(
Fig. 2 Single-crystal impedance analyses. (a) Complex-plane imped-
ance data for a tetrathiafulvalene nitrate crystal (42 mm wide ꢁ 4 mm
thick) were collected in order to demonstrate the microelectrode
method. The inset image illustrates the setup, and shows a crystal
contacted with microelectrodes having a 80 mm gap using a pressure
Fig. 1 Crystallographic analyses on the [Fe(ox)(H
PXRD patterns of two [Fe(ox)(H O) ] samples: 1 is single crystalline, and was measured. The data were analyzed using an equivalent circuit
contains stacking faults. The black lines show simulated PXRD consisting of the parallel combination of a resistance and a capaci-
patterns, and the red dotted lines show simulated PXRD patterns with tance. (b) Complex-plane impedance diagrams for a single crystal of 1
10 preferred orientation. The inset image shows the reciprocal lattice at room temperature under different relative humidities. The inset
images along the a* direction. (b–d) Structure images. The FeO
compares resistance vs. relative humidity of 1 (23 mm wide ꢁ 23 mm
octahedra are connected by oxalate along the b axis (b). The unit cell is thick), 2 (26 mm wide ꢁ 26 mm thick) and the blank cell. The
composed of 4 different [Fe(ox)(H O) ] chains (c and d). The stacking measurements were performed along the b axis of the crystal, which
2 2
O) ] samples. (a) sensitive adhesive, PSA. The conductivity along the a axis of the crystal
2
2
2
0
6
2
2
faults in 2 are indicated to be along the c* axis.
was confirmed by X-ray diffraction.
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This conrms the reliability of the conductivity measurements Since the solubility of [Fe(ox)(H
2
O)
2
] is negligible (ESI†), the
using the microelectrode method. Fig. 2b shows impedance protons must be released from the particle surface. This is
1
9
spectra of 1, measured along the b axis of the crystal, i.e., along consistent with the Gouy-Chapman Theory of diffuse layers,
1
1
the [Fe(ox)(H O) ] chains. The resistance is >10 U, and slightly which shows that protons can exist as far away as ꢄ900 nm from
2
2
20
decreases with increasing humidity. As shown in the inset, 1, 2 xed counter anions.
and the blank cell have almost the same resistance, indicating In order to explore the generality of interparticle conduction,
that the conductivity originates from the quartz surface, not we have surveyed the literature on proton conductors more
from [Fe(ox)(H O) ]. In addition, the spectra of 1, 2 and the widely. Fig. 4 compares s vs. E for proton conductivity in MOFs
2
2
0
a
blank cell are similar (Fig. S5†), conrming that the large in comparison with classical solid-state materials (details are
resistances are attributable to the bulk of the materials, not to available in the ESI; Table S2–S4 and Fig. S10†). We can classify
the contact between the materials and the electrodes. The them into two types on the basis of their activation energies:
proton conductivity through the [Fe(ox)(H O) ] framework is type 1 has E < 0.4 eV, and type 2 has E > 0.5 eV. In classical
2
2
a
a
ꢀ
9
ꢀ1
ꢂ
judged to be ꢃ5.4 ꢁ 10 S cm (20 C, 80%RH) for 1. In other solid-state materials, most of the gels, polymer electrolytes and
words, it is an insulator. The discussion on the hydrogen bond interparticle conductors are type 1, while most of anhydrous or
networks also supports the conclusion (Fig. S6†).
We then carried out conductivity measurements of 1 and 2 drous conductor, CsHSO
using powder samples under saturated water vapour (Fig. S7– disorder of SO , and thus protons can move as in a liquid.
From the extensive studies of classical hydrated proton
S cm ); conductors, such as Naon, we can establish some fundamental
microporous conductors are type 2. The exception is the anhy-
4
, which has signicant rotational
21
4
ꢀ
4
ꢀ1
ꢀ12
S9†). The conductivity values are 5.7 ꢁ 10 S cm for 1 and
ꢀ3
ꢀ1
ꢂ
ꢀ1
1
.5 ꢁ 10 S cm for 2 at 20 C (dry samples, <10
16
these are in good agreement with the previous measurements.
aspects of proton conduction in hydrated MOFs. Naon has 18–
˚
The Arrhenius plots (Fig. 3) show that proton conduction of 1 35 A water nanochannels (11 wt% H
2
O) lined by –SO
With dehydration, the channel
increases up to 0.4–0.5 eV due
can be accounted for
by the restriction of the molecular dynamics in nanopores,
3
H groups
ꢂ
˚
10,22
proceeds via a constant mechanism in the range 20–90 C, as at separations of ꢄ8 A.
˚
indicated by the linear relationship (activation energy, E
a
¼ 0.17 becomes narrower (<10 A), and E
a
2
ꢀ1 ꢀ1
17
eV, and pre-exponential factor, s
has similar kinetics in the range 20–60 C (E
0
¼ 1.5 ꢁ 10 S cm
K ). 2 to the inuence of the walls. This high E
a
ꢂ
a
¼ 0.14 eV and s ¼
0
2
ꢀ1 ꢀ1
1
.2 ꢁ 10 S cm
K ), but the gradient decreases at higher because the diffusion of molecules in nanopores becomes
temperatures, indicating that some change has taken place. slower than in the bulk when the pores are smaller than about
This is most likely due to the annealing out of the stacking ten molecular diameters;
23
furthermore, the Grotthuss
faults which takes place in this temperature range. The E
a
17
values are close to those for proton conduction in liquids,
ꢄ0.1 eV. In our opinion, therefore, the protons are transported
E
a
¼
18
through an aqueous phase that is formed between the particles.
Fig. 4 Comparison of proton conductivity in metal-organic frame-
works (MOFs) and classical solid-state materials. The relationship
between the pre-exponential factor s
the activation energy E are plotted. MIL53(M) is [M(OH)(14bdc)], where
bdc is benzenedicarboxylate anion. Im@ refers to the incorporation of
] 1 imidazole. –NH , –OH and –CO H correspond to functional groups
0
of the Arrhenius Equation and
a
Fig. 3 Arrhenius plots of the proton conductivity in [Fe(ox)(H
O)
2 2
3
2
and 2 obtained by powder impedance analyses. The AC impedance of modified bdc molecules. ndc ¼ 1,4-naphthalenedicarboxylate
data were obtained at different temperatures under saturated water anion. btc ¼ 1,3,5-benzenetricarboxylate anion. Tz ¼ 1,2,4-triazole.
vapour.
dtmp ¼ hexamethylenediamine tetramethylenephosphonate anion.
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Mechanism does not operate due to the lack of a long-range
Conclusions
24
hydrogen bonded network. Such inuence of the walls is
also observed in microporous materials; for example, zeolite-A On the basis of experimental evidence and a literature survey, it
+
˚
˚
(
ꢄ4.5 A aperture) containing NH (2.8 A diameter) with 11 is proposed that hydrated interparticle phases make a consid-
4
25
wt% water shows E
a
¼ 0.63 eV.
erable contribution to proton conduction in many MOFs when
Compared with these classical materials, the porous MOFs of measurements are made on pellets. MOFs are quite new
type 1 have exceptionally low E values as if there is no inuence materials in terms of the history of proton conductors, and the
a
of the interior walls. For example, MIL-53 ([M(OH)(bdc)], bdc ¼ next stage in this area is to acquire a better understanding of
˚
˚
8,26
1
,4-benzenedicarboxylate anion) has 8.5 A ꢁ 8.5 A channels,
and MIL-53 functionalized with –CO H groups shows E ¼ 0.21
proton conduction mechanisms in such systems.
2
a
6
eV. [Ca(btcH)(H O)]$H O (btc ¼ 1,3,5-benzenetricarboxylate
2
2
Methods
Synthesis
˚
27,28
29
anion) has <6.4 A water channels,
and shows E ¼ 0.18 eV.
a
[La(dtmpH )]$7H
5
2
O (dtmp ¼ hexamethylenediamine tetrame-
28
˚
thylenephosphonate anion) has water channels of <11.7 A, and Diethyl oxalate (1 mmol; Aldrich, >99%) and iron sulfate hep-
30
shows E
a
¼ 0.28 eV. These low E
a
values in microporous MOF tahydrate (2 mmol; Aldrich, >99%) were dissolved in 20 mL of
16
materials probably indicates the presence of interparticle distilled water. The solution was heated by microwave radia-
ꢂ
2 2
O) ]. tion (Biotage Initiator) at 90 C for 2 h, slowly cooled down to
conduction as found in the case of polycrystalline [Fe(ox)(H
Thus, we would argue that the observed conductivities are room temperate, and then aged overnight. These crystals, 2
probably extrinsic rather than intrinsic.
(yield ¼ 72%, oxalate-base), were further treated in the mother
ꢂ
By contrast, most of the anhydrous porous MOFs containing solution at 90 C for 3 days, 1 (yield ¼ 82%). Yellow prisms (10–
heterocycles (i.e., imidazole and triazole) as guest carriers are in 100 mm thick and 100–1000 mm long) were recovered by vacuum
type 2. For example, MIL-53 containing guest imidazole mole- ltration and rinsed with distilled water.
8
˚
˚
cules (4.3 A ꢁ 3.7 A) has E
a
of 0.9 eV (while free imidazole and
The chemical compositions were determined by elemental
analysis at the Department of Chemistry, University of Cam-
31
imidazole in polymer electrolytes have a lower E
We speculate that the high E
a
of 0.25 eV).
a
of the former points to some bridge (found, calcd), for C H O Fe:C (13.31, 13.35), H (2.20,
2 4 6
inuence from pore walls and thus the conductivities of these 2.22), N (0.00, 0.00) for 1; C (13.34, 13.35), H (2.12, 2.22), N (0.00,
anhydrous MOFs are probably due to intrinsic mobility through 0.00) for 2. The H O content was determined by thermal anal-
2
the micropores. Another example of an intrinsic conductor is ysis (Fig. S11†) (found, calcd): 1 (21.4, 20.0), 2 (20.4, 20.0).
the anhydrous, dense MOF, [Zr(HPO
4
)
2
(Tz)
2
] (Tz ¼ 1,2,4-tri-
azole), whose conductivity was the same for both single crystal
Structure determinations
32
and powder samples. In the light of the present work, we
regard this material as an analogue of KH PO and the high E
is probably due to the unsuitable hydrogen bond lengths.
The crystal structures were determined at room temperature by
single crystal diffractometry using an Oxford Diffraction Gemini
A Ultra X-ray diffractometer with Mo Ka radiation. The structure
was solved by direct methods and then rened by the least
2
4
a
33
In order to be certain of obtaining reliable data, we propose
that one should measure the conductivity using single crystal
samples. Whenever powder samples are used, it is hard to
remove the inuence of interparticle phases, especially in
hydrated conductors. In addition, there is further concern
about the reliability with powders due to the possibility of phase
transitions associated with the pelletizing pressure. The impact
of the pelletizing process is oen underestimated, but can be
39
squares methods using the SHELX program within the Olex2
40
interface. All non-hydrogen atoms were rened anisotropically
and hydrogen atoms were rened isotropically. Hydrogen atoms
on water molecules were then added to calculated positions.
Visualization of structures was carried out using the VESTA
41
program. CIF les are available (CCDC 1010845-1010846).
The phase purity was conrmed by powder X-ray diffraction
34
very important in yielding reproducible results. The pressure
needed for making a stable pellet form of a MOF, e.g., 0.2–0.5
GPa in our own experience, might be too high for some MOFs to
retain the original structure of interest. For example, MIL-53(Cr)
(PXRD). The PXRD data were collected using a Bruker-AXS D8
diffractometer with Cu Ka radiation in the Bragg–Brentano
geometry. The patterns were analyzed by the Pawley method
42
using the GSAS-II soware, and plotted with zero-shi
35
undergoes a phase transition at around 55 MPa, zinc imida-
correction and background subtraction. The PXRD patterns
3
6
zolate undergoes phase transition in the range 0.54–0.85 GPa,
43
were simulated using the Mercury program with the March-
37
[Zn
4 3
O(14bdc) ] (MOF-5) is amorphized at 3.5 MPa, and an
44
Dollase parameter of 3.0 for modeling the inuence of
preferred orientation. For the simulations, the high quality
crystallographic structure reported by Echigo et al. was used for
erbium formate framework undergoes a phase transition at 0.6
38
GPa. It is better not to pelletize porous MOFs but to prepare
8
the sample using milder methods, as was done for Im@MIL53;
16
1,
and the crystallographic information obtained by single
PXRD patterns should also be checked aer the measurements.
We conrmed that 0.5 GPa does not change the crystal structure
crystal diffraction in the present work was used for 2.
2 2
of the dense framework of [Fe(ox)(H O) ]. We believe that the
difference between our powder results and the previously
Single-crystal impedance measurements
6
ꢀ1 ꢀ1
reported ones (E
a
¼ 0.37 eV and s
0
¼ 10 S cm
K
at 98% The single-crystal impedance measurements are carried out by
13
RH) is due to the pelletization procedure or the RH difference. mounting the crystal on a quartz chip having microelectrodes
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with a 80 mm gap (Fig. 2a, inset). Humidity was controlled by
Paper
14
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1
1
1
1
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ꢂ
vapour. Aer equilibration at 90 C for a few days, impedance
ꢂ
ꢂ
ꢂ
spectra were collected at every 10 C from 90 C to 20 C with an
equilibration time of >8 h. The reproducibility was conrmed
using different pellets prepared in the same way.
2
3
Infrared spectroscopy
4
Fourier-transform infrared spectroscopy (FTIR) was carried out 17 M. Eikerling and A. A. Kornyshev, J. Electroanal. Chem., 2001,
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ꢀ
1
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1
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Acknowledgements
This work was supported by an Advanced Investigator Award to
AKC from the European Research Council (ERC) and by the 20 S. Tominaka, C.-W. Wu, K. Kuroda and T. Osaka, J. Power
World Premier International Research Center Initiative on
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“
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