Jongkook
Hwang
,
Runyu
Yan
,
Martin
Oschatz
and
Bernhard V. K. J.
Schmidt
*
Department of Colloid Chemistry, Max-Planck Institute of Colloids and Interfaces, Am Mühlenberg 1, 14476 Potsdam, Germany. E-mail: Bernhard.schmidt@mpikg.mpg.de
First published on 23rd October 2018
While downscaling metal–organic frameworks (MOFs) into a nanosize regime is highly relevant to meet their growing demand in various potential applications, a simple synthesis of nano-MOF under ambient conditions still remains a difficult task. Here we report a room temperature synthesis of 3D MOF, [Zn2(bdc)2dabco]n (ZBD) (bdc = benzene-1,4-dicarboxylic acid and dabco = 1,4-diazabicyclo[2.2.2]octane) with controlled polymorphism, size, and morphology by changing the kind and composition of solvents. The solvents function as both templates and crystal modulators. Dimethylformamide (DMF) preferably forms a hexagonal rod MOF (ZBDh) while methanol (MeOH) leads to the formation of a tetragonal plate MOF (ZBDt) via a solvent template effect (i.e., polymorph control). The size and morphology can be further controlled using DMF and MeOH as cosolvents with various volume ratios. DMF and MeOH work competitively, and the solvent with a weaker template effect under the given conditions acts as a crystal modulator that lowers the rate of nucleation and increases the size of the crystals. With an increase of MeOH amount, the morphology changes from 1D rods to 2D plates. Protic MeOH reduces the reactivity of nucleophilic dabco and suppresses crystal growth along Zn-dabco [001], thereby leading to the formation of 2D ZBDt plates. To help understand the fundamental morphology–volumetric capacitance relationships in energy storage devices, the resulting ZBDs are conformally pyrolyzed to hexagonal rod- and tetragonal plate-nanoporous carbons and used as electrodes for supercapacitors. Thanks to a 2D morphology and relatively high packing density, tetragonal plate carbon delivers two times higher volumetric capacitance than hexagonal rod carbon, despite their nearly similar gravimetric capacitances.
To meet the ever-growing demand of MOFs in a myriad of potential applications, considerable research efforts have been devoted to controlling the size and morphology of MOFs in a nanosize regime, which can tailor the physical/chemical properties without changing the chemical compositions.11,12 Structuring MOFs at the nano- and mesoscopic scale has been accomplished by various methodologies11–13 such as microwave assisted synthesis,14 mechanochemical synthesis,15 interfacial assembly,16,17 emulsion templating,18 and coordination modulation.19,20 However, these approaches usually require specific reactors and equipment, or additional additives and modulators that are sometimes restricted to work in a narrow synthesis parameter window only. In addition, less attention has been paid to control framework topology (i.e., polymorphism), although it can potentially provide new opportunities for controlling the intrinsic properties and morphologies of MOFs. In this regard, development of a simple strategy that simultaneously enables control over the size, morphology, and framework topology of MOFs under ambient synthetic conditions is highly relevant.
The solvent is one of the critical parameters in MOF synthesis, as its properties such as polarity and solubility of the building blocks greatly influence the nucleation and growth of MOFs.11 Thus, the solvent can not only act as a reaction medium but also a structure directing agent that affects the structures and topologies of the final MOFs by a solvent template effect.21,22 Considering the fact that the pores of the as-made MOFs are generally filled with the solvents as guest molecules, the solvent template effect can become an effective means for preparing MOFs with the desired framework topology. To date, a few systematic studies about the solvent effect on MOF formation have been reported,22–26 however, most of the studies are usually MOF specific and limited to polymorph control only, and therefore considerably lack the ability to adjust the particle size and morphology (e.g., aspect ratio). While a simple solvent mediated strategy toward a designable nano-MOF is highly desirable, it still remains largely unexplored due to the lack of efficient solvent systems and incomplete understanding of the associated MOF formation mechanisms.
Electric double-layer capacitors (EDLCs) have aroused tremendous research interest for decades because of their fast charging capabilities, long cycle life, high stability and safe operation.27–30 One of the most important components in such devices is the carbonaceous electrode material as its ability for electrosorption of electrolyte ions is crucial for the specific capacitance and thus the energy density.31,32 A wide variety of novel carbonaceous materials such as activated carbon,33 carbide-derived carbon,34 zeolite-templated carbon,35 carbon nanotubes,36 and graphene37 have been extensively investigated as electrodes in EDLCs. The capacitance of carbon-based electrode materials and supercapacitors in general has been effectively enhanced by (i) increasing the specific surface area and tailoring the pore size and distribution,32,38,39 (ii) producing composites with pseudocapacitive materials (e.g., metal oxides and conductive polymers),40 and (iii) doping with heteroatoms or incorporating functional moieties.41
Recently, volumetric performance has become a more and more important criterion that gives a realistic picture of the charge-storage capacity in the limited space of energy storage devices, particularly in next generation portable electronic devices and electric vehicles.38,42,43 Although nanoporous carbons provide a stable and reliable performance, they show a relatively low volumetric performance due to their inherent low density. This in turn leads to the presence of significant “dead volume” in EDLC electrodes, that is, the electrode volume that has to be filled with the electrolyte but is not used for energy storage. One promising strategy to overcome this limitation is the preparation of carbon particles with controlled size and morphology that can maximize the materials' packing density.42 In this respect, MOF-derived carbons (MDCs) are highly advantageous because a wide variety of inherent crystal morphologies of MOFs can be transferred through pseudomorphic transformation.6–10 Moreover, MDCs can have high surface area, adjustable porosity, and can be doped with heteroatoms without post-treatments. These properties are very attractive for the development of high performance EDLC electrodes. Therefore, morphology controlled MDCs are highly relevant to understanding the fundamental morphology–performance relationships in electrochemical energy storage, which in turn potentially enables the targeted development of electrodes with high volumetric efficiency.
Herein, we report a solvent mediated room temperature synthesis of a 3D MOF, [Zn2(bdc)2dabco]n (ZBD) (bdc= benzene-1,4-dicarboxylic acid and dabco = 1,4-diazabicyclo[2.2.2]octane) with controlled size, morphology and polymorphism by changing the kind and composition of the solvents used (Scheme 1). ZBD is a prototypical dabco MOF having unique guest-dependent framework flexibility and continuous 1D nanochannels,44,45 which are suitable for a broad range of applications such as gas storage, adsorption, separation and controlled polymerization.46–51 Dimethylformamide (DMF) and methanol (MeOH) induce the selective formation of a hexagonal rod MOF with Kagome nets (ZBDh) and a tetragonal plate MOF with square-grid nets (ZBDt), respectively. The size and morphology can be further controlled by using DMF and MeOH with different ratios as a cosolvent, which can tailor the polarity and solubility of the reactants and thus the nucleation and growth of MOFs. Conformal transformation of MOF precursors via carbonization at 900 °C leads to the preparation of hexagonal rod- and tetragonal plate-nanoporous carbon, which are subsequently employed as electrodes for EDLCs. Benefitting from a 2D morphology and a high packing density, tetragonal plate MDC delivers two times higher volumetric capacitance than hexagonal rod MDC, despite their similar gravimetric capacitances. The result highlights the importance of particle morphology for fabrication of high volumetric capacitance EDLC electrodes.
Cyclic voltammetry (CV) tests were performed at a cell voltage of 0–2.5 V and at scan rates of 10–500 mV s−1. The carbon integral volumetric capacitance, C (F cm−3), was calculated according to the following equation:
(1) |
Galvanostatic charge/discharge with potential limitation (GCPL) was applied at specific currents between 0.1 and 10 A g−1 in a voltage range from 0 to 2.5 V.
For long-term stability tests, the voltage of the cell was kept at 2.5 V for 100 h, and the specific capacity was measured every 10 h by galvanostatic cycling at 1 A g−1.
The morphologies of the as-synthesized ZBDh(t)-DxMy were investigated by electron microscopy (Fig. 1). ZBDh-D10 has an anisotropic hexagonal 1D rod-like morphology with hundreds of nanometers in length (Fig. 1a). As the volume ratio of MeOH increased from D10 to D9M1 to D7M3 to D6M4, the particle size increased from 0.58 to 1.2 to 13 to 75 μm (Fig. 1a–d), accompanied by morphology changes from nanorods to microplates. Hexagonal faces along the [001] direction were clearly observable in SEM. An intermediate mixture of both morphologies was observed in a narrow window of reaction conditions, e.g., D5.5M4.5 (Fig. S2†). Further increase in MeOH fraction ≥50vol% (i.e., 66 mol%) led to crystal phase transition from hexagonal to tetragonal framework topology. With an increase of MeOH ratio from D5M5 to D3M7 to D1M9 to M10, the particle size decreased from 3.4 to 0.72 to 0.65 to 0.43 μm (Fig. 1e–h), and the morphology changed from cubic microparticles to 2D-like cubic nanoplates. The changes in particle size and aspect ratio of ZBDh(t)-DxMy are summarized in Table S1 and Fig. S3.†
The selective formation of ZBDh and ZBDt was further supported by PXRD and N2 physisorption analysis (Fig. 2). PXRD patterns revealed that the synthesized ZBD can be assigned to one of the three classes, which correspond to hexagonal ZBDh-(D10, D9M1, D7M3, D6M4), an intermediate transition phase (D5.5M4.5), and tetragonal ZBDt (D5M5, D3M7, D1M9, M10) (Fig. 2a). The results were further confirmed by N2 physisorption at 77 K and pore size distribution calculated by a slit pore NLDFT equilibrium model. The two polymorphs ZBDh and ZBDt have distinct surface areas and pore sizes, which indicate phase pure ZBD synthesis. In general, ZBDh has a larger surface area and pore size than ZBDt, as calculated from single crystal structure analysis and simple geometric considerations.52 The representative N2 sorption isotherms for ZBDh-D10 and ZBDt-M10 are shown in Fig. 2b. N2 sorption isotherms show abrupt gas uptake and saturation at a very low relative pressure <0.05, which is typical type-I behavior of microporous materials (Fig. 2b and S4†). The BET surface areas of ZBDh and ZBDt were in the range of 1880–2160 m2 g−1 and 1480–1760 m2 g−1, respectively. Such values are comparable to or even larger than those of their bulk counterparts. The pore size distributions also show distinct micropores of ZBDh (1.3–1.8 nm) and ZBDt (1.0–1.2 nm) (Fig. 2c). Note that the estimated pore size from NLDFT is highly dependent on the equilibrium model used and is not as accurate as that from the single crystal analysis, but the characteristic micropores of ZBDh and ZBDt can be distinguished by using the same NLDFT model, which indicates the phase selective synthesis of ZBD. TGA data show that the synthesized ZBDh and ZBDt have thermal stabilities similar to those of the reference materials reported in the literature (Fig. S5†).44,52 They lost the guest solvents at temperatures below 200 °C, and started to decompose at 300 °C. The physicochemical properties of ZBD are summarized in Table S2.†
All results are in good agreement with the SEM observations and clearly support the successful synthesis of ZBD with controlled polymorphism, size and morphology by simply adjusting the kind and composition of the solvent used for MOF synthesis. It should be mentioned that the synthesis of nanoZBDt with controlled size is reported, to the best of our knowledge, for the first time in the present work.
The slight change in the size of the solvent template (guest molecule) can lead to the formation of ZBD with different framework topologies, i.e., the bulkier solvent molecule has bigger steric hindrance and may thus induce the formation of a MOF polymorph with larger pore size. As expected from the bulk density (ZBDh 0.730 vs. ZBDt 0.870 g cm−3) and pore size (ZBDh 1.5 vs. ZBDt 0.75 nm),44,52 bulkier DMF prefers ZBDh and MeOH prefers ZBDt. As confirmed in Fig. 2d and S3,† formation of ZBDh dominates over ZBDt formation in a broad range of MeOH mole fractions (0–0.6). ZBDt could be obtained when a MeOH mole fraction >0.66 was employed. Thus, it can be inferred that DMF has a stronger template effect than MeOH. DMF and MeOH work competitively rather than cooperatively. One overwhelms the other and preferably induces the formation of one particular crystal phase.
The competitive relationship between DMF and MeOH further enables the control over the size of the resulting ZBD crystals. For instance, in the case of ZBDh synthesis in a DMF/MeOH system, MeOH hinders the formation of ZBDh nuclei and decreases the rate of nucleation. Therefore fewer nuclei grow into larger crystals with increasing amount of MeOH (Fig. 2d). Likewise, the size of ZBDt increases with increasing amount of DMF. This phenomenon was further supported experimentally by the time dependent observation of crystal growth of ZBDh-D10 and ZBDh-D6M4. In a pure DMF system (ZBDh-D10) most of the ZBDh nuclei were rapidly generated and grew/saturated to a few hundred nanometer-sized crystals after reaction for 5 h, which suggests a fast and homogeneous nucleation at the early stage of the reaction (Fig. S6†). In contrast, in ZBDh-D6M4, the nucleation was largely retarded by MeOH, and therefore the secondary heterogeneous nucleation occurred continuously (Fig. S7†). The newly generated nuclei/small nanocrystals and large crystals existed together even after reaction for 32 h (Fig. S7a–f†). With increasing reaction time, the initially formed nuclei grew into large crystals at the expense of small nanocrystals via Ostwald ripening, leading to the formation of pure hexagonal microplates after reaction for >48 h (Fig. S7g and h†).
The particle morphology (e.g., aspect ratio) is also greatly affected by solvent composition (Table S1; Fig. S3†). As the amount of MeOH increased, the anisotropic hexagonal nanorods (ZBDh-D10) transformed into hexagonal microplates (ZBDh-D6M4) to tetragonal microcubes (ZBDt-D5M5) to 2D-like tetragonal nanoplates (ZBDt-M10) (Fig. 1 and 2d). This unique morphology evolution can be understood by the intrinsic crystal growth modes of ZBD and the polar protic nature of MeOH. ZBD has two growth modes (Zn-bdc and Zn-dabco) (Fig. S1†). Zn-bdc forms 2D hexagonal or tetragonal layers which are three dimensionally extended by dabco pillars. As a result, ZBDh consists of two hexagonal (001) faces terminated by Zn-dabco bonds and the other six faces terminated by Zn-bdc bonds (Scheme 2). The relatively high energy Zn-dabco (001) surfaces disappear through preferential anisotropic growth along the [001] direction in ZBDh-D10.20 However, such Zn-dabco growth can be largely suppressed by addition of polar protic MeOH, which can form hydrogen bonds with the nucleophilic dabco ligand, creating a shell of MeOH molecules around dabco (Scheme 2).55 The lone pair electrons of dabco possibly interact with the electron-poor hydrogen atoms of MeOH. As a result, polar protic MeOH can decrease the reactivity of dabco and hinder the growth of Zn-dabco (001) surfaces, whereas polar aprotic DMF cannot, thereby resulting in the formation of 2D-like nanoplates of ZBDt-M10.
The present solvent mediated approach enables the selective isolation of the pure crystalline ZBD phase with the desired structures by simply choosing the appropriate solvent composition. In addition, it is more energy efficient and environmentally friendly and requires less demanding synthesis equipment than conventional synthesis methods. In terms of framework stability and pore activation, the solvent template is highly preferred over the typical organic template (additive). The solvent generally has no direct interaction with the framework and is relatively easy to remove by volatile solvent exchange and thermal activation, thereby generating permanent porosity. In contrast, organic templates such as amines and polymers usually induce strong host–guest interactions and thus are often difficult to remove which might result in structural collapse.56–58
The morphologies of the parent ZBDs were well preserved after carbonization. Randomly yet homogeneously distributed mesopores were generated throughout the particles (Fig. 3a and b). ZBDh-D10-900 and ZBDt-M10-900 show type IV N2 sorption isotherms and an obvious hysteresis loop between P/P0 0.4 and 0.9, suggesting the presence of a substantial amount of mesopores (Fig. 3c). ZBDh-D10-900 has higher BET surface area and larger porosity than ZBDt-M10-900 (1490 vs. 1230 m2g−1), which is a similar trend to that in the parent ZBD precursors (Fig. 2d).
Fig. 3 Morphology conserved carbonization of ZBD. SEM and TEM images of (a) ZBDh-D10-900 and (b) ZBDt-M10-900, and their (c) N2 sorption isotherms and (d) pore size distributions. |
All Raman spectra have two broad peaks at 1342 and 1600 cm−1 that correspond to the D and G bands, respectively (Fig. S8†). The G-band arises from the stretching of any pair of sp2 sites whether in rings or chains. The D-band originates from the breathing mode of sp2 sites in rings of defects and disorders. Both carbons have an almost identical ID/IG intensity ratio near 0.97 and a broad G line width of 110 cm−1, suggesting the presence of graphitic cluster sizes in amorphous carbon smaller than 10 Å.59,60 The results reveal that both carbons have amorphous characteristics with a similar degree of graphitization and thus similar electrical conductivity. In addition, they have a comparable nitrogen content of 4.2 wt% which originated from dabco ligands, as determined by CHN elemental analysis (Table S3†). The N1S XPS spectra can be deconvoluted into pyridinic-N (398.2 eV), pyrrolic-N (399.5 eV), and graphitic-N (401.2 eV). The relative proportions of nitrogen species in both carbons are almost identical to each other (Fig. S9†). ICP-OES confirmed that most of the Zn (residual amount < 0.03 wt%) was removed during carbonization at 900 °C by carbothermal reduction and evaporation (Table S3†).
To demonstrate the morphology–volumetric capacitance relationships of ZBD derived carbons for EDLCs, ZBDh-D10-900 and ZBDt-M10-900 were fabricated into electrodes and characterized by cyclic voltammetry (CV) tests, in 1 M solution of tetraethylammonium tetrafluoroborate/acetonitrile (1 M TEABF4/AN), which is a common organic electrolyte.
Both ZBD-derived carbons present rectangle-like CV curves (0–2.5 V) without an obvious distortion even at a scan rate as high as 500 mV s−1, indicating a typical capacitive behavior with excellent rate capability (Fig. 4a and b). This is further underlined by the galvanostatic charging/discharging curve of ZBDt-M10-900 showing a symmetric triangular shape at 0.1 A g−1 and 10 A g−1 indicating the typical capacitor behavior (Fig. S10†). The small voltage drop even at 10 A g−1 reveals the low resistance of the EDLC. A voltage floating stability test performed for the EDLC cell with ZBDt-M10-900 as the electrode for 100 h at 2.5 V shows sufficient stability of the device (Fig. S11†). From a gravimetric perspective, due to the comparable surface area, conductivity and surface chemistry, the performances of the two carbons are rather similar to each other, exhibiting comparable specific capacitance and rate capability (around 80 F g−1 at 10 mV s−1, 50 F g−1 at 500 mV s−1) (Fig. 4c). However, the volumetric performance shows a significant difference. At the scan rate of 20 mV s−1 and 500 mV s−1, the differential current density of ZBDt-M10-900 is almost twice as high as that of ZBDh-D10-900 (Fig. 4a and b). This is further confirmed by the integral volumetric capacitance (Fig. 4d), where the volumetric capacitance of ZBDt-M10-900 is pronouncedly higher than that of ZBDh-D10-900 at all scan rates. This is attributed to the different pore structures and particle morphology. The meso- and macro-porosities of ZBDt-M10-900 are much lower compared with those of ZBDh-D10-900, leading to a higher density of the carbon materials in the electrodes. In addition, the 2D plate morphology of ZBDt-M10-900 can facilitate preferential and uniform alignment along the basal plane with relatively high packing density, as demonstrated by the high volumetric capacitances of various 2D materials.37,61,62 Simulation studies also show that particle elongation can lead to a decrease in packing density because of the increase in the orientationally excluded volume.63 Therefore, despite a similar gravimetric specific capacitance, ZBDt-M10-900 can provide higher volumetric capacitance. On the other hand, the lower meso–macroporosity of ZBDt-M10-900 results in the faster decay of capacitance at higher rates (Fig. 4c), as the meso/macropores tend to facilitate ion transport inside carbon electrodes. Note that ZBDt-M10-900 still provides higher volumetric capacitance over the entire range of scan rates investigated (Fig. 4d).
The general importance of the particle morphology and porosity on the volumetric efficiency of supercapacitor electrode materials is further shown by using non-commercial ordered mesoporous carbon CMK-3 as an additional reference material.64 CMK-3 typically shows a hexagonal rod-like shape and particles with sizes in the range of 1–3 μm.30 Its typical mesoporous structure is shown by the nitrogen physisorption isotherm (Fig. S12a†) with a pore volume of 1.23 cm g−1. The gravimetric capacitance is ∼80 F g−1, thus comparable to that of the two ZBD-derived carbons (Fig. S12c†). However, due to the mesoporous feature and different particle textures leading to moderate packing density, its volumetric capacitance is lower than that of ZBDt-M10-900 with the highest packing density, and higher than that of ZBDh-D10-900 with the lowest packing density (Fig. S12b†). Hence, the packing density generally plays a vital role in the volumetric capacitance. It has to be pointed out that in comparison to other recently reported supercapacitor electrode materials, ZBDt-M10-900 does not show a particularly impressive volumetric or gravimetric capacitance but the MOF-derived carbons are a suitable model system to evaluate the importance of the morphology of the electrode material particles for the volumetric capacitance of supercapacitors (Table S4†).
Footnote |
† Electronic supplementary information (ESI) available: Supplementary schematic representation, particle size distribution, aspect ratio, SEM images, N2 isotherms, TGA, and Raman spectra are included. See DOI: 10.1039/c8ta07700b |
This journal is © The Royal Society of Chemistry 2018 |