Download - On/off switchable electronic conduction in intercalated metal … · be used to induce the Li-intercalated state in 2,6-Naph(COOLi) 2, which is black in this state. Although exfoliation

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Toyota Central R&D Laboratories Inc., Nagakute, Aichi 480-1192, Japan.*Corresponding author. Email: [email protected]

Ogihara, Ohba, Kishida, Sci. Adv. 2017;3 : e1603103 25 August 2017

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On/off switchable electronic conduction in intercalatedmetal-organic frameworksNobuhiro Ogihara,* Nobuko Ohba, Yoshihiro Kishida

The electrical properties of metal-organic frameworks (MOF) have attracted attention for MOF as electronicmaterials. We report on/off switchable electronic conduction behavior with thermal responsiveness in interca-lated MOF (iMOF) with layered structure, 2,6-naphthalene dicarboxylate dilithium, which was previously reported as areversible Li-intercalation electrode material. The I-V response of the intercalated sample, which was prepared usinga chemically reductive lithiation agent, exhibits current flow with sufficiently high electronic conductivity, eventhough it displays insulating characteristics in the pristine state. Calculations of band structure and electron hoppingconduction indicate that electronic conduction occurs in the two-dimensional p-stacking naphthalene layers whenthe band gap is decreased to 0.99 eV and because of the formation of an anisotropic electron hopping conductionpathway by Li intercalation. The structure exhibiting electronic conductivity remains stable up to 200°C and revertsto an insulating structure, without collapsing, at 400°C, offering the potential for a shutdown switch for batterysafety during thermal runaway or for heat-responsive on/off switching electronic devices.

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INTRODUCTIONThe physical properties of several types of metal-organic frameworks(MOFs), which comprise an organic linker and metal ions/clusters,have been studied using tunable and well-defined porous specimensto explore their applicability in gas adsorption and separation (1).Currently, their electrical properties (2–4), such as their conductivity(5–10), magnetism (11, 12), photonic characteristics (13, 14), andenergy storage properties (15–18), are also being investigated. Untilrecently, various kinds of electrically conductive MOFs have beenreported: (i) tetrathiafulvalene-based MOFs with one-dimensional(1D) p-stacking interaction channels (5, 6), (ii) Ni-dithiolene–basedMOFs using p-type semiconductors via a redox mechanism (13),(iii) hexagonally packed dihydroxyterephthalate-based MOFs with1D charge mobility using metal-oxygen (or metal-sulfur) chains(7–9), (iv) benzentricarboxylate-based MOFs infiltrated withtetracyanoquinodimethane exhibiting tunable electrical conductivity(10), and (v) highly conducting 2D MOFs by stacking sheets withexpanded p conjugation (19–21). These electrically conductingMOFsare potential candidates for future models of reconfigurable electronicsand sensors. These also could be useful as an alternative approach inorganic electronic devices to meet the safety requirements of thermallyresponsive switching performance at high temperature.

We recently reported the synthesis of intercalated MOF (iMOF)electrode materials through an electrochemically reversible chargetransfer Li-intercalation reaction (15–17). As shown in Fig. 1, 2,6-naphthalene dicarboxylate dilithium [2,6-Naph(COOLi)2], a repre-sentative iMOF, is composed of aromatic dicarboxylate lithium andforms an organic-inorganic layered structure based on p-stacked naph-thalene packing layers and tetrahedral LiO4 units, respectively. It alsoexhibits two–electron transfer–related reversible intercalation of Li ions,corresponding to a reversible specific capacity of ca. 220 mA·hour g−1,at a constant potential of 0.8 V (versus Li/Li+), without dissolving in theorganic electrolyte. This operating potential enables both maintenanceof high voltage and suppression of lithium dendrite deposition, generatingan internal short circuit for safety issues at the same time (15). Further-more, 2,6-Naph(COOLi)2 also exhibits another interesting property in

that it undergoes a very small change in its unit cell volume (0.33%)while maintaining its framework during Li intercalation. Moreover,the powder of this material, which consists of micrometer-sized whiteparticles, exhibits insulating characteristics. However, the material iselectrochemically fully reversible when used to prepare electrodes; thisis true even at low carbon content (17). Although there are a numberof reports on aromatic dicarboxylate salts that are similar to electrodematerials (22–24), these materials have been considered as havingpoor electronic conductivity, and the essential mechanisms ofelectronic conduction responsible for the electrochemical reversibilityremain to be elucidated. To clarify the reason for the electrochemicalreversibility of iMOF electrodes, we performed experimental andtheoretical studies of the electrical properties of these electrodes. Ourresults provide direct evidence for the switching electronic conductivityof 2,6-Naph(COOLi)2while retaining its ionic conductivity and thermalstability, and suggest that these materials would be highly suited forproduction of new class of heat-responsive organic switching devices.

RESULTSWe experimentally characterized the electrochemical behavior of 2,6-Naph(COOLi)2 electrodes (fig. S1A). From the discharge curvescorresponding to the Li-intercalation reaction, the first dischargepotential of these electrodes was found to be ca. 100 mV lower than thatduring the second and subsequent cycles. As shown in fig. S1B, thedifferential capacity (dQ/dV) curves exhibited a decrease in polarizationwith repeated cycling, with the potential change between the first andsecond discharge cycles being the largest. The current-voltage (I-V)resistance after the second cycle was half the value of the first one(fig. S1B, inset). This suggested that the decrease in the internal resistanceis significantly related to Li intercalation, especially during the first cycle,and that the extent of Li intercalation is high enough to change thematerial from an insulating one to a conducting one.

The detailed Li-intercalation/deintercalation processes of the2,6-Naph(COOLi)2 electrodes were analyzed using the galvanostaticintermittent titration technique (GITT) to investigate the cause of theinternal resistance (fig. S2A). The open circuit potential (OCP) for theinitial cycle and that after 10 cycles were almost identical, as shown infig. S3. What was remarkable was the difference in the polarization

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behavior of DEi when a current was applied at the OCP (for a detailedexplanation, see fig. S2B). The polarization in the initial state (fig. S3A)was twice as large as that after the 10th cycle (fig. S3B). Similarly, the IRdrop resistance in the initial state (200 to 300 ohms) (fig. S3A, inset),whose primary component was the electronic resistance from theviewpoint of the time constant, was two to three times larger than thatafter the 10th cycle (ca. 100 ohms) (fig. S3B, inset). Therefore, it canbe concluded that the change in internal resistance contributed tothe electronic conduction, which was significantly related to the Li-intercalation state.

Next, we performed simulations using first-principles calculationsbased on the density functional theory to understand the mechanismof electronic band conduction before and after Li intercalation. Wehad previously determined two models for the positions of the inter-calated Li by x-ray diffraction (XRD) measurement and Rietveld re-finement using first-principles calculation (15): In model 1, theintercalated Li ions are located in the vicinity of the tetrahedral LiO4

network layers, whereas inmodel 2, themain structure is located in thep-stacked naphthalene layers (fig. S4). In the pristine state (Fig. 2A),the band gap is relatively wide and has an energy (Eg) of 3.27 eV. Asper models 1 and 2 corresponding to Li intercalation (Fig. 2, B and C,respectively), the Fermi level moves up to the conduction band afterLi intercalation, owing to the increase in the electron density becauseof the organic unit (fig. S5). Therefore, the calculated Eg values formodel 1 (Eg = 1.51 eV) and model 2 (Eg = 0.99 eV) are lower, indicat-ing that Li intercalation transforms 2,6-Naph(COOLi)2 from aninsulating material into a semiconducting one; this is especially thecase for model 2. Although a 2D channel-type MOF with high con-ductivity exhibits a low band gap (Eg = 0.13 to 0.23 eV) (21), thecalculated Eg value for model 2 is relatively lower than the Eg valuesof previously reported 1D channel-typeMOFs, such as cubic arrays ofZn4O(CO2)6-type frameworks (Eg = 1.3 to 6.8 eV) (25, 26) or hexag-onally packed cylindrical channels of M3O3(CO2)3-type frameworks(Eg = 1.47 to 2.60 eV) (9). The CBMof pristine and that ofmodel 1 arelinear (red lines in Fig. 2, A and B). On the other hand, the CBM inmodel 2 is curved and has dispersion (Fig. 2C). This suggests that elec-trons in conduction band in model 2 are delocalized and thatelectronic mobility increases. The calculated 2D and 3D electron den-sity distributions of the CBM indicated the following: The pristinestructure is localized intramolecularly along the cross-sectional direc-


tion of the p-stacked naphthalene packing layers (Fig. 2, D and G). Incontrast, in the case of model 1, the Li-intercalated structure is de-localized intramolecularly between C1 and C1′ (Fig. 2, E and H). Asper model 2, it is delocalized intermolecularly between C1 and C2 (orC1′ and C2′), in addition to being intramolecularly delocalized (Fig. 2,F and I). This reveals that, in model 2, Li intercalation provides anelectron hopping conduction path via the 2D p-stacked naphthalenepacking organic layers.

To further investigate electron conduction between adjacent mol-ecules, the hopping conduction parameters, including the transfer in-tegral (V), diffusion coefficient (D), and drift mobility of hopping (mh),for the pristine structure and the Li-intercalated one correspondingto model 2 were determined on the basis of the Marcus theory (for adetailed explanation, see Materials and Methods) (27). To be able todescribe hopping transport before and after the Li intercalation ofthe 2,6-Naph(COOLi)2 crystal, we considered the following threehopping paths in which charge could only be transferred betweenadjacent molecules, as shown in Fig. 3 (A and B): (i) p-conjugatedmolecules arranged in a layered basal herringbone packing structure(P), (ii) basal-edge packing (T), and (iii) edge-edge packing acrossthe tetrahedral LiO4 layer (L). The transfer integrals of all threehopping pathways corresponding to the Li-intercalation state werelarger than those before Li intercalation, as shown in Fig. 3C. In con-trast, the diffusion coefficient (Fig. 3D) and drift mobility of hopping(Fig. 3E) for path T increased markedly, by three orders of magni-tude, after Li intercalation. Furthermore, those for path P exhibitedhigh values, whereas those for path L decreased in value. These resultssuggest that anisotropic electron hopping conduction occurs in thep-stacked naphthalene packing layer along the b-c planar directionbut not in the organic-inorganic stacking layer along the b-c verticaldirection; this gives rise to anisotropic 2D electronic transport withinthe organic layer of the 2,6-Naph(COOLi)2 crystal. Moreover, theseresults agree with the calculated electron density distribution of theband structure, as shown in Fig. 2. The hopping conduction calcula-tions performed in this work can be used as a molecular design tool tofind high–electron conduction iMOF materials.

To gather direct evidence for electronic conduction in 2,6-Naph(COOLi)2, a Li-intercalated sample was prepared using lithiumnaphthalenide, which is a known reductive lithiation agent (28, 29),and electronic conductivity measurements were performed on the

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Fig. 1. Schematic illustration of redox mechanism for crystalline 2,6-Naph(COOLi)2. Orange, intercalated Li; green, pristine Li; red, O; blue, C; white, H.

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synthesized sample. With the reductive lithiation agent and compos-ite electrode sample of conducting carbon free (see Fig. 4A), a Li-intercalated sample could be prepared without having to use anyelectrochemical techniques. The sample changed from white in thepristine state to black in the Li-intercalated state (Fig. 4B). Furthermore,the XRD patterns corresponding to the pristine and Li-intercalationstates were different and confirmed that the respective crystal structuresof the samples were similar to those reported previously (fig. S6) (15).This confirmed that naphthalene-catalyzed reductive lithiation canbe used to induce the Li-intercalated state in 2,6-Naph(COOLi)2,


which is black in this state. Although exfoliation occurs when naph-thalene is incorporated in the interlayer when other layered materials,such as graphite (29) and chalcogenides (30), are made to react withlithium naphthalenide, the reaction of 2,6-Naph(COOLi)2 and lithiumnaphthalenide did not cause the exfoliation of the crystal structure,which resulted in amorphous naphthalene.

The I-V curves of the electrodes of the Li-intercalated sample usingclosed parallel-electrode cell (fig. S7) were linear, as shown in Fig. 4C,indicating that the Li-intercalated sample exhibited a favorable currentresponse at an applied bias voltage of 0.5 V; in contrast, in the pristine


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Fig. 2. Predictive calculation of electronic band conduction. (A to C) Calculated electronic band structures obtained from the HSE06 functional for the pristine sample(A) and the Li-intercalated sample corresponding to model 1 (B) and model 2 (C). The valance band maximum and conduction band minimum (CBM) are represented inthe band structures as blue and red lines, respectively. (D to F) 3D electron density distribution (obtained with the PBEsol functional) of the CBM of the 2,6-Naph(COOLi)2crystal structure for the pristine sample (D) and the Li-intercalated sample corresponding to model 1 (E) and model 2 (F). Orange, intercalated Li; green, pristine Li; red, O;blue, C; white, H. (G to I) 2D electron density distribution of the CBM across the p-stacking packing naphthalene layer via the C1, C1′, C2, and C2′ carbons corresponding tothe figure above for the pristine sample (G) and the Li-intercalated sample obtained using model 1 (H) and model 2 (I).

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state, the sample exhibited a negligible current response. To investigatethe detailed electrical characteristics, we performed electrochemicalimpedance spectroscopy (EIS) using the same cell. In the Bode plotsshown in Fig. 4D, the Li-intercalated sample showed ca. 105 ohms.This value is four orders of magnitude lower than that of the pristinesample (ca. 109 ohms), as the frequency becomes lower. In theNyquistplots shown in Fig. 4E, blocking behavior corresponding only to ionicconductionwithout electronic conductionwas observed in the pristinesample, and a depressed semicircle was observed in the Li-intercalatedsample, which denotes amixed electronic and ionic conduction (fig. S8)(31–33). Because the low-frequency resistance component of the Li-intercalated sample corresponds to an electronic resistance (Re), the ob-served current response of the Li-intercalated sample shown in Fig. 4Cis due to a decrease in the Re, suggesting that a switchable electronic con-duction path is formed by the Li intercalation of 2,6-Naph(COOLi)2.In the states with white and black colors (Fig. 4B), 2,6-Naph(COOLi)2


exhibited insulating and conducting characteristics, respectively. Theresults of inductively coupled plasma optical emission spectroscopy(ICP-OES) showed that the quantity of the intercalated Li can becontrolled by the treatment time with lithium naphthalenide; in addi-tion, electronic conductivity was improved depending on the quantityof the intercalated Li (fig. S9).


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Fig. 3. Predictive calculation of electron hopping conduction. (A and B)Schemes of electron hopping pathways between adjacent molecules for the pristinesample (A) and the Li-intercalated samples corresponding to model 2 (B). Orange,intercalated Li; green, pristine Li; red, O; blue, C; white, H. (C to E) Comparison ofthe predicted electronic transfer integral (C), diffusion coefficient (D), and driftmobility of hopping (E) before and after Li intercalation.

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Fig. 4. Experimental confirmation of switchable electronic conduction. (A) Schematicof chemical reductive lithiation of 2,6-Naph(COOLi)2 using lithium naphthalenide.(B) Photographs of 2,6-Naph(COOLi)2 electrode without conductive carbon before(left) and after (right) being immersed in a lithium naphthalenide/tetrahydrofuran(THF) solution. (C to E) Comparison of I-V curves (C), Bode plots (D), and Nyquistplots (E) for 2,6-Naph(COOLi)2 samples before (blue) and after (red) beingimmersed in the same solution. In (E), the inset is the equivalent circuit for amixed electronic and ionic conductive material. The intersections at high (100 kHz)and low (1 Hz) frequencies with the real axis are the mixed electronic and ionicresistances in parallel (ReRi/Re + Ri) and the only electronic resistance (Re), respectively.(F) Arrhenius plots of each conductivity obtained from I-V curves and EIS at varioustemperatures (fig. S10), se (red) and si (green), which are calculated using theequivalent circuit model shown in the inset of (E) for the mixed electronic and ionicconductive material, and polarization measurement, sI-V (blue). All conductivityvalues were calculated using percolation theory (see Materials and Methods).

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Both the I-V curves and Nyquist plots of the Li-intercalated sampleshowed lower resistance behavior as the temperature increased (fig. S10).The Nyquist plots of the Li-intercalated sample showed depressedsemicircle at all temperatures. The bulk electronic conductivity (se)and ionic conductivity (si) were calculated separately from Nyquistplots of highest and lowest values using the equivalent circuit modelfor mixed electronic and ionic conductors (fig. S8B). The obtained seand si were 10

−6 to 10−5 S cm−1 and 10−5 to 10−4 S cm−1 within atemperature range of 25° to 60°C, respectively. The se agrees with thatcalculated from the I-V curves (sI-V). The obtained se is comparableto the conducting MOFs, such as hexagonally packed cylindricalchannels of M3O3(CO2)3-type frameworks (10−16 to 10−6 S cm−1)(9), tetrathiafulvalene-based frameworks (10−6 to 10−4 S cm−1) (6),tetracyanoquinodimethane-based frameworks (10−5 to 10−1 S cm−1)(34), and dithiolene-based frameworks (10−11 to 10−5 S cm−1) (13).In addition, compared with se (10

−8 to 10−4 S cm−1) and si (10−5

to 10−3 S cm−1) of pristine oxide-based electrode materials for batteries,which are known as mixed electronic and ionic conductors (33, 35, 36),the proposed material has comparable electronic and ionic conduc-tions. Both se and si display Arrhenius behavior (Fig. 4F). The acti-vation energies (Ea) for se and si calculated from Arrhenius plots areca. 18.3 kJ mol−1 (or ca. 190 meV) and ca. 3.5 kJ mol−1 (or ca. 36 meV),respectively. Although the Ea for se is slightly larger than that in theprevious reports of high-mobility conducting MOF materials (10, 20),this value is one order of magnitude smaller than that of the pristine


electrode materials (36, 37). The Ea for si was lower than that of theionic conductors of well-known solid electrolytes (38–40). The low Eaensures the least fluctuation of ionic conductivity and therefore favorsthe practical applications in a broad operating temperature range.

When the sample was used as an electrode material, the disappearanceof its electronic conductivity during elevated temperature process atapproximately 200°C was expected to act as a shutdown switch forbattery safety, because its insulating nature would increase safety duringthermal runaway. To verify this phenomenon, we measured the XRDpatterns of the Li-intercalated sample obtained during heating (Fig. 5A).The results showed that the crystal structure exhibiting electronicconductivity remains intact until approximately 180°C. Thereafter,the crystal structure changed to the initial one (blue pattern in fig. S6)exhibiting insulating characteristics via a two-phase coexistence regionin a temperature range of 180° to 250°C, where the initial structure withinsulating property was maintained until a temperature of 400°C. Inaddition, direct evidence of disappearance of electronic conductivityabove 200°C was given by EIS and I-V curve of Li-intercalated sampleafter heating (Fig. 5, B to D). In the heating sample at 100°C, whichkeeps the Li-intercalated crystal state, although slight resistance in-crease was confirmed at a low frequency from the Bode plots (Fig. 5B),the current response was observed from the I-V curve (Fig. 5C). In theNyquist plot, a new semicircle appeared in the low-frequency range(Fig. 5D). This indicates the presence of a grain-boundary resistanceby growth of the Li-containing film, which was formed during the

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Fig. 5. Heat-responsive on/off switching of electronic transport characteristics. (A) Temperature dependence of in situ powder XRD patterns of Li-intercalated 2,6-Naph(COOLi)2for temperatures of up to 400°C. (B to D) Change in Bode plots (B), I-V curves (C), and Nyquist plots (D) for pristine, Li-intercalated samples and Li-intercalated samples preparedby heat treatment at temperatures of 100° and 200°C.

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chemically reductive lithiation at the grain boundary (fig. S9). Theequivalent circuit of this Nyquist plot can be explained in fig. S8C(32). There is almost no change in bulk electronic resistance whilegenerating grain-boundary resistance in the sample heated at 100°C.In contrast, in the heating sample at 200°C, which returns to the initialcrystal state, from the impedance behavior, the electronic resistanceincrease was confirmed and the current response of the I-V curvedrastically disappeared.

To investigate the Li that disappeared from the main structure duringheat treatment, we compared the XRD patterns of pristine nondopedsample and Li-doped sample after heating, as shown in fig. S11. In theLi-intercalated sample after heating, several new peaks correspondingto LiO, Li2O, Li2O H2O, and Li2CO3 appeared in the high-angleregion. These peaks appeared from temperatures higher than 200°Ccorresponding to the start of the Li release (fig. S12). These areLi-containing components that become resistant phase. Therefore,the XRD results imply that the Li released from the Li-doped sampleduring heating is incorporated into the Li-containing components andbecomes a grain-boundary resistance component. In addition, in theLi-intercalated sample after heating, although the (011) reflectioncorresponding to the inorganic tetrahedral LiO4 network (16) didnot change, the (102) reflection corresponding to the p-stacked naph-thalene packing layer shifted to the lower angle. This result means theexpansion of the p-stack distance of naphthalene and thus the inhibi-tion of the hopping conduction of the main structure. Therefore, theheat treatment of the Li-doped sample includes the following three-electron conduction blocking mechanisms: (i) the disappearance ofthe main crystal structure responsible for hopping conduction accom-panying Li release, (ii) the inhibition of hopping conduction accom-panying the spreading of stacked naphthalene distance in the maincrystal structure, and (iii) the formation of Li-containing insulatinglayer with the released Li and heating at the grain boundary. Thiselectronic conduction blocking mechanism will be a useful shutdownfunction during thermal runaway. Although there are many indirectapproaches to battery safety using electrolyte (41, 42), separators(43), and peripheral materials (44, 45), the proposed material is re-garded as a direct approach to the safety of the electrode that storeselectrical energy.

, 2020

DISCUSSIONThis study is the first to provide theoretical and experimental evidencefor mixed electronic and ionic conductivity induced in electrodematerials based on aromatic dicarboxylate salts such that the materialsexhibit a reversible redox response and should facilitate the devel-opment of MOFs that exhibit switchable anisotropic electronic con-duction. Although the electronic conduction observed in previouslyreported MOF materials has been attributed to metal-metal or metal-ligand interactions (2–4, 7–10, 12, 18), our results demonstrate hoppingconduction in an iMOFmaterial based on ligand-ligand interactions be-tween the p-stacking aromatic units owing to reductive lithium doping;the organic framework design will be the key to improving theelectronic conductivity.

The mechanism of electronic conduction triggered by Li intercalationwhile keeping ionic conductivity is such that it allows for electrochemicalreversibility and can be exploited to fabricate low-resistance electrodesfor batteries and supercapacitors that exhibit high-power perform-ances. We also envisage that the switchable electronic conduction/insulation properties of this Li-doped iMOF at approximately 200°C


will be exploited to fabricate heat-responsive on/off switching electronicdevices, such as battery electrode with shutdown function or organictransistors for producing current interrupters that improve safety.To achieve organic transistors, technology of pellets and thin-filmfabrication with excellent responsiveness using the proposed materialwould be necessary.

MATERIALS AND METHODSSample synthesisThe synthesis of 2,6-Naph(COOLi)2 was reported previously (15).2,6-Naph(COOLi)2 was synthesized from lithium hydroxide mono-hydrate (4.85 g, 115.6 mmol) and 2,6-naphthalenedicarboxylic acid(10.0 g, 46.3 mmol) by heating using reflux in 300 ml of methanol.The solution was initially clear; however, a white precipitate formedover the course of 30 min. The suspension was stirred under thereflux conditions for 12 hours. The liquid in the suspension wasevaporated, and the resulting solid was washed with methanoland dried under ambient conditions. The product obtained wasin the form of needle-shaped white crystals [96.8% yield based on2,6-Naph(COOLi)2].

Chemical Li intercalationA lithium naphthalenide solution (0.1 mol dm−3) was prepared usinga previously described procedure (28, 29). Naphthalene was dissolvedin anhydrous THF to prepare a 0.1 M solution, and an equimolaramount of lithium was added to the 0.1 M naphthalene solution understirring at room temperature in an Ar-filled glove box. The solutionturned dark green after the addition of lithium. The quantity of Lielement in the pristine and reduced samples was confirmed by ICP-OES(CIROS 120EOP, Rigaku and SPECTRO).

In situ powder XRDThe XRD pattern of the Li-intercalated 2,6-Naph(COOLi)2 samplefree of conductive carbon was obtained using a Debye-Scherrercamera at the BL5S2 beamline, Aichi Synchrotron Radiation Center,Aichi prefecture, Japan (proposal no. 201601031). The data wererecorded with a PILATUS 100K detector (DECTRIS Ltd.) with aresolution of 1.00011 Å (calibrated using standard CeO2 powder)for 2q values of 10° to 24°; the step size was 0.01°. The sample washeated with a high-temperature N2 gas blower at a rate of 5°C min−1

to temperatures of up to 400°C, and the XRD pattern was recordedevery 10°C after a wait time of 2 min for each temperature. Before thein situ XRD measurements, the Li-intercalated 2,6-Naph(COOLi)2sample was prepared by soaking the 2,6-Naph(COOLi)2 electrode freeof conductive carbon in the lithium naphthalenide/THF solution(0.1 mol dm−3) in an Ar-filled glove box. The 2,6-Naph(COOLi)2electrodes free of carbon as a conductive agent were also prepared bydispersing 98 weight % (wt %) 2,6-Naph(COOLi)2 and 2 wt % car-boxymethylcellulose (CMC) in water and then by coating the slurryon a Cu foil; the electrode loading rate was ca. 3 mg cm−2. After theelectrode had been prepared, the sample was removed from theelectrode, placed in a capillary tube, and sealed to prevent oxidationby moist air during the XRD measurements, which were performedin an Ar-filled glove box.

Electrochemical measurementsThe galvanostatic charge/discharge and GITT measurements wereperformed using a two-electrode system of Li/2,6-Naph(COOLi)2 cells

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assembled in an Ar-filled glove box. The 2,6-Naph(COOLi)2 electrodeswere prepared by coating a dispersion of 81 wt % 2,6-Naph(COOLi)2,14 wt % carbon black as a conductive agent, 2 wt % CMC as thebinder, and 3 wt % styrene butadiene rubber (also used as a binder)in water on a Cu foil. The electrode loading rate was ca. 3 mg cm−2.The electrolyte used was 1 M LiPF6 dissolved in a solution of ethylenecarbonate, dimethyl carbonate, and ethyl methyl carbonate (volumeratio of 30:40:30). A microporous polypropylene film was used as theseparator. During the galvanostatic charge/discharge measurementsto determine the specific capacity and electrochemical reversibilityof the 2,6-Naph(COOLi)2 electrodes at 25°C, coin-type cells werecycled between 0.5 and 2.0 V (versus Li/Li+) at a rate correspondingto the full charging to the theoretical capacity of 2,6-Naph(COOLi)2per 10 hours.

Electrical measurementsAll electronic transport characteristics were performed using closedparallel-electrode cell with a four-electrode system (12962A, Solartron).Cell resistance without sample is smaller by seven to eight orders ofmagnitude when compared to cell with sample and can therefore beignored. The I-V characteristics and EIS were measured at 25°C with aSolartronMultiStat system (1470E, Solartron) using 2,6-Naph(COOLi)2electrodes free of conductive carbon; the measurements were madeboth before and after the electrodes were immersed in the lithiumnaphthalenide/THF solution. After the removal of the immersedsolution, the electrode was washed with THF. Before measurement,the electrodes were dried. The closed parallel-pellet cells were thenprepared with electrodes in a dry state (fig. S7). The electrical connec-tionwas formedusing a piece ofCu foil. The I-V curves weremeasuredat a scan rate of 100 mV s−1 within the polarization range of −0.5 to0.5V using a two-electrode system assembled in anAr-filled glove box.EIS measurements of the same cell were performed at OCP. The fre-quency was varied from 100 kHz to 1 Hz with a perturbation ampli-tude of 0.5 V (peak to peak). The measurements were performedbetween−30° and 60°C. Before themeasurement of the heat treatmentsamples, the Li-intercalated sample was annealed at temperatures of100° and 200°C for 5 min under Ar atmosphere.

Percolation theory in electronic propertiesThe effective electronic conductivity (seff) of conductor-insulatorcomposites is defined as follows

seff ¼ 1Rx

LS

ð1Þ

where Rx, L, and S are resistance, composite electrode thickness, andelectrode geometric area, respectively. The actual conductivity can beestimated from percolation theory (46–48). According to the percolationtheory, the effective electronic conductivity (seff) of conductor-insulatorcomposites can be expressed as follows

seff ¼ s0ðp� pcÞn ð2Þ

where s0 is the actual conductivity, p is the weight fraction of theconductive material, pc is the critical weight fraction at percolationthreshold, and n is the critical exponent that depends on the


topological dimensionality of the percolating system. The p valueof this composite electrode in this study is calculated to be ca.0.466. The pc value of the conductor-insulator composites is re-ported as 0.32 (49, 50). In the case of random void model, in whichinsulating penetrable spheres are randomly immersed in a conduct-ing continuum, the exponent of n is found to be 2.38 (49).

Kinetic interpretation of individual conductivities ofLi-intercalated 2,6-Naph(COOLi)2 composite electrodesTo determine the kinetic parameters of the composite electrodes, theactivation energies of each conductivity were evaluated using theArrhenius equation

s ¼ A exp

�� EaRT

�ð3Þ

where A, Ea, R, and T are the frequency factor, activation energy,gas constant, and absolute temperature, respectively.

Calculation of electronic band structure and densityof statesFirst-principles calculations for structural optimization and analysis ofthe electronic properties were performed with the Vienna ab initiosimulation package (51) using the projector-augmented wave method.The cutoff points for the kinetic energy and k-point spacing were setto 500 eV and approximately 0.1 Å−1, respectively. During the optimizationof the crystal structures, Gaussian smearing with a width of 0.2 eV wasadopted for the integral over the Brillouin zone. For the exchange-correlation energy within the density functional theory, the PBEsol func-tional (52), which revises the Perdew-Burke-Ernzerhof–generalizedgradient approximation functional to improve the equilibrium prop-erties of solids and surfaces, was applied. For the electronic structurecalculations, the tetrahedron method with Blöchl corrections and theHSE06 functional (53, 54) were used for the integral over the Brillouinzone and the exchange-correlation energy, respectively.

Calculation of electron hopping conduction usingMarcus theoryThe Marcus theory was widely used to calculate the hopping mobilityin organic semiconductors (27, 55, 56). The electron hopping ratebetween neighboring molecules (W) can be determined on the basisof the Marcus theory as shown

W ¼ V2

ħp

lkBT

� �1=2

exp

�� l4kBT

�ð4Þ

where V is the transfer integral between the initial and final states (inelectron volts), l is the reorganization energy (which is defined asthe energy change associated with the geometry relaxation duringcharge transfer), ħ is Plank’s constant, and kB is the Boltzmann con-stant. The diffusion coefficient (D) is calculated from the hoppingrates (cm2 s−1) as

D ¼ 12n∑ir2i WiPi ð5Þ

where n is the dimensionality, i represents a specific hopping pathway,ri is the hopping distance between two molecules, and Pi is the

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hopping mobility, which is calculated by

Pi ¼ Wi

∑iWi

ð6Þ

The drift mobility of hopping (mh) can then be evaluated from theEinstein relation (cm2 V−1 s−1)

mh ¼e

kBTD ð7Þ

where e is the electronic charge.The transfer integral (V) was calculated using the site-energy

correction method (57) and the DiPro approach (58). The reorganizationenergy (l) was obtained from the sum of the following two energydifferences (56)

l ¼ l1 þ l2 ð8Þ

l1 ¼ E*eq0 � E*

eq* ð9Þ

l2 ¼ E0eq* � E0

eq0 ð10Þ

whereE*eq0 ðE*

eq*Þ is the energy of a charged molecule in the equilibriumneutral (charged) geometry andE0

eq0ðE0eq*Þ is that of a neutral molecule

in the equilibrium neutral (charged) geometry. The calculations forthe transfer integral and reorganization energy were performed usingthe Gaussian 09 package (59), with the B3LYP functional and the6-311++G* basis set for C12H6O4Li2(Li2) molecules extracted from theexperimental crystal structure.

on June 29, 2020org/

SUPPLEMENTARY MATERIALSSupplementary material for this article is available at http://advances.sciencemag.org/cgi/content/full/3/8/e1603103/DC1fig. S1. Charge/discharge curves and differential capacity (dQ/dV) curves.fig. S2. The profile of GITT against time.fig. S3. Electrochemical behavior before and after Li intercalation.fig. S4. Schematic illustrations of the crystalline structures corresponding to the pristine andLi-intercalated states.fig. S5. DOS and PDOS profiles.fig. S6. Ex situ powder XRD patterns without conductive carbon before (left) and after (right)being immersed in a lithium naphthalenide/THF solution.fig. S7. Schematic illustration of closed parallel-pellet cell used for electrical measurements.fig. S8. Interpretation of electron and ion transport mechanism of the proposed material.fig. S9. Dependence of electrical properties on the treatment time of the chemical Liintercalation.fig. S10. Electrical properties of the Li-intercalated 2,6-Naph(COOLi)2 samples at varioustemperatures.fig. S11. Powder XRD patterns of pristine nondoped sample (black line) and Li-doped sampleafter heating (red line).fig. S12. Temperature dependence of in situ powder XRD patterns of Li-intercalated2,6-Naph(COOLi)2 for temperatures of up to 400°C in the high-angle region.

REFERENCES AND NOTES1. H. Furukawa, K. E. Cordova, M. O’Keeffe, O. M. Yaghi, The chemistry and applications of

metal-organic frameworks. Science 341, 1230444 (2013).2. G. Givaja, P. Amo-Ochoa, C. J. Gomez-Garcia, F. Zamora, Electrical conductive

coordination polymers. Chem. Soc. Rev. 41, 115–147 (2012).


3. V. Stavila, A. A. Talin, M. D. Allendorf, MOF-based electronic and opto-electronic devices.Chem. Soc. Rev. 43, 5994–6010 (2014).

4. L. Sun, M. G. Campbell, M. Dincă, Electrically conductive porous metal–organicframeworks. Angew. Chem. Int. Ed. 55, 3566–3579 (2016).

5. T. C. Narayan, T. Miyakai, S. Seki, M. Dinca, High charge mobility in a tetrathiafulvalene-basedmicroporous metal–organic framework. J. Am. Chem. Soc. 134, 12932–12935 (2012).

6. S. S. Park, E. R. Hontz, L. Sun, C. H. Hendon, Aron Walsh, T. V. Voorhis, M. Dincă,Cation-dependent intrinsic electrical conductivity in isostructural tetrathiafulvalene-basedmicroporous metal–organic frameworks. J. Am. Chem. Soc. 137, 1774–1777 (2015).

7. L. Sun, T. Miyakai, S. Seki, M. Dinca, Mn2(2,5-disulfhydrylbenzene-1,4-dicarboxylate): Amicroporous metal-organic framework with infinite (-Mn-S-)͚ chains and high intrinsiccharge mobility. J. Am. Chem. Soc. 135, 8185–8188 (2013).

8. A. F. Cozzolino, C. K. Brozek, R. D. Palmer, J. Yano, M. Li, M. Dincă, Ligand redoxnon-innocence in the stoichiometric oxidation of Mn2(2,5-dioxidoterephthalate)(Mn-MOF-74). J. Am. Chem. Soc. 136, 3334–3337 (2014).

9. L. Sun, C. H. Hendon, M. A. Minier, A. Walsh, M. Dincă, Million-fold electrical conductivityenhancement in Fe2(DEBDC) versus Mn2(DEBDC) (E = S, O). J. Am. Chem. Soc. 137,6164–6167 (2015).

10. A. A. Talin, A. Centrone, A. C. Ford, M. E. Foster, V. Stavila, P. Haney, R. A. Kinney, V. Szalai,F. E. Gabaly, H. P. Yoon, F. Léonard, M. D. Allendorf, Tunable electrical conductivity inmetal-organic framework thin-film devices. Science 343, 66–69 (2014).

11. P. D. C. Dietzel, Y. Morita, R. Blom, H. Fjellvåg, An in situ high-temperature single-crystalinvestigation of a dehydrated metal–organic framework compound and field-inducedmagnetization of one-dimensional metal–oxygen chains. Angew. Chem. Int. Ed. Engl. 44,6354–6358 (2005).

12. L. E. Darago, M. L. Aubrey, C. J. Yu, M. I. Gonzalez, J. R. Long, Electronic conductivity,ferrimagnetic ordering, and reductive insertion mediated by organic mixed-valence in aferric semiquinoid metal–organic framework. J. Am. Chem. Soc. 137, 15703–15711 (2015).

13. Y. Kobayashi, B. Jacobs, M. D. Allendorf, J. R. Long, Conductivity, doping, and redoxchemistry of a microporous dithiolene-based metal−organic framework. Chem. Mater. 22,4120–4122 (2010).

14. L. E. Kreno, K. Leong, O. K. Farha, M. Allendorf, R. P. Van Duyne, J. T. Hupp, Metal–organicframework materials as chemical sensors. Chem. Rev. 112, 1105–1125 (2012).

15. N. Ogihara, T. Yasuda, Y. Kishida, T. Ohsuna, K. Miyamoto, N. Ohba, Organic dicarboxylatenegative electrode materials with remarkably small strain for high-voltage bipolarbatteries. Angew. Chem. Int. Ed. Engl. 53, 11467–11472 (2014).

16. T. Yasuda, N. Ogihara, Reformation of organic dicarboxylate electrode materials forrechargeable batteries by molecular self-assembly. Chem. Commun. 50, 11565–11567 (2014).

17. N. Ogihara, Y. Ozawa, O. Hiruta, A self-assembled intercalated metal–organic frameworkelectrode with outstanding area capacity for high volumetric energy asymmetric capacitors.J. Mater. Chem. A 4, 3398–3405 (2016).

18. M. L. Aubrey, J. R. Long, A dual–ion battery cathode via oxidative insertion of anions in ametal–organic framework. J. Am. Chem. Soc. 137, 13594–13602 (2015).

19. M. Hmadeh, Z. Lu, Z. Liu, F. Gándara, H. Furukawa, S. Wan, V. Augustyn, R. Chang, L. Liao,F. Zhou, E. Perre, V. Ozolins, K. Suenaga, X. Duan, B. Dunn, Y. Yamamto, O. Terasaki,O. M. Yaghi, New porous crystals of extended metal-catecholates. Chem. Mater. 24,3511–3513 (2012).

20. T. Kambe, R. Sakamoto, T. Kusamoto, T. Pal, N. Fukui, K. Hoshiko, T. Shimojima, Z. Wang,T. Hirahara, K. Ishizaka, S. Hasegawa, F. Liu, H. Nishihara, Redox control and highconductivity of nickel bis(dithiolene) complex p-nanosheet: A potential organictwo-dimensional topological insulator. J. Am. Chem. Soc. 136, 14357–14360 (2014).

21. S. Chen, J. Dai, X. C. Zeng, Metal–organic Kagome lattices M3(2,3,6,7,10,11-hexaiminotriphenylene)2(M = Ni and Cu): From semiconducting to metallic by metal substitution. Phys. Chem.Chem. Phys. 17, 5954–5958 (2015).

22. M. Armand, S. Grugeon, H. Vezin, S. Laruelle, P. Ribière, P. Poizot, J.-M. Tarascon,Conjugated dicarboxylate anodes for Li-ion batteries. Nat. Mater. 8, 120–125 (2009).

23. A. Abouimrane, W. Weng, H. Eltayeb, Y. Cui, J. Niklas, O. Poluektova, K. Amine, Sodiuminsertion in carboxylate based materials and their application in 3.6 V full sodium cells.Energy Environ. Sci. 5, 9632–9638 (2012).

24. A. Choi, Y. K. Kim, T. K. Kim, M.-S. Kwon, K. T. Lee, H. R. Moon, 4,4′-Biphenyldicarboxylatesodium coordination compounds as anodes for Na-ion batteries. J. Mater. Chem. A 2,14986–14993 (2014).

25. A. Kuc, A. Enyashin, G. Seifert, Metal−organic frameworks: Structural, energetic,electronic, and mechanical properties. J. Phys. Chem. B 111, 8179–8186 (2007).

26. L. Valenzano, B. Civalleri, S. Chavan, S. Bordiga, M. H. Nilsen, S. Jakobsen, K. P. Lillerud,C. Lamberti, Disclosing the complex structure of UiO-66 metal organic framework:A synergic combination of experiment and theory. Chem. Mater. 23, 1700–1718 (2011).

27. W.-Q. Deng, W. A. Goddard III, Predictions of hole mobilities in oligoacene organicsemiconductors from quantum mechanical calculations. J. Phys. Chem. B 108, 8614–8621(2004).

28. M. Shindo, R. Koretsune, W. Yokota, K. Itoh, K. Shishido, Practical synthesis of ynolateanions: Naphthalene-catalyzed reductive lithiation of a,a-dibromo esters. TetrahedronLett. 42, 8357–8360 (2001).

8 of 9

http://advances.sciencemag.org/cgi/content/full/3/8/e1603103/DC1

http://advances.sciencemag.org/cgi/content/full/3/8/e1603103/DC1




ag.org/D

ownloaded from

29. H. Jung, S. J. Yang, T. Kim, J. H. Kang, C. R. Park, Ultrafast room-temperature reduction ofgraphene oxide to graphene with excellent dispersibility by lithium naphthalenide.Carbon 63, 165–174 (2013).

30. J. Zheng, H. Zhang, S. Dong, Y. Liu, C. T. Nai, H. S. Shin, H. Y. Jeong, B. Liu, K. P. Loh, Highyield exfoliation of two-dimensional chalcogenides using sodium naphthalenide. Nat.Commun. 5, 2995 (2014).

31. V. Thangadurai, R. A. Huggins, W. Weppner, Use of simple ac technique to determine theionic and electronic conductivities in pure and Fe-substituted SrSnO3 perovskites.J. Power Sources 108, 64–69 (2002).

32. J. Jamnik, Impedance spectroscopy of mixed conductors with semi-blocking boundaries.Solid State Ion. 157, 19–28 (2003).

33. C. Wang, J. Hong, Ionic/electronic conducting characteristics of LiFePO4 cathodematerials: The determining factors for high rate performance. Electrochem. Solid StateLett. 10, A65–A69 (2007).

34. C. Avendano, Z. Zhang, A. Ota, H. Zhao, K. R. Dunbar, Dramatically different conductivityproperties of metal–organic framework polymorphs of Tl(TCNQ): An unexpectedroom-temperature crystal-to-crystal phase transition. Angew. Chem. Int. Ed. 50, 6543–6547(2011).

35. J. Ma, C. Wang, S. Wroblewski, Kinetic characteristics of mixed conductive electrodes forlithium ion batteries. J. Power Sources 164, 849–856 (2007).

36. J. Lu, G. Oyama, S.-i. Nishimura, A. Yamada, Increased conductivity in the metastableintermediate in LixFePO4 electrode. Chem. Mater. 28, 1101–1106 (2016).

37. D. Young, A. Ransil, R. Amin, Z. Li, Y.-M. Chiang, Electronic conductivity in the Li4/3Ti5/3O4–Li7/3Ti5/3O4

system and variation with state-of-charge as a Li battery anode. Adv. Energy Mater. 3,1125–1129 (2013).

38. G. Sahu, E. Rangasamy, J. Li, Y. Chen, K. An, N. Dudney, C. Liang, A high-conduction Gesubstituted Li3AsS4 solid electrolyte with exceptional low activation energy. J. Mater.Chem. A 2, 10396–10403 (2014).

39. Z. Liu, W. Fu, E. A. Payzant, X. Yu, Z. Wu, N. J. Dudney, J. Kiggans, K. Hong, A. J. Rondinone,C. Liang, Anomalous high ionic conductivity of nanoporous b-Li3PS4. J. Am. Chem. Soc.135, 975–978 (2013).

40. N. Kamaya, K. Homma, Y. Yamakawa, M. Hirayama, R. Kanno, M. Yonemura, T. Kamiyama,Y. Kato, S. Hama, K. Kawamoto, A. Mitsui, A lithium superionic conductor. Nat. Mater. 10,682–686 (2011).

41. D. H. C. Wong, J. L. Thelen, Y. Fu, D. Devaux, A. A. Pandya, V. S. Battaglia, N. P. Balsara,J. M. DeSimone, Nonflammable perfluoropolyether-based electrolytes for lithiumbatteries. Proc. Natl. Acad. Sci. U.S.A. 111, 3327–3331 (2014).

42. J. Wang, Y. Yamada, K. Sodeyama, C. H. Chiang, Y. Tateyama, A. Yamada,Superconcentrated electrolytes for a high-voltage lithium-ion battery. Nat. Commun. 7,12032 (2016).

43. K. Liu, W. Liu, Y. Qiu, B. Kong, Y. Sun, Z. Chen, D. Zhuo, D. Lin, Y. Cui, Electrospun core-shellmicrofiber separator with thermal-triggered flame-retardant properties for lithium-ionbatteries. Sci. Adv. 3, e1601978 (2017).

44. W. Ji, F. Wang, D. Liu, J. Qian, Y. Cao, Z. Chen, H. Yang, X. Ai, Building thermally stableLi-ion batteries using a temperature-responsive cathode. J. Mater. Chem. A 4,11239–11246 (2016).

45. Z. Chen, P.-C. Hsu, J. Lopez, Y. Li, J. W. F. To, N. Liu, C. Wang, S. C. Andrews, J. Liu, Y. Cui,Z. Bao, Fast and reversible thermoresponsive polymer switching materials for saferbatteries. Nat. Energy 1, 15009 (2016).

46. J.-F. Zou, Z.-Z. Yu, Y.-X. Pan, X.-P. Fang, Y.-C. Ou, Conductive mechanism of polymer/graphiteconducting composites with low percolation threshold. J. Polym. Sci. Part B Polym. Phys.40, 954–963 (2002).

47. B. Abeles, H. L. Pinch, J. I. Gittleman, Percolation conductivity in W-Al2O3 granular metalfilms. Phys. Rev. Lett. 35, 247–250 (1975).

48. J. Zhang, M. Mine, D. Zhu, M. Matsuo, Electrical and dielectric behaviors and their originsin the three-dimensional polyvinyl alcohol/MWCNT composites with low percolationthreshold. Carbon 47, 1311–1320 (2009).


49. N. Johner, C. Grimaldi, I. Balberg, P. Ryser, Transport exponent in a three-dimensionalcontinuum tunneling-percolation model. Phys. Rev. B 77, 174204 (2008).

50. N. Johner, C. Grimaldi, T. Maeder, P. Ryser, Optimal percolation of disordered segregatedcomposites. Phys. Rev. E 79, 020104 (2009).

51. G. Kresse, J. Furthmüller, Efficient iterative schemes for ab initio total-energy calculationsusing a plane-wave basis set. Phys. Rev. B 54, 11169–11186 (1996).

52. J. P. Perdew, A. Ruzsinszky, G. I. Csonka, O. A. Vydrov, G. E. Scuseria, L. A. Constantin,X. Zhou, K. Burke, Restoring the density-gradient expansion for exchange in solids andsurfaces. Phys. Rev. Lett. 100, 136406 (2008).

53. J. Heyd, G. E. Scuseria, M. Ernzerhof, Hybrid functionals based on a screened Coulombpotential. J. Chem. Phys. 118, 8207–8215 (2003).

54. J. Heyd, G. E. Scuseria, M. Ernzerhof, Erratum: “Hybrid functionals based on a screenedCoulomb potential” [J. Chem. Phys. 118, 8207 (2003)]. J. Chem. Phys. 124, 219906 (2006).

55. H. Li, R. Zheng, Q. Shi, Theoretical study of charge carrier transport in organicsemiconductors of tetrathiafulvalene derivatives. J. Phys. Chem. C 116, 11886–11894(2012).

56. H. Kobayashi, N. Kobayashi, S. Hosoi, N. Koshitani, D. Murakami, R. Shirasawa, Y. Kudo,D. Hobara, Y. Tokita, M. Itabashi, Hopping and band mobilities of pentacene, rubrene, and2,7-dioctyl[1]benzothieno[3,2-b][1]benzothiophene (C8-BTBT) from first principlecalculations. J. Chem. Phys. 139, 014707 (2013).

57. E. F. Valeev, V. Coropceanu, D. A. da Silva Filho, S. Salman, J.-L. Brédas, Effect of electronicpolarization on charge-transport parameters in molecular organic semiconductors. J. Am.Chem. Soc. 128, 9882–9886 (2006).

58. B. Baumeier, J. Kirkpatrick, D. Andrienko, Density-functional based determination ofintermolecular charge transfer properties for large-scale morphologies. Phys. Chem.Chem. Phys. 12, 11103–11113 (2010).

59. M. J. Frisch, G. W. Trucks, H. B. Schlegel, G. E. Scuseria, M. A. Robb, J. R. Cheeseman,G. Scalmani, V. Barone, B. Mennucci, G. A. Petersson, H. Nakatsuji, M. Caricato, X. Li,H. P. Hratchian, A. F. Izmaylov, J. Bloino, G. Zheng, J. L. Sonnenberg, M. Hada, M. Ehara,K. Toyota, R. Fukuda, J. Hasegawa, M. Ishida, T. Nakajima, Y. Honda, O. Kitao, H. Nakai,T. Vreven, J. A. Montgomery Jr., J. E. Peralta, F. Ogliaro, M. Bearpark, J. J. Heyd, E. Brothers,K. N. Kudin, V. N. Staroverov, T. Keith, R. Kobayashi, J. Normand, K. Raghavachari,A. Rendell, J. C. Burant, S. S. Iyengar, J. Tomasi, M. Cossi, N. Rega, J. M. Millam, M. Klene,J. E. Knox, J. B. Cross, V. Bakken, C. Adamo, J. Jaramillo, R. Gomperts, R. E. Stratmann,O. Yazyev, A. J. Austin, R. Cammi, C. Pomelli, J. W. Ochterski, R. L. Martin, K. Morokuma,V. G. Zakrzewski, G. A. Voth, P. Salvador, J. J. Dannenberg, S. Dapprich, A. D. Daniels,O. Farkas, J. B. Foresman, J. V. Ortiz, J. Cioslowski, D. J. Fox, Gaussian 09, Revision B. 01(Gaussian Inc., 2010).

Acknowledgments: We thank S. Kosaka for ICP-OES measurements as well as H. Nakano andM. Ohashi for heat treatment of the samples. XRD measurements were performed at the BL5S2beamline, Aichi Synchrotron Radiation Center, Aichi prefecture, Japan (proposal no.201601031). Funding: This work was supported by the Toyota Central R&D Laboratories Inc.Author contributions: N. Ogihara conceived, designed, and performed the experiments andprepared the manuscript. Y.K. performed the XRD measurements. N. Ohba performed thecalculations for the band structures and hopping conduction. Competing interests: Theauthors declare that they have no competing interests. Data and materials availability: Alldata needed to evaluate the conclusions in the paper are present in the paper and/or theSupplementary Materials. Additional data related to this paper may be requested from the authors.

Submitted 7 December 2016Accepted 26 July 2017Published 25 August 201710.1126/sciadv.1603103

Citation: N. Ogihara, N. Ohba, Y. Kishida, On/off switchable electronic conduction inintercalated metal-organic frameworks. Sci. Adv. 3, e1603103 (2017).

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On/off switchable electronic conduction in intercalated metal-organic frameworksNobuhiro Ogihara, Nobuko Ohba and Yoshihiro Kishida

DOI: 10.1126/sciadv.1603103 (8), e1603103.3Sci Adv

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