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A “nanopore lithography” strategy towards hierarchically micro/mesoporous carbon from ZIF-8/graphene oxide hybrids for electrochemical energy storage Francisco Julian Martin-Jimeno, Fabian Suárez-García, Juan Ignacio Paredes, Marina Enterria, Manuel Fernando R. Pereira, José Inácio Martins, José L. Figueiredo, Amelia Martinez-Alonso, and Juan Manuel D. Tascon ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b16567 • Publication Date (Web): 07 Dec 2017 Downloaded from http://pubs.acs.org on December 10, 2017

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ACS Applied Materials & Interfaces

A “nanopore lithography” strategy towards hierarchically micro/mesoporous carbon from ZIF-8/graphene oxide hybrids for electrochemical energy storage F.J. Martín-Jimenoa, F. Suárez-Garcíaa, J.I. Paredesa*, M. Enterríab, M.F.R. Pereirab, J.I. Martinsc,d, J.L. Figueiredob, A. Martínez-Alonsoa, J.M.D. Tascóna a

Instituto Nacional del Carbón, INCAR-CSIC, C/Francisco Pintado Fe 26, 33011 Oviedo Spain *

b

E-mail address: [email protected]

Laboratório de Processos de Separação e Reacção – Laboratório de Catálise e

Materiais (LSRE-LCM), Departamento de Engenharia Química, Faculdade de Engenharia, Universidade de Porto, R. Dr. Roberto Frias s/n, 4200-465 Porto, Portugal c

Departamento de Engenharia Química, Faculdade de Engenharia, Universidade de Porto, R. Dr. Roberto Frias s/n, 4200-465 Porto, Portugal

d

LAB2PT-Laboratório de Paisagens, Património e Território, Universidade do Minho, 4710-057 Braga, Portugal

Keywords: Porous carbon; Metal-organic framework; Mesoporosity; Graphene oxide; Electrochemical energy storage 1 ACS Paragon Plus Environment

ACS Applied Materials & Interfaces 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Abstract Porous carbons derived from metal-organic frameworks (MOFs) are promising materials towards a number of energy- and environment-related applications, but their almost exclusively microporous texture can be an obstacle to their performance in practical uses. Here, we introduce a novel strategy for the generation of very uniform mesoporosity in a prototypical MOF, namely the zeolitic imidazolate framework-8 (ZIF-8). The process, which is referred to as “nanopore lithography”, makes use of graphene oxide (GO) nanosheets enclosing ZIF-8 particles as masks or templates for the transfer of mesoporous texture to the latter. Upon controlled carbonization and activation steps, nanopores created in the GO envelope serve as selective entry points for the localized etching of the carbonized ZIF-8, so that such nanopores are replicated in the MOF-derived carbonaceous structure. The resulting porous carbons are dominated by uniform mesopores ~3-4 nm in width and possess specific surface areas of ~1300-1400 m2 g-1. Furthermore, we investigate and discuss the specific experimental conditions that afford the mesopore templating action of the GO nanosheets. Electrochemical characterization revealed an improved capacitance as well as a faster, more reversible charge/discharge kinetics for the ZIF-8 derived porous carbons obtained through nanopore lithography, relative to their counterparts with standard activation (no GO templating), thus indicating the potential practical advantage of the present approach in capacitive energy storage applications.

1. Introduction Boasting such attractive characteristics as high surface area, large pore volume, good electrical conductivity and chemical/mechanical stability, as well as a wide availability of sources for their production, porous carbons play a prominent role in the ongoing efforts to advance a number of key technologies that will contribute to securing a sustainable future.1 These technologies encompass, for example, the fields of electrochemical energy storage (electrodes for supercapacitors, Li-ion or Li-S batteries, etc),2,3 gas adsorption (storage/separation of energy- and environment-relevant gases, including H2, CO2 and CH4),3,4 or catalysis (e.g., catalysts and catalyst supports for the synthesis of industrially relevant organic compounds or for energy conversion processes in fuel cells, such as the oxygen reduction reaction).5,6

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Although many approaches for the preparation of porous carbons have been developed to this day,7 in recent years researchers have paid particular attention to both direct and indirect methods based on the use of metal-organic frameworks (MOFs) as templates and precursors of the carbon material.8-12 MOFs are crystalline solids formed through coordination of metal ions (Zn2+, Cu2+, Co2+, etc) with certain organic ligands (e.g., terephthalic acid, trimesic acid or 2-methylimidazole), and often exhibit a permanent porosity, typically in the micropore range (< 2 nm), that is useful for a broad range of practical applications.13 High-temperature carbonization (pyrolysis) of such porous MOFs as MOF-5, MOF-74, ZIF-8, ZIF-69 or Al-PCP, either alone or in the presence of certain infiltrated organic compounds acting as a secondary carbon source, has been demonstrated to afford highly porous carbons, the porous texture of which can be further developed through a subsequent activation step (usually with KOH as the activating agent).10,11 There are a number of important assets associated to the use of these coordination networks as precursors for porous carbons, including (1) the ability to generate a spatially uniform and homogeneous porosity in the carbon material inherited from the high crystallinity of the MOF precursor, (2) the possibility to readily incorporate (also with spatial uniformity) metal nanoparticles or clusters derived from the coordinating ions, and (3) the prospect to access porous carbons doped with heteroatoms (e.g., nitrogen) by choosing organic ligands that incorporate the desired element.11 These appealing features have triggered many studies aimed at exploring the potential of MOF-derived porous carbons towards different applications, mainly in the field of electrochemical energy storage and conversion.14-19 On the other hand, it is well known that carbons obtained from MOFs tend to be mostly microporous in nature, even if the carbonized MOF precursor is subjected to an activation process.10,11,20-23 Unfortunately, the absence of a sizable proportion of larger pores (especially mesopores) is detrimental to the performance of these materials in many of their target applications (e.g., electrodes for supercapacitors or electrocatalysts for oxygen reduction), as the transport of molecules and ions along the narrow microporous passages is expected to be sluggish. To address this question, a number of novel strategies conducive to generating a hierarchical micro/mesoporosity in MOFderived carbons have been recently proposed and investigated, mainly using ZIF-8 as the precursor. Such strategies are based on (1) synthesis of the MOF precursor in an ultrasonic environment,24 (2) carbonization of the MOF with simultaneous solvent evaporation,25 (3) nanowire- or nanotube-templated growth of the MOF,26-28 (4) 3 ACS Paragon Plus Environment


carbonization of the MOF under polymer confinement,29,30 (5) high-temperature etching of the carbonized MOF in NH3 atmosphere,31 or (6) use of heteroatom-rich organic ligands in the MOF to generate a large amount of gases (e.g., NO2 or CO2) upon pyrolysis that will trigger pore expansion in the remaining carbon framework.32,33 However, these approaches usually lead to carbons dominated by large mesopores (> 10 nm) and broad mesopore size distributions that even extend to the macropore range. Attaining MOF-derived carbons with uniform-sized mesopores, particularly in the 2-10 nm range, would be highly desirable, as it would afford materials with more controlled characteristics,34,35 but for the most part such a goal has remained elusive. In order to overcome this issue, here we propose and implement a novel concept, which we have termed “nanopore lithography”, for the generation of uniform mesoporosity in MOF-derived porous carbons. Such a strategy is based on carbonization and subsequent activation of MOF particles sheathed by porous graphene oxide (GO) nanosheets, where the latter act as masks or templates for the transfer of mesoporous texture to the carbonized MOF particles during the activation step. As will be shown below, this process leads to high surface area porous carbons dominated by a uniform mesoporosity (~3-4 nm in width), which is created by the nanopore lithography process, coexisting with the typical microporosity of carbon materials obtained from MOFs. Furthermore, to explore the potential applications of these novel hierarchically porous carbons, we have investigated their electrochemical performance with a view to their use as electrodes for electric double-layer capacitors. Specifically, a combination of conventional electrochemical analysis (i.e., cyclic voltammetry and electrochemical impedance spectroscopy) and potential step voltammetry enabled a fine characterization of the pseudocapacitive storage phenomena taking place on the porous carbon electrodes. The results indicated that such carbons exhibit a better performance than that of their counterparts prepared by carbonization and activation of the MOF in the absence of GO nanosheets, thus highlighting the practical advantage of the proposed strategy.

2. Results and discussion 2.1. Rationale of the use of activated graphene oxide as a mask for nanopore lithography Fig. 1 shows a schematic illustration of the rationale behind the generation of uniform mesopores in MOF-based carbon by the nanopore lithography approach 4 ACS Paragon Plus Environment

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proposed here, which makes use of GO nanosheets as masks or templates and ZIF-8 as the MOF precursor. GO is a well-known, highly oxidized derivative of graphene, the carbon lattice of which is heavily decorated with different types of oxygen functional groups, mainly hydroxyls and epoxides on its basal plane and carbonyls and carboxylic acids along its edges.36 On the nanometer scale, the spatial distribution of these functional groups is quite uneven, so that the GO nanosheet can be viewed as a fine patchwork of domains with a very high density of oxygen functionalities coexisting with regions essentially devoid of oxygen.37 Such a structural disparity implies that the oxygen-rich domains must be much more chemically reactive than their oxygen-free counterparts. Indeed, extensive studies have shown that reaction of both GO and reduced graphene oxide (RGO, a less oxidized version of GO) with oxidizing agents (e.g., H2O2 at moderate temperature, O2/air at high temperature or oxygen plasma) leads to preferential etching of their strongly oxidized domains, leaving behind graphene nanosheets punctured by large numbers of nanometer-sized holes.38-41 Similarly, heat treatment of GO and RGO under vacuum or inert gases generates atomic vacancies/tiny holes (i.e., incipient porosity) on the nanosheets as a result of decomposition of the oxidized domains into CO and CO2.42,43 This behavior has inspired the development of a novel family of porous carbons produced by activation of GO/RGO, mostly (although not exclusively) using KOH as the activating agent.44-48 If the process is carried out under mild conditions (e.g., at relatively low temperatures and/or KOH/precursor weight ratios), the activation reactions can be limited to creating discrete holes (nanopores) in the two-dimensional nanosheets (Fig. 1a), thus avoiding their overetching and conversion into a three-dimensional porous carbon structure.49 As shown in Fig. 2, chemical activation of GO nanosheets with KOH under controlled conditions afforded a mildly activated graphene-based carbon material that retained the two-dimensional, lamellar morphology of its precursor, possessed a moderate surface area with a porous texture dominated by micropores and small mesopores, and exhibited a well-developed graphitic character. As demonstrated by atomic force microscopy (see Supporting Information), the individual starting GO nanosheets were single-layered objects with typical lateral sizes between 100 and 500 nm. The preparation of the activated product from these nanosheets involved a heat treatment step at 700 ºC under inert atmosphere followed by chemical activation proper (e.g., at 800 ºC) in the presence of KOH (see Experimental section for details). To prevent the graphene lattice from an excessive attack during activation as well as for 5 ACS Paragon Plus Environment


other reasons that will be discussed below, a moderate amount of KOH was used (KOH/precursor weight ratio of 1). We note that such an amount was much lower than that commonly reported in the preparation of high surface area activated graphene materials (typical KOH/precursor weight ratios between 4 and 9).44,45,47,48 Field emission scanning electron microscopy (FE-SEM) imaging (Fig. 2a) indicated that the mildly activated material was comprised of relatively loose collections of twodimensional nanosheets, the general morphology of which was essentially indistinguishable from that of their non-activated precursor (images not shown). As revealed by transmission electron microscopy (TEM), the nanosheets appeared to be decorated by small holes with sizes below 5 nm (see, e.g., some holes labeled by green circles in Fig. 2b). N2 adsorption/desorption measurements of the activated (800 ºC) GO nanosheets at -196 ºC (Fig. 2c, red trace) yielded isotherms with a combination of types I and IV, as well as an H2 hysteresis loop, which denote the presence of micropores and small sizedmesopores [Brunauer-Emmett-Teller (BET) surface area of 566 m2 g-1].50 Indeed, the corresponding pore size distribution (PSD) obtained by the quenched solid density functional theory (QSDFT) method (inset to Fig. 2c) evinced that ultramicropores (< 1 nm in size) and supermicropores (up to around 2 nm) coexisted with a large amount of small mesopores of well-defined sizes (~3-4 nm, with a peak maximum in the PSD at ~3.4 nm). This porous texture was mostly developed during the activation step at 800 ºC, as the material prior to activation proper (i.e., only heat-treated at 700 ºC in the absence of KOH) lacked significant porosity (see isotherm drawn in black in Fig. 2c). The typical Raman spectrum of the mildly activated material is presented in Fig. 2d (red trace) together with that of the starting GO (black trace), which is shown for comparison purposes. The spectrum was dominated by three main features that are characteristic of graphitic carbons.

51,52

namely, the G band (~1582 cm-1), the defect-related D band

(~1348 cm-1) and the second-order 2D band (~2692 cm-1). The so-called D´ band (~1612 cm-1), which is also defect-related, was observed as a shoulder on the high wavenumber side of the G band. We note that the D and G bands of the activated GO exhibited a relatively sharp and narrow appearance, suggestive of a reasonably welldeveloped graphitic character. By contrast, these bands were broader and partially overlapped with each other in the case of the starting GO, pointing to a more disordered structure. Furthermore, the 2D band for the latter was broad and weak, whereas it became much more intense and sharper after activation, which is also a sign of 6 ACS Paragon Plus Environment

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increased overall graphitic order in the activated GO.42 We interpret this result to be mainly a consequence of the selective removal of highly oxidized and disordered domains from GO leading to the generation of small-sized holes in the nanosheets (as seen by TEM, Fig. 2b), so that the carbon structures left behind in the activated product are on average significantly more graphitic and less disordered than their starting GO counterpart. It is also worth noting that activation of GO nanosheets at different temperatures (e.g., 600 and 1000 ºC) led to the generation of mesopores of the same size as those generated at 800 ºC (see Supporting Information). We believe this result to be related to the high reactivity of the extensively oxidized domains present in the nanosheets, so that reaction with KOH at 600 ºC suffices to remove single domains completely, whereas higher activation temperatures cannot enlarge the resulting holes (mesopores) due to the graphitic, unreactive nature of the carbon lattice surrounding the holes. Based on this knowledge, we hypothesized that enclosing ZIF-8 particles by GO nanosheets would provide the former with a protective sheath that could fundamentally change the process of their chemical activation by KOH (Fig. 1b). More specifically, access of the KOH reagent to the (carbonized) ZIF-8 particles would take place selectively through the nanopores created by the activation reactions in the GO envelope, rather than indiscriminately through the whole external surface of the particles (the latter scenario would be avoided by the graphitic, unreactive domains of the activated GO nanosheets acting as a barrier). Under such conditions, those sections of the carbonized ZIF-8 particles located immediately adjacent to the GO nanopore openings would be readily gasified by the very large local influx of KOH (very high local ratios of KOH to carbonized ZIF-8),53 thus inducing the generation of porosity in the ZIF-8-derived carbon with dimensions dictated by those of the GO nanopores that served as templates. If the chemical activation process creates nanopores in the GO nanosheets different in size to those typically generated by carbonization/activation of the stand-alone ZIF-8 particles (e.g., mesopores), then we should expect such a porosity to be translated to the GO-sheathed ZIF-8 particles upon activation.

2.2. Generating mesoporosity in ZIF-8-derived carbons with sheathing graphene oxide nanosheets as nanopore masks To test the idea of using activated GO as a mask for the generation of mesoporosity in ZIF-8-derived carbons, we prepared ZIF-8/GO hybrids in which the ZIF-8 particles 7 ACS Paragon Plus Environment


were mostly enclosed by the GO nanosheets, and then such hybrids were subjected to carbonization and KOH activation. Although the combination of ZIF-8 particles with GO nanosheets has been previously investigated, in most cases researchers have focused on using the nanosheets as a support for the particles in “naked” form.54-57 However, albeit less commonly reported, the generation of GO-sheathed ZIF-8 particles is also possible.58 In our case, the ZIF-8/GO hybrids were prepared by coordination of Zn2+ ions with 2-methylimidazole (Hmim) in a mixed water/methanol solution (1/9 vol/vol) in the presence of GO nanosheets as well as polyvinylpyrrolidone (PVP). The latter was mainly included to promote the interaction of the ZIF-8 particles with GO,54 so that they could be efficiently wrapped by the submicrometric nanosheets. Although different concentrations and/or ratios of these chemical species and reagents were tested (see Supporting Information), optimal results were obtained using a GO concentration in the water/methanol medium of 0.5 mg mL-1, Hmim/PVP/GO weight ratio of 50/7.5/1 and Hmin/Zn molar ratio of 7.5/1 (see Experimental section for details on the synthesis). Fig. 3a and b shows typical FE-SEM images of the optimized ZIF-8/GO hybrid. On a large scale, the material exhibited a homogeneous and slightly rough appearance (Fig. 3a). Upon closer inspection, polygonal objects ascribed to ZIF-8 particles became more readily apparent (Fig. 3b). These particles were typically around 100 nm in size and were veiled by very thin lamellar objects that were attributed to the GO nanosheets (see TEM image in Fig. 3c). By contrast, when the synthesis was carried out in the absence of GO only “naked” particles of a larger size were obtained (Fig. 3d). As noticed from Fig. 3e, X-ray diffraction (XRD) confirmed that these particles (both the “naked” ones and those hybridized with GO) were indeed ZIF-8 crystals.54 Pyrolysis of the optimized ZIF-8/GO hybrid at 700 ºC under inert atmosphere followed by activation of the resulting carbonized product with KOH afforded hierarchically porous carbons with quantitative features that depended on the particular activation conditions but with a qualitatively invariant mesoporous texture coincident with that present in the activated GO nanosheets. The pyrolysis step was carried out to generate a carbonaceous structure from the ZIF-8 particles that was suitable for activation. Thermogravimetric analysis under nitrogen flow indicated that pyrolysis of ZIF-8 started at ~550 ºC (see Supporting Information). In order to attain a carbonized ZIF-8 material with a highly disordered or even amorphous structure that was readily amenable to activation (etching), we selected a carbonization temperature as close as possible to this onset temperature (i.e., 700 ºC). Pyrolysis at higher temperatures (e.g., 8 ACS Paragon Plus Environment

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800, 900 or 1000 ºC) could be expected to give rise to more graphitic, less reactive materials that are harder to etch away upon activation. The pyrolyzed ZIF-8/GO hybrid was a low surface area (~250 m2 g-1) carbon material that exhibited an incipient porosity, both microporosity and mesoporosity. Most likely, such an incipient porosity was largely generated by the (partial) thermal etching of the highly oxidized domains present in the GO nanosheets43 and should serve as the nucleating sites for the subsequent enlargement of porosity in the nanosheets and the localized etching of the carbonized ZIF-8 particles upon KOH activation (see Fig. 1). Due to their poor porosity development, the pyrolyzed-only, non-activated ZIF-8/GO hybrids were of limited interest with a view to their use in molecular or ionic adsorption processes (e.g., in capacitive energy storage). We also note that when the ZIF-8 particles themselves, rather than their carbonized counterparts, were subjected to KOH activation, a significantly lower extent of porosity development was observed in the resulting carbon products. Such an outcome was probably related to the instability of this MOF under the extremely basic conditions generated by the activating agent. Fig. 4a and b shows representative FE-SEM images on the submicrometric scale of the activated hybrids prepared at 800 (a) and 1000 (b) ºC with a KOH/precursor weight ratio of 1. These samples, as well as those prepared at activation temperatures of 600, 700 and 900 ºC (images not shown), were characterized by a uniform, network-like morphology with nanometer-sized voids. Such a morphology was obviously different to that of the porous carbons prepared by activation of ZIF-8 particles in the absence of GO (e.g., Fig. 4c for ZIF-8 activated at 800 ºC with a KOH/precursor weight ratio of 1), but also differed from the markedly lamellar make-up of the material obtained by activation of GO alone (Fig. 2a). As disclosed by TEM (Fig. 4d), on a smaller scale the activated hybrids showed the presence of dark objects that were similar to those previously reported for carbons derived from ZIF-8/GO hybrids and could be ascribed to the carbon component inherited from the ZIF-8 particles.54-56 Furthermore, close inspection of the images revealed the presence of (mesoporous) voids in these objects with sizes around 3-4 nm (e.g., green circles in Fig. 4d). The presence of a large component directly derived from the ZIF-8 particles in the activated materials obtained from the ZIF-8/GO hybrid was also made apparent from Raman spectroscopy measurements (Fig. 4e). While the activated GO-only carbon displayed a reasonably well-developed graphitic structure (Fig. 2d), the carbon lattice of the ZIF-8/GO hybrids activated at different temperatures (e.g, 600, 800 and 1000 ºC 9 ACS Paragon Plus Environment


with a KOH/precursor ratio of 1) was much more disordered, as attested by their broader and overlapping D and G bands as well as by their very weak (or even negligible) 2D band. In fact, such features were mostly characteristic of the carbon structures inherited from the MOF component, as exemplified by the Raman spectrum of ZIF-8 particles grown in the absence of GO and activated at 800 ºC (KOH/precursor ratio of 1) that is shown also in Fig. 4e. Referring specifically to the samples shown in Fig. 4e, their measured values of the integrated intensity ratio of the D and G bands (ID/IG ratio) were ~2.1, 1.7, 1.7 and 1.3 for ZIF-8 (1|800), ZIF-8/GO (1|600), ZIF-8/GO (1|800) and ZIF-8/GO (1|1000), respectively, whereas the value for GO activated in the absence of ZIF-8 (sample GO (1|800), Fig. 2d) was ~0.9. The lower ID/IG value determined for the activated GO-only material was consistent with its expected well developed graphitic character, as mentioned above. By contrast, the large ID/IG value of the activated ZIF-8 sample (no GO) was indicative of a much more disordered (even amorphous) structure, which in turn was in agreement with what can be expected for a carbon material derived from a non-aromatic organic precursor (i.e., Hmim). As could also be anticipated, the ID/IG values for the activated ZIF-8/GO hybrids were in-between those of their individual GO- and ZIF-8-derived components. Furthermore, we believe that the structural imperfections responsible for the emergence of the D band were of a different nature in the case of the GO-derived and ZIF-8-derived components. For the former, the structural defects in the carbon lattice were most likely dominated by the large fraction of internal edges associated to the nanometer-sized holes created in the nanosheets upon activation with KOH, while for activated ZIF-8 many defect sites can be ascribed to sp3 centers derived from its organic Hmim precursor. In any case, we can conclude that the structure of the activated hybrids was dominated to a large extent by the carbon component directly stemming from the ZIF-8 particles. The N2 adsorption/desorption isotherms and the corresponding PSDs for ZIF-8/GO hybrids activated at different temperatures (600, 700, 800, 900 and 1000 ºC) as well as for ZIF-8 (no GO) activated at 800 ºC are presented in Fig. 5. Similar to the case of the activated material derived from GO alone (Fig. 2c), the activated hybrids displayed a combination of type I and type IV isotherms with a large H2-type hysteresis loop (for samples activated at 700 ºC and above), suggesting a large contribution of uniform mesopores to their porous texture in addition to the presence of micropores (Fig. 5a). On the other hand, the activated ZIF-8 sample (no GO) clearly exhibited a different adsorption behavior, namely, a type Ib isotherm (Fig. 5b) that almost lacked hysteresis, 10 ACS Paragon Plus Environment

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which indicated that supermicropores (sizes up to around 2 nm) dominated the porous texture of this material and larger pores were largely absent. The BET surface areas of the activated hybrids were around 1200-1300 m2 g-1 except for the material activated at 600 ºC (383 m2 g-1), which exhibited a rather limited porosity development (see Table 1 for their textural parameters). These BET values were lower than those obtained for the activated materials prepared from ZIF-8 alone (e.g., 4234 m2 g-1 for ZIF-8 activated at 800 ºC; Table 1), but were nonetheless comparable to (or even larger than) those of previously reported ZIF-8-derived porous carbons with hierarchical micro/mesoporosity generated by different strategies.27-32 Concerning the PSDs (Fig. 5c), while the activated materials derived from ZIF-8 alone were dominated by supermicropores, the most salient feature of the activated hybrids was the presence of very uniform mesoporosity with sizes between 3 and 4 nm. Significantly, the pore volume associated to such mesopores increased dramatically, both in absolute and relative terms (i.e., relative to the total pore volume), with activation temperature up to 900 ºC and then started to decrease at 1000 ºC (Table 1). We note that the size of the mesopores generated by activation of the ZIF-8/GO hybrids (Fig. 5c) and the GO nanosheets alone (inset to Fig. 2c) were essentially indistinguishable. Furthermore, comparison of the mesopore volumes of the hybrid and the GO nanosheets activated at the same temperature (e.g. 800 ºC, mesopore volumes of 0.795 vs 0.318 cm3 g-1; see Table 1) strongly suggested that the mesopore volume arising from the carbon component of the hybrid directly derived from the ZIF-8 particles must be quite large (i.e., larger than 0.795 cm3 g-1), indicating that mesopore generation in such a component was achieved. The presence of ~3-4 nm wide mesopores in this carbon component was indeed directly visualized by TEM (see Fig. 4d). Because such a uniform and well-defined mesoporosity was not generated through activation of the ZIF-8 particles alone, we have to conclude that its emergence in the carbon component derived from the ZIF-8 particles in their hybrids with GO nanosheets must be a consequence of the mesopore-templating action of the nanosheets, as schematically depicted in Fig. 1b. We stress that activation of the ZIF-8/GO hybrids was carried out using a KOH/precursor weight ratio of 1, which is a rather modest value compared with the ratios that have been typically employed not only in the activation of GO nanosheets, but also of ZIF-8 particles alone. For example, ratios of 4-5 are common in the production of activated carbons from the latter.20-22 In our case, such high amounts of KOH were seen to have a negative impact on the porous texture development of the 11 ACS Paragon Plus Environment


activated ZIF-8/GO hybrids. Fig. 6 compares the N2 adsorption/desorption isotherms (a) and their corresponding PSDs (b) for hybrids activated at 800 ºC using different KOH/precursor weight ratios (1, 2, 3 and 4), with their textural parameters being given in Table 1. The samples activated with a weight ratio of 1 and 2 exhibited similar BET surface areas (~1300-1350 m2 g-1) and total pore volumes (~1.1-1.2 cm3 g-1), although the abundance of templated mesopores ~3-4 nm in size was noticeably lower for the carbon with a weight ratio of 2. On the other hand, the use of higher KOH/precursor ratios (i.e., 3 and 4) led to a remarkable decrease in BET surface areas (down to ~200 m2 g-1) and total pore volumes (< 0.2 cm3 g-1). We interpret this result to arise from the particular mode of activation of the ZIF-8 particles wrapped by GO nanosheets. Because access of KOH to the carbonized ZIF-8 particles is expected to take place only through the holes present in the sheathing GO nanosheets, there will be a large influx of activating agent at those entry points, whereas those regions on the surface of the particles protected by the graphene lattice will be mostly kept from reacting with KOH (see Fig. 1b). This situation will generate a large spatial heterogeneity in the etching of each individual particle during activation (unlike the case of stand-alone, non-sheathed particles, for which spatially homogeneous and uniform etching is to be expected). For moderate amounts of activating agent, localized etching of the particle can be expected to generate pores similar in size to the holes in the GO sheets that served as entry points for the KOH reactant, thus giving rise to carbon materials with uniform, well-developed mesoporosity (e.g. samples with KOH/precursor ratios of 1 and 2 in Fig. 6). However, if the amount of KOH used for activation is too large (e.g., KOH/precursor ratios of 3 and 4), then an excessive supply of the activating agent at the entry points will probably lead to the indiscriminate removal of large portions of the particle starting from such entry points, and so to a collapse of the porous texture in the resulting carbon. Consequently, the optimal activation conditions for GO-sheathed ZIF-8 particles involve the use of relatively low amounts of activating agent. The surface chemistry characteristics of the activated hybrids were investigated by X-ray photoelectron spectroscopy (XPS). Fig. 7 shows high resolution N 1s (a,c,e,g) and O 1s (b,d,f,h) core-level spectra for ZIF-8/GO hybrids activated at 600 (a,b), 800 (c,d) and 1000 (e,f) ºC, as well as for ZIF-8 alone activated at 800 ºC (g,h). The corresponding spectra for the hybrids activated at 700 and 900 ºC are given in the Supporting Information (Fig. S6), as well as the C 1s core-level spectra of all the activated samples (Fig. S7). From the survey spectra (not shown), only carbon, oxygen, 12 ACS Paragon Plus Environment

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nitrogen and/or zinc were detected in the different samples. The N/C, O/C and Zn/C atomic ratios determined from such spectra are given in Table 2, together with the results of peak-fitting the high resolution O 1s and N 1s spectra into different components. The amount of nitrogen present in the samples decreased with activation temperature up to 800 ºC. For instance, the N/C ratio was 0.097 at 600 ºC but was 0.015-0.019 between 800 and 1000 ºC. A similar trend was observed in the case of oxygen: although the O/C ratio decreased with increasing activation temperature up to 800 ºC (i.e., from 0.178 at 600 ºC down to 0.046 at 800 ºC), it tended to level off (or increase slightly) at higher temperatures. Regarding zinc, although it was present in a non-negligible amount at 600 ºC, it was barely detectable at higher temperatures. It is also worth noting that even though a similar level of nitrogen was present in the ZIF8/GO and ZIF-8-only materials activated at the same temperature (800 ºC), the same was not true in the case of oxygen, with the latter carbon material being more oxidized (O/C ratio of 0.130). Such a discrepancy could arise, at least in part, from the different modes of activation of the ZIF-8 particles depending on the presence/absence of sheathing GO nanosheets that were discussed above. Because the process of activation with KOH leads to some oxidation of the carbon material,53 we can expect that a significant fraction of the surface of ZIF-8 particles covered with GO nanosheets will be left unreacted, and therefore unoxidized, during activation. By contrast, all the surface of the “naked”, uncovered ZIF-8 particles will react with, and get oxidized by, KOH (or its reaction products and by-products). As a result, the latter should present a higher overall amount of surface oxygen. Analysis of the peak-fitted O 1s and N 1s spectra (Fig. 7 and Fig. S6 of the Supporting Information) provided some insight into the specific types of heteroatom configurations (Table 2). The O 1s band was fitted to two main components, located at about 531.5 eV (oxygen single-bonded to carbon, as in ethers and hydroxyls) and 533 eV (oxygen double-bonded to carbon as in carbonyls and quinones), together with minor components at ~534.5 eV (oxygen in acid anhydrides) and ~530 eV (oxygen in zinc oxide).59 The latter two components were only present at an activation temperature of 600 ºC due to the absence of zinc at higher temperatures and the relatively low decomposition temperature of acid anhydrides.60 The relative weight of the component at 533 eV increased with activation temperature up to 900 ºC to become clearly dominant, and then decreased at 1000 ºC. Concerning the N 1s band, it was fitted to four components, centered at about 398.5 eV (pyridinic nitrogen), 400 eV (pyrrolic 13 ACS Paragon Plus Environment


nitrogen), 401.5 eV (graphitic nitrogen) and 402.5 eV (oxidized nitrogen groups, as in pyridine oxide).61 At the lowest activation temperature, pyridinic and pyrrolic groups were by far the most abundant. However, consistent with its higher thermodynamic stability,62 the fraction of graphitic nitrogen increased remarkably at higher temperatures, becoming dominant at 900 ºC.

2.3. Electrochemical characterization of the ZIF-8/graphene oxide-derived porous carbons Finally, with a view to their prospective use in energy storage applications (e.g., supercapacitor devices), we carried out a detailed study of the faradaic and capacitive phenomena occurring on the activated ZIF-8/GO carbon materials. First, the operational voltage window and the electrochemical storage phenomena occurring on the prepared materials were studied by cyclic voltammetry (CV, see Fig. 8). We investigated the ZIF-8/GO hybrids activated at 700, 800 and 900 ºC using a KOH/precursor weight ratio of 1, as well as GO alone (no ZIF-8) and ZIF-8 alone (no GO) activated at 800 ºC. The activated GO material [sample GO(1|800)] and the activated ZIF-8/GO hybrids [samples ZIF-8/GO(1|700), ZIF-8/GO(1|800) and ZIF-8/GO(1|900)] exhibited a slightly broader voltage window than that of the activated ZIF-8 material [sample ZIF-8(1|800)], namely, 0.9 vs 0.8 V (Fig. 8). While the ZIF-8/GO hybrid activated at 800 ºC boasted a rather large capacitance (246 F g-1, see Table 3), the one activated at 900 ºC performed quite poorly (81 F g-1), the remaining samples having capacitances between 150 and 200 F g-1. The activated GO displayed three resolved anodic humps centered at -0.1, 0.1 and 0.34 V. With the exception of the latter, these peaks were not well resolved in the cathodic sweep (Fig. 8a). The activated ZIF-8 electrode presented one noticeable faradaic process at ~0.25 V in both the anodic and cathodic scans. Concerning the activated ZIF-8/GO hybrids (Fig. 8b), they all displayed essentially square-shape symmetrical curves regardless of activation temperature, indicating a good capacitive behavior, but with a clear pseudocapacitive charge storage contribution. Samples ZIF8/GO(1|700) displayed wide humps that revealed the presence of a variety of redoxactive groups on its surface. Nevertheless, the maximum contribution was provided by the phenomenon occurring at 0.27 V. In contrast, sample ZIF-8/GO(1|800) showed a square voltammogram with an intense charge transfer around 0.31 V. It can therefore be assumed that this porous carbon has a larger capacitive storage than that of ZIF8/GO(1|700) and more defined redox activity towards the quinone/hydroxyquinone 14 ACS Paragon Plus Environment

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redox pair.63,64 On the other hand, sample ZIF-8/GO(1|900) presented a notable decrease in both total capacitance and faradaic activity. The electron conduction (conductivity) and the molecular transport (electrolyte diffusion) of the porous carbon electrodes were probed by electrochemical impedance spectroscopy (EIS). The Nyquist plots for all the samples discussed in Fig. 8 exhibited three well-defined segments (Fig. 9): (i) a semicircle at high frequencies related to the electronic conduction of the electrode, (ii) a 45º linear part at medium frequencies (Warburg region), which can be related to molecular diffusion, and (iii) a 90º linear part at low frequencies that indicates the formation of the double layer on the electrode surface. The semicircle extending from the high to the medium frequencies region is associated to charge transfer phenomena taking place at the electrode-electrolyte interface. As observed in the high frequency region (inset to Fig. 9), all the electrodes have a similar cell resistance but different resistive behavior caused by the internal structure of the electrode. Samples GO(1|800), ZIF-8/GO(1|700) and ZIF-8/GO(1|800) possessed equivalent series resistance (ESR) values (~0.28-0.31 Ω; see Table 3) much lower than those of ZIF-8(1|800) (0.58 Ω). This result could be due to the positive effect of the (reduced) GO sheets with a relatively well developed graphitic structure present in these materials. Conversely, the electronic conduction of the ZIF-8/GO(1|900) sample was largely hindered (ESR value of 1.67 Ω). The origin of such a decrease in electrode conductivity is currently not understood, but it might be related to some overetching of this carbon material during activation that could degrade the electrical contact between the different parts of the sample. Regarding the medium frequency region of the Nyquist plots, we conclude that electrolyte diffusion was best facilitated in samples ZIF-8/GO(1|700) and ZIF-8/GO(1|800). The hierarchical micro-/mesoporous structure of these carbons should favor the diffusion of the electrolyte molecules into the electrode. Nevertheless, ZIF-8/GO(1|700) exhibited a significant difficulty in switching from resistive to capacitive behavior at low frequencies, which was probably caused by its lower mesopore volume (see Table 1). Square wave voltammetry (SWV) was also carried out for a better understanding of the faradaic phenomena occurring in the electrodes. These were anodically scanned and the current response was analyzed from -0.2 to 0.6 V for the forward scan (Fig. 10a). An attempt to assign the different currents in the SWV curves to specific faradaic phenomena known from the literature is proposed in Table S1 (Supporting Information). The GO(1|800) curve displayed a range of overlapped signals that 15 ACS Paragon Plus Environment


revealed a variety of redox-active functional groups. The signals at lower potentials are believed to arise from the reversible oxidation of aldehyde moieties (PP1, Table S1),65,66 since carboxylic acid groups are not active in strong acidic media. The peak observed at ~0.15 V was ascribed to pyrenes (PP4, Table S1),67-69 whereas the one at ~0.25 V was associated to the activity of the hydroxyquinone/quinone redox pair (PP5, Table S1). 64,70

In the case of ZIF-8(1|800), a broad SWV curve also indicated the contribution of

different species. The maximum intensity for the current response was registered at ~0.25 V, the potential at which quinone groups are active (PP5). Compared with GO(1|800), the ZIF-8(1|800) electrode exhibited a larger current response for potentials between -0.1 and 0.2 V. This low potential activity is thought to arise from the presence of both oxygen functionalities with a low oxidation state (PP1) and nitrogen-containing functional groups. Regarding the hybrid materials, sample ZIF-8/GO(1|700) yielded a similar SWV profile to that of ZIF-8(1|800), but with larger contributions from faradaic reactions occurring at low (-0.1 to 0.05 V) and high (0.2 to 0.5 V) oxidation potentials. Both phenomena are believed to arise from the irreversible oxidation products of pyrrolic-type structures. Thus, the low potential oxidation (PP2) is attributed to nucleophilic attack of water molecules to pyrrolic structures in acidic media, causing the hydroxylation of β-positions.71,72 The oxidation of pyrrolic-type moieties, deriving from imidazole units in the ZIF-8 precursor, can go through further oxidation to quinone-type forms at higher potentials (PP6),71,72 giving rise to the current peak in the 0.2-0.5 V range. On the other hand, the current signal observed from 0 to 0.2 V for both ZIF8(1|800) and ZIF-8/GO(1|700) samples can be ascribed to the reversible oxidation of pyridine (PP3)72-74 or pyrene-like functionalities (PP4). The largest faradaic activity was observed for ZIF-8/GO(1|800), which showed an intense current peak at 0.3 V (PP5). Compared to ZIF-8/GO(1|700), this electrode had a lower activity arising from the oxidation of pyrrolic groups (PP2). Finally, the ZIF-8/GO (1|900) electrode experienced a large decrease in current response all over the studied potential range. Indeed its faradaic activity almost vanished at low potentials, while some current response arising from hydroxyquinone oxidation was detected in the 0.2-0.3 V range (PP5). The observed differences in electrochemical behavior between ZIF-8(1|800), ZIF8/GO(1|700), ZIF-8/GO(1|800) and ZIF-8/GO(1|900) can be explained by their textural (Fig. 5, Table 1) and surface chemistry characteristics (Fig. 7, Table 2). Samples ZIF8(1|800) and ZIF-8/GO(1|700) possess the largest amount of pyrrolic functional groups on their surface, which accounts for their high faradaic response arising from oxidation 16 ACS Paragon Plus Environment

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of such a species at low and high voltages (PP2 and PP6, Fig. 10). This oxidation peak is not so noticeable in ZIF-8(1|800), because the latter has a relatively low N/C ratio and the diffusion of electrolyte molecules could be hindered by its narrower mesoporosity. On the other hand, sample ZIF-8(1|800) is significantly oxidized (O/C atomic ratio of 0.130) and a large proportion of its oxygen functional groups are of carbonyl type (C=O). This result is in agreement with the intense faradaic activity of this material observed at higher potentials (PP5). As a consequence of its higher activation temperature, ZIF-8/GO(1|800) not only possesses a lower overall amount of nitrogen compared to ZIF-8/GO(1|700), but also its relative fraction of pyrrole-type nitrogen is much lower, thus explaining the decreased current in the -0.1 to 0.05 V range (PP2). By contrast, a surface rich in pyridinic and C=O functional groups provides this electrode with a large faradaic activity in the 0-0.4 V range. On the other hand, the activity registered at potentials higher than 0.4 V for ZIF-8/GO(1|700) and ZIF-8/GO(1|800) is believed to arise from the reversible oxidation of dihydroxylated forms derived from oxidized pyrrolic groups (PP7).71,72 It is worth noting that the activity of stabilized quinones (with OH groups in β positions) was also reported in previous works64,67 at similar potential values and, therefore, they may eventually contribute to the current response within the higher potential region (PP8, Table S1). Concerning the ZIF8/GO(1|900) sample, its low pseudocapacitive storage can be ascribed to the transformation or pyridinic and pyrrolic functional groups into quaternary nitrogen and nitrogen oxides (see Table 2) at such a high activation temperature. The latter groups are more related to electronic conduction than they are to redox activity. Rather, the faradaic activity of this electrode can be assigned to a large proportion of C=O functional groups (PP5). Fully reversible redox reactions are highly desired in electrochemical energy storage in terms of suitable performance and long-term stability. The simultaneous study of the front and reverse pulse of the anodic sweep in SWV should afford some insight into this question (Fig. 11). Analysis of the curves revealed that faradaic phenomena taking place in GO(1|800) (Fig. 11a), ZIF-8/GO(1|800) (Fig. 11d) and ZIF8/GO(1|900) (curves not shown) were fast and highly reversible, as they displayed almost symmetric response during the anodic (front) and cathodic (reverse) potential pulses. On the contrary, electrodes deriving from ZIF-8(1|800) (Fig. 11b) and ZIF8/GO(1|700) (Fig. 11c) displayed positive currents during the cathodic pulse, indicating the occurrence of oxidation reactions with a very low rate of electron transfer. The 17 ACS Paragon Plus Environment


inversion of the pulse using ZIF-8(1|800) as the electrode led to two different current peaks (Fig. 11b): one centered at 0.23 V and the other one at potentials below 0.05 V. These signals were coincident with the non-reversible oxidation of pyrrole-like functional groups proposed in Table S1 (PP2 and PP6), which supports the hypothesis proposed above. To a lesser extent, PP2 and PP6 reactions could be therefore responsible for the low reversible charge storage in the ZIF-8/GO(1|700) electrode (Fig. 11c). The faster kinetics observed for ZIF-8/GO(1|800) as compared with ZIF-8(1|800) probably resulted from a better electrolyte diffusion (well-developed mesoporous system) and larger conductivity (presence of graphene sheets). Overall, these results indicate the ZIF-8/GO(1|800) hybrid material to be a good candidate for electrochemical energy storage applications, since it exhibited a better performance than that of any of its two components alone (i.e., either GO or ZIF-8 activated at 800 ºC).

3. Conclusions We have demonstrated a novel strategy, which we refer to as “nanopore lithography”, for the generation of very uniform mesoporosity (~3-4 nm in width) in porous carbons derived from ZIF-8, giving rise to hierarchically micro/mesoporous materials. Such a strategy relied on the use of GO nanosheets enclosing individual ZIF8 particles as masks or templates for the transfer of mesoporous texture to the MOFderived carbonaceous material on the basis of a mild carbonization and chemical activation process. During this process, nanometer-sized holes were created on the GO envelope, which then acted as selective entry points for the activating agent (KOH) to reach and locally etch the enclosed (carbonized) ZIF-8 particles, thus resulting in pores with sizes mainly determined by those of the holes in the GO nanosheets. The success of this mesopore templating approach was seen to be critically dependent on the activation conditions, and particularly on the use of a relatively low amount of activating agent. Larger amounts of KOH, including those typically used for chemical activation of ZIF-8, led to a collapse of the porous texture in the MOF-derived carbon, which was attributed to localized overetching of its structure. A detailed electrochemical characterization of the porous carbons indicated that the generation of uniform and relatively narrow mesoporosity by the approach reported here had a positive impact on their capacitance and charge/discharge kinetics, thus pointing to their potential utility in electrochemical energy storage applications. Finally, provided that a proper encasing by GO nanosheets can be attained, it is expected that the method proposed in this work will 18 ACS Paragon Plus Environment

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be applicable for the templating of mesoporosity not only in carbons derived from other MOFs but also from a number of other organic precursors.

4. Experimental section 4.1. Preparation of materials Unless otherwise stated, all the chemicals and reagents used for the preparation of the different materials were acquired from Sigma-Aldrich. The activated ZIF-8/GO hybrids were obtained through a procedure that involved three steps: (1) growth of ZIF8 particles sheathed by GO nanosheets; (2) carbonization of the resulting ZIF-8/GO hybrids, and (3) chemical activation of carbonized hybrids with KOH as the activating agent. First, GO nanosheets colloidally dispersed in water were obtained through ultrasound-assisted exfoliation of graphite oxide prepared via the Hummers method as reported elsewhere.75,76 The concentration of GO nanosheets in the aqueous dispersions was determined by UV–vis absorption spectroscopy following a previously reported procedure,76 and was typically in the range between 5 and 10 mg mL-1. To prepare GOsheathed ZIF-8 particles, different synthesis conditions were tested, which are discussed in the Supporting Information. However, a typical preparation procedure was the following: zinc acetate (1.86 g) was dissolved in 25 mL of an aqueous solution containing GO sheets (5 mg mL-1) and PVP (37.5 mg mL-1), and then the mixture was bath-sonicated for 1 h to ensure a homogeneous dispersion of the different components. Subsequently, this aqueous solution was poured into a solution of Hmim (6.25 g) in 225 mL of methanol (VWR Chemicals) that was pre-heated to 40 ºC. After being allowed to react at 40 ºC for 2 h under moderate stirring in the mixed water/methanol solvent, the resulting ZIF-8/GO hybrid product was recovered by centrifugation, washed by redispersion in methanol with stirring for 16 h, filtered and finally dried under vacuum at 275 ºC for 6 h to remove the Hmim molecules occluding the pores of the ZIF-8 particles. For comparison purposes, stand-alone ZIF-8 particles were also prepared following the same procedure but excluding GO from the reaction mixture. For the preparation of activated ZIF-8/GO hybrids, the as-prepared and purified ZIF-8/GO materials were first carbonized at 700 ºC for 30 min in a horizontal tubular furnace under nitrogen flow (99.999% purity, 300 mL min-1). Then, the carbonized hybrids were chemically activated with KOH. To this end, the carbonized material was physically mixed with KOH powder (Panreac) at different weight ratios (KOH/carbon weight ratios of 1, 2, 3 and 4) and subsequently heat-treated at a target temperature 19 ACS Paragon Plus Environment


between 600 and 1000 ºC for 1 h under nitrogen flow. The resulting activated materials were extensively washed with water in a Soxhlet extractor until the ionic conductivity of the washing water was below 3 µS cm-1. Finally, the samples were dried at 100 ºC for 24 h under vacuum. The stand-alone, GO-free ZIF-8 particles were also activated using the same protocol. Furthermore, activated GO nanosheets (without ZIF-8) were also obtained. In this case, an aqueous GO dispersion (5 mg mL-1) was first hydrothermally treated at 180 ºC for 2 h to obtain a three-dimensionally assembled, hydrogel-like product that was dried under ambient conditions, carbonized at 700 ºC and then chemically activated with KOH using the same protocols as those followed for the ZIF8/GO hybrids.

4.2. Characterization techniques The samples were characterized by UV–vis absorption spectroscopy, field emission scanning electron microscopy (FE-SEM), transmission electron microscopy (TEM), Raman spectroscopy, X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), N2 physisorption as well as electrochemical techniques, namely cyclic voltammetry (CV), electrochemical impedance spectroscopy (EIS) and square wave voltammetry (SWV). UV–vis absorption spectra (for the starting GO dispersions) were recorded with a Heλios α spectrophotometer (Thermo Spectronic). FE-SEM images were obtained on a Quanta FEG 650 microscope (FEI Company) operated at 30 kV, whereas for TEM a JEOL 2000 EX-II apparatus operated at 160 kV was employed. Specimens for TEM were prepared by dispersing the sample in ethanol via a short sonication treatment and then drop-casting ~40 µL of the resulting dispersion onto a copper grid (200 square mesh) covered with a lacey carbon film. Raman spectroscopy was accomplished with a LabRam instrument (Horiba Jobin–Yvon) using a laser excitation wavelength of 532 nm and using a low incident laser power (0.5 mW) to avoid damaging the samples. XRD was carried out in a D5000 diffractomer (Siemens) at angles (2θ) between 5 and 80º. XPS was performed on a SPECS apparatus working at a pressure of 10-7 Pa with a non-monochromatic Mg Kα X-ray source (1253.6 eV) operated at 11.81 kV and 100 W. N2 adsorption/desorption isotherms were recorded at 196 ºC on a volumetric apparatus (Autosorb 1, from Quantachrome). To this end, the activated carbon samples were previously outgassed at 150 ºC overnight. For the (noncarbonized) ZIF-8/GO and stand-alone ZIF-8 materials, outgassing was carried out at


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250 ºC overnight to ensure the complete removal of non-reacted Hmim that could be occluding the pore structure of the ZIF-8 particles. For the electrochemical characterization, using absolute ethanol as the solvent, electrode pastes were prepared by mixing 10 wt% PTFE as a binder (60 wt% dispersion in H2O), 10 wt% carbon black as a conductive additive (Degussa, Printex 200) and 80 wt% of the carbon materials previously prepared as the active component. Circular electrodes were obtained by pressing 10 mg of the paste in a mold (1 t for 15 s). The resulting circular pellets (0.78 cm2) were sandwiched in a stainless steel mesh (w/w%: 67.77% Fe, 0.08% C; 0.04% P; 1.95% Mn; 0.50% Si; 0.03% S; 10.30% Ni; 17.65% Cr; 1.68% Mo) by pressing (5 t for 15 s). The counter electrode was prepared in the same manner but using 25 mg of paste prepared with Norit DLC Supra 50 carbon as the active component. All the prepared electrodes were vacuum-degassed for 4 h at 120 ºC, soaked into the electrolyte and kept immersed in the solution for one day. The electrochemical behavior of the prepared carbon electrodes was examined in a beakertype cell in 1 M H2SO4 with a Ag/AgCl (1 M KCl) reference electrode by means of an Autolab (PGSTAT302N) potentiostat controlled by Nova 1.10 software. Both the working and counter electrodes were arranged in a stack, separated by filter paper and connected to two Pt wires as current collectors. Nitrogen gas was bubbled through the electrolyte for 20 min prior to measurement in order to remove any oxygen present in the electrolyte solution. CV measurements were accomplished at 2 mV s-1 in a voltage window from -0.3 to 0.6 V. The specific gravimetric capacitance (CS, F/g) of a single electrode was calculated by the formula:

=

(1)

, where r is the scan rate (V s-1), m is the mass of active material in the working electrode (g), Vw is the voltage window (V) and I is the current (A). SWV measurements were carried out using a step potential of 1 mV at a fixed amplitude and frequency to 10 mV and 1 Hz, respectively.

Acknowledgements Financial support from both the Spanish Ministerio de Economía y Competitividad (MINECO) and the European Regional Development Fund (ERDF) through project 21 ACS Paragon Plus Environment


MAT2015-69844-R is gratefully acknowledged. Partial funding by both Plan de Ciencia, Tecnología e Innovación 2013-2017 del Principado de Asturias and the ERDF through project GRUPIN14-056 is also acknowledged. This work was also partially funded by Project “AIProcMat@N2020–Advanced Industrial Processes and Materials for a Sustainable Northern Region of Portugal 2020” (Grant NORTE-01-0145-FEDER000006), supported by Norte Portugal Regional Operational Program (NORTE 2020) under the Portugal 2020 Partnership Agreement through the ERDF, and Project POCI01-0145-FEDER-006984–Associate Laboratory LSRE-LCM, funded by the ERDF through COMPETE2020–Programa Operacional Competitividade e Internacionalização (POCI), and by national funds through FCT–Fundação para a Ciência e a Tecnologia.

Supporting information. AFM of the starting GO nanosheets; Activation of GO nanosheets at different temperatures; Selection of synthesis conditions for the preparation of the ZIF-8/GO hybrids; Thermogravimetric analysis of ZIF-8; Additional XPS of the activated materials; Proposed faradaic processes taking place in the porous carbons.

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Table 1. Textural parameters of the ZIF-8/GO activated at KOH/precursor weight ratios of 1, 2, 3 and 4, and at temperatures of 600, 700, 800, 900 and 1000 ºC [e.g., sample ZIF-8/GO (1|600) denotes a sample activated with a KOH/precursor weight ratio of 1 at 600 ºC. PSD – QSDFT(c) (cm3/g) SAMPLE

SBET(a)

VT (b)

(m /g)

(cc/g)

2

Vµp

Vµp

Vmeso

(dp

research 1..16 - ACS Publications - American Chemical Society

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