Structure‐guided Capacitance Relationships in Oxidized Graphene Porous Materials Based Supercapacitors

Supercapacitors formed from porous carbon and graphene‐oxide (GO) materials are usually dominated by either electric double‐layer capacitance, pseudo‐capacitance, or both. Due to these combined features, reduced GO materials have been shown to offer superior capacitance over typical nanoporous carbon materials; however, there is a significant variation in reported values, ranging between 25 and 350 F g−1. This undermines the structure (e.g., oxygen functionality and/or surface area)‐performance relationships for optimization of cost and scalable factors. This work demonstrates important structure‐controlled charge storage relationships. For this, a series of exfoliated graphene (EG) derivatives are produced via thermal‐shock exfoliation of GO precursors and following controlled graphitization of EG (GEG) generates materials with varied amounts of porosity, redox‐active oxygen groups and graphitic components. Experimental results show significantly varied capacitance values between 30 and 250 F g−1 at 1.0 A g−1 in GEG structures; this suggests that for a given specific surface area the redox‐active and hydrophilic oxygen content can boost the capacitance to 250–300% higher compared to typical mesoporous carbon materials. GEGs with identical oxygen functionality show a surface area governed capacitance. This allows to establish direct structure‐performance relationships between 1) redox‐active oxygen functional concentration and capacitance and 2) surface area and capacitance.

DOI: 10.1002/eem2.12637 Supercapacitors formed from porous carbon and graphene-oxide (GO) materials are usually dominated by either electric double-layer capacitance, pseudo-capacitance, or both. Due to these combined features, reduced GO materials have been shown to offer superior capacitance over typical nanoporous carbon materials; however, there is a significant variation in reported values, ranging between 25 and 350 F g À1 . This undermines the structure (e.g., oxygen functionality and/or surface area)-performance relationships for optimization of cost and scalable factors. This work demonstrates important structure-controlled charge storage relationships. For this, a series of exfoliated graphene (EG) derivatives are produced via thermal-shock exfoliation of GO precursors and following controlled graphitization of EG (GEG) generates materials with varied amounts of porosity, redox-active oxygen groups and graphitic components. Experimental results show significantly varied capacitance values between 30 and 250 F g À1 at 1.0 A g À1 in GEG structures; this suggests that for a given specific surface area the redox-active and hydrophilic oxygen content can boost the capacitance to 250-300% higher compared to typical mesoporous carbon materials. GEGs with identical oxygen functionality show a surface area governed capacitance. This allows to establish direct structureperformance relationships between 1) redox-active oxygen functional concentration and capacitance and 2) surface area and capacitance.
conditions (i.e., an identical C:O ratio) also shows significantly varied capacitance values between 50 and 350 F g À1 . For example, hydrothermally-generated RGO materials from GO solutions, without incorporating additive materials or heteroatom dopant precursors, show capacities in the range of 100-350 F g À1 (Table S1, Supporting Information). Similarly, thermal-shock exfoliated graphene networks, directly derived from GO solids (again no additive precursors related to heteroatom dopants being introduced), have also exhibited capacities between 100-300 F g À1 . The capacitance values in the lower range (e.g., 25-250 F g À1 ) but with significant variation are linked to RGO materials processed under other chemical and thermal reductions and their combinations. Such capacitance variation is linked, in part, to pseudo-capacitance contributions from surface oxygen functional groups. This is to say, RGO samples generated under hydrothermal reduction or thermal-shock exfoliation or other chemical reduction routes result in an RGO structure that remains rich in oxygen functional groups with concentrations between 2 and 12 at% (Figure 1b and Table S1, Supporting Information). Here, it is worth mentioning that for the given low porosities or surface areas of RGO materials, which are usually limited to about 800 m 2 g À1 , such capacities are comparatively higher than the typical EDLC in NPCs (Figure 1a; Figure S1 and Table S2, Supporting Information). Therefore, the capacities in RGO are dominated by the pseudo-capacitance contribution of between 100 and 300% to the EDLC. This can be understood from the capacitance derived from those highly reduced or graphitized RGO samples with low levels of oxygen residues. For Specific capacitance values against specific surface area a) and C:O atomic ratio of RGO materials b) reported in the literature. This data represents more common values to less common; this could be linked to the dominant and low-level pseudo-capacitance contribution of mildly reduced RGO or RGO with lower C:O ratio values and highly reduced RGO or RGO with high C:O ratio values, respectively. The RGO samples with a larger C:O ratio are usually produced by processing under elevated temperatures of more than 500°C. c) PXRD patterns, d) Raman spectra, e) XPS C 1s spectra, and f) XPS survey spectra of EG and GEG samples. The arrow marks indicate changes in peak potion and/or intensity.
Energy Environ. Mater. 2023, 6, e12637 2 of 12 example, RGOs further reduced under elevated temperatures of over 400°C have been shown to yield much lower capacitance values of below 150 F g À1 relative to more than 200 F g À1 for mildly reduced RGOs. [6,13,16,20,22,29,[33][34][35][36][37][38][39][40][41][42][43][44][45][46][47][48][49] In this case, the capacitance variation is linked to reduction (i.e., inherent oxygen functional groups of RGO) levels as well as accessible graphene layer networks of different porosity, conductivity and density. However, in all these analysis scenarios, there is no clear-cut understanding of the overall capacitance variation against key structural parameters of either 1) oxygen functional concentration linked pseudo-capacitance and 2) surface area promoted EDLC. These are important factors for designing materials and applications, but no such relationships are established in the literature for RGO or relevant materials applied in SCs. Therefore, for the first time, this work identifies and demonstrates the relevant structure-guided capacitance relationships for both residual oxygen concentration and specific surface area, by strictly controlling the respective structural parameters in numerous RGOs synthesized to have varying degrees of reduction and porosity.
In this study, different types of GO precursors and their RGO derivatives are synthesized and examined for electrochemical charge storage. RGO structures (named as EGs) are generated via direct solid-state thermal reduction of GO by rapid thermal-shock exfoliation. [18,29,30,38] Following this, EGs are heat-treated at different temperatures; this process not only offers a further reduction in oxygen concentration but also generates graphitized samples (named as GEG, i.e., graphitized EG) with enhancement in electrically conducting sp 2 (with respect to the high concentration of electrically poor sp 3 carbon in GO or mildly reduced RGO materials) carbon components. These GEG samples generated under controlled exfoliation and graphitization, as well as precursor GO synthesis conditions, produce extensively interconnected graphene networks with a controllable enhancement of open porosity; SSA between 300 and 850 m 2 g À1 , PSD between 1.0 and 50.0 nm and specific pore volume between 3.0 and 7.7 cm 3 g À1 . Electrochemical tests yield capacitance values between 30 to 250 F g À1 at 1.0 A g À1 for the EG and GEG samples under alkaline electrolyte. Enhancement in the graphitization leads to considerable improvement in their charge and electron transfer kinetics due to a decrease in the relevant impedance that also contributes to enhancing their rate capacitance. EG samples show the highest capacitance, which decreases systematically with the reduction of oxygen residual functional concentration or enhancement of graphitization in GEG samples. Interestingly, GEG materials with identical reduction or oxygen functional groups show an SSA-governed linear capacitance relationship (with a slope of 0.12 or capacitance ffi 0.12 times of SSA, meaning that an increase of capacitance~60 and~110 F g À1 for an SSA of 500 and 900 m 2 g À1 , respectively). This indicates that most of the discrepancies in the literature-reported capacitance to be linked to relative sp 2 -sp 3 carbon and coordinated oxygen components in the RGO structures. The porosity in the samples is a key defining parameter, but its direct influence can only be seen in highly reduced or graphitized RGO structures; for example, EG structures dominated by significant oxygen functional group concentration show up to 3-times higher capacitance values to their graphitic GEG counterparts (e.g., 200 and 70 F g À1 in EG and GEG, respectively). These oxygen functional group polar sites also enhance the hydrophilicity/wettability or electrolyte ions adsorption in electrode structures thus facilitating efficient charge storage. However, high concentrations of such O-functionalization become detrimental to the rate capacitance and energy storage due to an increased charge transfer resistance in the structures that limits their rate capacitance values during rapid charge-discharge processes. Moreover, by controlling the reduction degree in the EG/GEG samples a linear relationship is also established between the concentration of oxygen groups and capacitance values. This shows a linear decrease of capacitance against a decrease in oxygen content or increase in carbon concentration. These findings clarify the reasons for a significant variation in the literature-reported capacities and offer insights for direct structure-controlled capacities, such as the two key parameters of 1) oxygen functional group and 2) SSAguided capacitance relationships. Here, it is worth noting that the specific gravimetric capacitance and surface area normalized areal capacitance values in GEG materials are on par or superior with other typical mesoporous NPC structures produced under high-temperature pyrolysis routes, indicating that the advantage/potential of graphitized graphenic networks for compact energy storage applications.

Results and Discussion
Graphene-oxide precursors with a controlled degree of oxidation are synthesized via Hummer's method. [16,30] A rapid solid-state thermalshock exfoliation of GO samples is carried out at 300°C in a preheated furnace and as-produced exfoliated graphene materials are named as EG. [18,[28][29][30]38,49,50] GO samples with controllably increased oxidation degree and their thermal-shock derivatives are noted by GO 1 to GO 7 and EG 1 to EG 7 , respectively. Following this, further thermal annealing/graphitization is performed at 1000°C and resultant graphitized EG samples are named as GEG 1 to GEG 7 , in sequence. All the samples are characterized with various complementary techniques before conducting electrochemical tests to evaluate and understand their structurerelated capacitance performance. Figure 1 shows the synthesis and process of GO and RGO or EG and GEG materials and their structurerelated powder X-ray diffraction (PXRD), Raman spectroscopy and Xray photoemission spectroscopy (XPS) characterization results. PXRD patterns (Figure 1c) of EG suggest an exfoliation-induced destruction of GO layered structure ( Figure S2, Supporting Information). A full exfoliation can be seen for EG 6 sample derived from highly oxidized GO 6 , whereas the EG 1 and EG 3 samples derived from mildly oxidized GO 1 and GO 3 indicate some of the closely packed layers as evidenced by weak and broad diffraction peaks that appear at 2h values of around 24 and 43°. As shown in Figure 1d (and Figure S3, Supporting Information), the Raman spectra with a relatively broader G and newly formed D modes and disappearance of 2D mode in EG samples compared to the graphite indicating the highly defective graphene sheets with relative intensity ratios of I D :I G modes to >0.8. Such defective sp 3 carbon to graphitic sp 2 carbon components is revealed by XPS C 1s and O 1s core level spectra (Figure 1e and Figure S4, Supporting Information). Here, it is worth noting that XPS survey spectra (Figure 1f) show the presence of only two atomic components identified by C and O peaks in the EG samples and sp 3 C components are mostly formed by C-O and C=O coordination bonds. The C and O atomic concentrations in EGs are around 86.5-89 and 11-13.5 at%, respectively, indicating a significant reduction degree compared to GO precursors, which show C and O concentrations between 65.5-73 and 27-34.5 at%, respectively (Tables S3 and S4, Supporting Information).
Interesting graphitization features can be seen in GEG samples. PXRD results in Figure 1c show a partial restoration of graphitic nature in GEG 1 sample compared to its weak graphitic counterpart EG 1 . However, such graphitic-related features are eliminated or not formed in GEG samples generated from fully exfoliated EGs. Raman spectra ( sharper D and G modes with significantly increased D mode intensity compared to the modes observed in EGs (e.g., I D :I G ratios further increased to more than 1.1 compared to >0.8 in EGs) suggest for enhancement in both graphitization and defective carbon components. Here, it is worth noting that the Raman spectra without 2D mode formation indicate that GEG samples are in exfoliated state and are supported by PXRD results. Graphitization-induced enhancement in defective carbon concentration is attributed to the reduction in oxygen functional groups during the high-temperature annealing. In this case, the residual oxygen in the structure with C-O and C=O bonding can reduce via dissociating into CO and CO 2 species, thus resulting in carbon defects and some of which thermally rearrange into graphenic structure making sharper D and G modes. The true extent of reduction is identified by XPS survey spectra and C or O atomic concentrations in GEGs are in the range of 97-98 and 2-3 at%, respectively (Figure 1e,f; Figure S4 and Table S5, Supporting Information). Graphitization and oxygen reduction are exclusively revealed by C 1s and O 1s spectra with a shift of C-C peak position (e.g., sp 3 type C-C positioned at 285 eV in EGs to sp 2 type C=C positioned at~284.6 eV in GEGs) and disappearance C=O/COO À in O=C-C (at~287 eV), and -COOH/C-O band in esters and O-C=O (~289 eV) groups related peaks.
The true surface/interface morphology of GEG samples is presented in Figure 2. Scanning electron microscopy (SEM) images show interesting exfoliation-induced porous networking structures. GEG samples derived from GO precursors of increased oxidation degree show a gradually enhanced exfoliation with closely stacked layers to increasingly propped-up and networked layers. Transmission electron microscopy (TEM) images further reveal interfacial interconnected graphenes. Both SEM and TEM images further suggest a significant development of porosity in the samples and related parameters such as SSA, pore volume and PSD are obtained by measuring N 2 adsorption-desorption isotherms in the GEG samples (Figure 2k,l). The qualitative behavior of isotherms indicates highly hierarchical porosity, dominated by mesoporosity and macroporosity, in the samples. SSA and total pore volume (V t ) of the samples are between 275-850 m 2 g À1 and 2.7-7.7 cm 3 g À1 , respectively (Table S6, Supporting Information); both the values increase significantly from GEG 1 to GEG 7 , along with pore population and cumulative pore volume in the mesoporous region ( Figure 2l). Here, it is worth mentioning that the SSA of GEG samples is basically inherited from as-produced EG 1 -EG 7 , which exhibits SSA and V t values of 230-735 m 2 g À1 and 2.3-5.5 cm 3 g À1 , respectively (Table S7, Supporting Information). This was correlated to the oxidation degree of GO 1 -GO 7 precursors. [16,18,29,30,38] More oxygen functional groups in GO structures result in efficient exfoliation and highly porous network structures on the basis of generation of more volatiles of decomposed -C-O/-C=O groups that blast out/exfoliate the GO Figure 2. Surface/interface morphological and porosity characteristics of GEG samples. SEM (top panel) and TEM (bottom panel) micrographs for a, f) GEG 1 , b, g) GEG 3 , c, h) GEG 5 , and d, e, i, j) GEG 6 samples. k,m) N 2 adsorption-desorption isotherms of GEG 1 to GEG 7 and GEG 6@T series. l) Pore-size distribution (on the left axis) and cumulative pore volume (on the right axis) curves of GEG 1 to GEG 7 .
Energy Environ. Mater. 2023, 6, e12637 4 of 12 layers apart and simultaneously create defective/holey/crumpled graphene sheets during the rapid thermal-shock treatment. Structurally, in GO, the oxidized graphene layers are stacked upon interlayer hydroxyl or water molecules. When GO is subjected to sudden heat treatment the volatility or decomposition of edge and interlayer bound carboxyl and hydroxyl/water molecules creates rapid pressure build-up between layers and eventually leads to exfoliation/detonation of the GO structure into graphenic networks of hierarchical porous nature. [18,29,30,38] A nearly linear increase in both SSA and V t of EGs was observed with respect to the amount of C=O/COO group concentration in GO precursors. [30] The decomposition of carboxylic and epoxide groups further creates defect/holey carbon sites/graphene sheets and thus, develops microporosity in EGs. The decomposition of residual -C-O/-C=O groups in EGs and thermal rearrangement of graphene layers under high temperature (1000°C) annealing contribute to further improved SSA and V t values in GEG samples (Figure 1c-f; Tables S4-S7, Supporting Information). Therefore, the results are in accordance with the exfoliation induced by oxidation degree of GO precursors, and are well supported by PXRD, XPS, Raman spectroscopy, and SEM/TEM characteristics. Overall, characterizations reveal that GEG samples of distinctly different structural features, such as varied sp 2 :sp 3 carbon components, networking and porosity, thus also result in different structure density and electrical conductivity.
Additional EG samples are produced by varying thermal-shock exfoliation temperature between 160 and 1000°C in a preheated furnace. From the above understanding, two representative GO precursors, GO 3 and GO 6 , are selected and as-produced EG materials are named as EG 3@T and EG 6@T (where T is a temperature at which GO samples are exfoliated) and following their high-temperature annealing at 1000°C generates GEG 3@T and GEG 6@T samples. Here, it is worth mentioning that the actual exfoliation of GO materials occurs and completes in less than 5 min at 300°C and this exfoliation time is reduced to under a minute when exfoliation is carried out at 1000°C in a preheated furnace. Following this process, annealing is carried out at 1000°C for 6 h after heating at a rate of 5°C per minute from room temperature under air/moisture-free nitrogen atmosphere. The resultant porosity in GEG 3@T and GEG 6@T samples along with C and O content summarily presented in Figure 2m; Figure S5 and Tables S8-S9, Supporting Information, indicates similarly reduced samples to that of GEG 1 -GEG 7 . The SSA of samples increases with exfoliation temperature of between 160 and 300°C and samples exfoliated above this temperature show nearly identical SSA/porosity characteristics. It is important to note that the GO samples treated under slow-heating rate to avoid exfoliation show negligible porosity development in the samples (Table S8, Supporting Information).
Capacitance performance of GEG samples is evaluated using a three-electrode electrochemical test system in alkaline electrolyte (1.0 M KOH) by performing cyclic voltammetry (CV), galvanostatic charge-discharge (GCD) and impedance measurements. The working electrodes are fabricated by sandwiching dry powers of as-produced GEG samples between two-nickel foam current collector strips without the use of a binder, solvents or conducting carbon components. [38] This method results in efficient utilization of the graphenic interfacial structure for both electrolyte and current collector compared to slurry/ doctor-blade technique and/or solvent dispersion coating methods, and the morphology of the sandwiched GEG sample in nickel foam remains identical to GEG powder without any identifiable layered agglomeration/re-stacking ( Figure S6, Supporting Information, and notes for electrode fabrication methods in Supporting Information). Figure 3 shows characteristic CV and GCD curves and their derived rate capacitance curves. The nearly rectangular CV and triangular shape GCD curves represent the typical features of NPCs and suggest a dominant EDLC behavior with negligible contribution from pseudo-capacitance or Faradic capacitance of redox reactions. The area under the CV and GCD curves qualitatively indicates the charge storage capability of the samples. The following formula is used to estimate the capacitance: C = I/2(m 9 (ΔV/Δt)) or C = (I 9 Δt)/(m 9 ΔV), where C is specific gravimetric capacitance (F g À1 ), I is discharge current (A), m is mass of sample (g), Dt is time (s) of discharge, DV is potential (V) of discharge and ΔV/Δt is scan rate. Accordingly, an improvement in capacity from 30 to 110 F g À1 at 1.0 A g À1 for samples GEG 1 to GEG 7 is observed (Figure 3c). Like CV and GCD curves, the rate capacitance curves also show an identical trend against increased current density and suggest an equally graphitized degree in the samples (as demonstrated in Figure 1). Capacitance values at 1.0 A g À1 in the number of (about 20) GEG samples including GEG 1 -GEG 7 and GEG x@T series show a linear dependence on their SSA (Figure 3d). The specific surface area-related capacitance yields a relationship of GCD curves at a current density of 1.0 A g À1 . c) Rate capacitance curves deduced from GCD discharge data, including GEG 1 to GEG 7 and GEG x@T (x = 3 and 6) series. d) Specific capacitance (at 1.0 A g À1 ) variation against BET surface area in GEG samples.

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"Capacitance ffi 0.12 9 SSA", suggesting a superior structure-related capacitance in GEG samples compared to a slope of 0.10 or less in the typical NPCs with high porosity or oxidized carbon nanotube materials ( Figure S1, Supporting Information). [21,27,28,51] Here, it is worth mentioning that all the GEG 1 -GEG 7 and GEG x@T series samples appear to show identical reduction degree ( Figure S7; Tables S5 and S9, Supporting Information). Moreover, as shown in Figure S8, Supporting Information, the as-produced EG 1 to EG 7 samples of near identical reduction degree show specific capacities between 120 and 250 F g À1 at 1.0 A g À1 , and vary monotonically with their SSA values, within the experimental error. The capacities in GEG appear relatively lower than the RGO based structures generated under mild reduction conditions (Figure 1a). For example, the exfoliated graphene or RGO materials produced under plasma and laser/light irradiated conditions in literature show capacitance values between 30 to 300 F g À1 , and a significant drop in capacitance is often observed with further graphitization of the samples under high temperature, anything over 400°C (Table S1, Supporting Information). [2,29,33,38,39,[45][46][47][48][49][51][52][53][54][55][56][57][58][59][60][61][62][63][64][65][66][67] To understand such a significant variation in capacitance between RGO materials of mildly reduced and highly graphitized nature, four sets of RGO samples are further produced with varied reduction/graphitization degrees. In this case, thermal annealing is performed on low, medium and high porosity representative samples of EG 1 , EG 3 , EG 4 and EG 6 at different temperatures of 400, 600, 800, 1000 and 1100°C and resultant graphitized EG samples are named as GEG X-T (T is annealing temperature). As shown in Figure 4 ( Figure S9 and Table S10, Supporting Information), XPS of GEG samples graphitized at increased annealing temperatures exhibit a continuous reduction in oxygen content and improved graphitization. Particularly, the residual C-O and C=O bonded functional groups; C=O/COO À in O=C-C (at~287 and~531.5 eV), C-OH groups (at~532.6 eV) and -COOH/C-O band in esters and O-C=O (~289 and 533.5 eV) related peaks are reduced significantly at 800°C or higher annealing temperature. This shows a rapid loss of C=O functional groups initially with increase of graphitization temperature to 600°C, and above this temperature a significant loss of C-O component is seen. This means that C-O groups are thermally stable relatively to C=O. It is also noted that the active/ defective graphene structures, due to the loss of these functional groups at or above 800°C, show a dominant peak of adsorbed water molecules from atmospheric moisture. This reduction process produces defective carbon sites along with graphitization as is evidenced in Raman spectra (Figure 4c) with sharper and increased relative intensities of D to G modes (with I D :I G ratio increase from 0.8 to 1.2 in EG to GEG or 1.4 in GEG 4-1100 ). Both XPS and Raman results further confirm that considerable graphitization occurs at an annealing temperature of ≥800°C. N 2 adsorption-desorption isotherms and derived PSD and cumulative pore volume curves presented in Figure 4d-f and Figure S10, Supporting Information, reveal a continuously improved porosity; SSA, total pore volume (Table S11, Supporting Information) and PSD in these samples with respect to enhancement in graphitization or, in other words, due to reduction in oxygen concentration or induced defective carbon sites while annealing at increased temperatures. For example, the SSA in GEG X-T sample is increased tõ 100 m 2 g À1 with an increase in graphitization temperature to 1000°C.
The characteristic CV, GCD and rate capacitance curves of GEG X-T samples indicate rapidly decreased capacitance values for the continuously enhanced reduction/graphitization degree in all four sets of samples (Figure 5a-i and Figure S11, Supporting Information). It is interesting to note that the enhancement in porosity contribution is not evidenced here in their charge storage performance. The capacitance in highly reduced GEG samples is dropped to more than 3times lower than their as-produced EG counterparts with~12 AE 1 at % residual oxygen groups (vs.~2.5 AE 1 at% in GEG samples). The as-produced EG and its mildly graphitized GEG samples show quasi-rectangular and quasi-triangular CV and GCD curves, respectivelytypical characters of RGO based or other inorganic based networkssupporting a significant contribution from pseudocapacitance, particularly associated with residual oxygen functional groups (as evidenced in XPS characterization) related redox reactions ( Figure S12, Supporting Information, for relevant reactions). Thus, a reduction in such oxygen content can diminish the additional capacitance contribution to the EDLC. The rate capacitance is improved in the graphitized samples compared to EG counterparts (e.g., capacitance retention of~70% in EG 6 vs 80% in GEG 6 for an increase in current density between 1 and 10 A g À1 ; Figure 5g-i). This can be attributed to enhanced electron transfer reactions and it is evidenced in their electrochemical impedance spectra (Figure 5j-l; Figure S11, Supporting Information), which show a considerable improvement in the charge transfer phenomena by continuously reducing the charge transfer resistance (R CT ) at the electrolyte-electrode interface with enhanced graphitization in the samples. For example, the smaller semicircles of GEG samples relative to EGs in the high-frequency region suggest considerably enhanced ionic dynamics and an increase in straight-line slopes to nearly vertical lines in the low-frequency region supporting the rapid electron transfer reaction and is in line with the dominant EDLC for the improved graphitization degree in the samples. The minimal or no indication of Warburg region between high-and low-frequency regions also suggests rapid ionic dynamics without any diffusion barrier in the samples indicating that efficient ion/electrolyte adsorption/interaction. This can be attributed to the dominant mesoporous nature with abundant active/defective carbon sites of the samples. [1][2][3][4][5][6][7][8][9][10][11][12][13][14][15][16][17][18][19][20] Moreover, the resistance obtained by extrapolating the impedance curves linear part of graphitized samples shows a correlatively decreased resistance values with reduction in oxygen concentration ( Figure S13, Supporting Information).
Despite these improved charge/electron transfer kinetics, a considerable drop in capacitance (e.g., about three times in EG 3 (or EG 4 ) to GEG 3 (or GEG 4 ) or from 240 to 105 F g À1 at 1 A g À1 in EG 6 to GEG 6 ) while an increase in the SSA of graphitized samples of about 100 m 2 g À1 (Figure 6a). A well-agreed relationship can be seen between rapidly decreased capacitance values with increased degree of graphitization of the samples that shows small increase in SSA. A relatively large decrease of capacitance values in low and medium SSA/ porosity EG 1 -GEG 1 and EG 3 -GEG 3 samples compared to high porosity EG 4 -GEG 4 and EG 6 -GEG 6 series can be seen and is attributed to the low-level of active/defective carbon with high degree of reduction in graphitized samples limiting the electrolyte interaction/wettability, for example, in EG 1 -GEG 1 . Interestingly, as shown in Figure 6b, the enhancement in SSA of the graphitized samples is well-correlated with the reduction in oxygen concentration and thus also yields a wellcorrelated relationship between capacitance and concentration of oxygen or carbon in the samples; that is, a direct linear relationship can be established for this data. All four sets of samples show rapidly decreased capacitance values against reduced oxygen concentration (Figure 6c). XPS focusing on C 1s and O 1s features are further analyzed to obtain specific contribution of C-O and C=O functional groups and to understand their influence on the capacitance values. As shown in Figure 6d Energy Environ. Mater. 2023, 6, e12637 6 of 12 and Figure S14, Supporting Information, the increased C-C (a graphitic component that excluded C-O and C=O contributions) concentration shows linearly decreased capacitance. An inverse trend, or a rapid increase in capacitance, with respect to the amount of C-O and/or C=O components of C 1s XPS spectra is observed (Figures 4a and 6e,f and Figures S9 and S14, Supporting Information). Likewise, analysis of O 1s XPS spectra, which contain mainly C-O and C=O components further suggest specific C-O and/or C=O concentration dependent capacitance variation. In this case, C=O concentration mostly governs the capacitance in mildly/low-temperature reduced samples with low C:O ratios, whereas the high temperature-reduced samples capacitance largely depends on C-O concentration (Figure 6g,h and Figures S9 and S14, Supporting Information). Overall, these results suggest a considerable decrease in capacitance in the samples with respect to a small change in both C-O and C=O concentrations; specifically, irrespective of SSA enhancement in the sample, C=O amount in the samples of mildly reduced conditions, which reduces C=O functionality, primarily influence the achievable capacitance, whereas, C-O amount governs the capacitance in the highly reduced samples.
Another factor, the wettability of the structures can affect the electrolyte ionaccessible surface and thus capacitance. [16,28,68,69] In carbon/graphitic materials, the hydrophilicity or wettability is mainly achieved by creating surface Ofunctional polar groups and/or enhancing the surface area/porosity with defective carbon sites and slit-, narrow-, micro-and mesoporosity (via capillary action and transport channels and reservoirs). With these features, better wettability of the electrodes facilitates ions transport and thereby enhances the capacitance. The increased wettability also reduces charge transfer resistance, whereas poor wettability of carbon-based electrodes induces high resistance to the diffusion of ions within the micropores during the charge/discharge process and can limit the capacitance. GEG samples with a dominant fraction of micro-/mesoporosity with high active/defective carbon site concentration exhibit smaller impedance values for both ion diffusion and electron transfer and high rate capacitance compared to the EG counterparts with relatively high hydrophilic oxygen group content. This suggests that GEG samples under graphitized conditions still maintain their wettability or offer high surface adsorption efficiency to the electrolyte. The electrodes being under electric field (a kind of driving force for ion adsorption) can effectively promote electrolyte ion adsorption. Interestingly, under electrochemical conditions, these hydrophilic O-functional groups also take part in redox reactions and contribute to pseudo-capacitance. Therefore, these two factors in electrochemical studies are inseparable or offer similar results of enhancing capacitance. From this, one can conclude that both porosity and O-functional groups are main contributing factors for boosting the overall capacitance of the samples and, for a given SSA, the concentration of polar C-O/C=O groups can boost the capacitance values of up to 3 times and at any given oxygen group concentration the capacitance can increase at a rate of 0.1-0.4 times their SSA. Excessively oxygenated groups in the structures are found to be detrimental for charge storage performance, specifically causing the rapid deterioration in rate capacitance, due to the low electrical conductivity from the increased percentage of sp 3 type graphene carbon components of oxidized carbon species, for example, carboxylic and epoxide groups at the expense of electrically conducting sp 2 type graphene carbon.
As presented in Figure 1b, the high values of C:O ratio in RGO samples resulted in relatively lower capacities compared to the samples with low values of C:O ratio. Accordingly, the current work demonstrates possible reasons for capacitance variation in the graphitized RGO materials and offers underlying insights by presenting controllable design and critical analysis of numerous RGO based materials with varied porosities, along with residual oxygen functionality or more specifically, degree of reduction and graphitization. Interestingly, additional thermal annealing treatment of as-produced EG samples shows a significant reduction in oxygen concentration with increasing temperature from 400 to 1000°C. This process also results in enhancement in defective carbon sites and porosity in the graphitized samples and thus allows understanding of the effect of structure-related charge storage capacity in the RGO materials. The rate performance of SCs are enhanced with respect to the annealing temperature, which improves the C:O ratio and electrically conducting graphitic sp 2 components in the structures and enhances the overall charge/electron transport properties. Finally, as shown in Figure 6i (Table S2, Supporting . d-f) GCD curves at a current density of 1.0 A g À1 . g-i) Rate capacitance curves deduced from GCD discharge data. j-l) Impedance curves.

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Information), the SSA normalized capacitance values of RGO samples exhibit very high areal capacitance values of more than 10 lF cm À2 (or 0.12 F m À2 ) to 100 lF cm À2 compared to the typical NPCs with high SSA and maximum areal capacitance of 11 lF cm À2 . This indicates that graphene materials can be favorably considered for compact energy storage applications over other highly porous carbon structures produced under high-temperature chemical activation/template/pyrolysis conditions.

Conclusions
This work is designed to reveal specific reasons for significantly varied capacitance values in RGO-based structures and to establish a direct structure-controlled capacitance performance relationship. The conclusions are drawn by synthesizing numerous RGO structures with controllably varying the degree of oxidation, reduction and graphitization. The resultant materials not only offered tunable porosities; for example, specific surface area variation between 250 and 850 m 2 g À1 , but also yielded a systematic reduction in surface functional oxygen group concentration between 2-13 at% along with simultaneously increased graphitization and defective carbon concentration. RGO samples processed under controlled reduction conditions revealed a systematic loss in the capacitance with the removal of redox-active oxygen groups and ultimately the highly reduced or graphitized RGO materials show significantly lower capacitance values, about 30% or 3-times lower than RGO counterparts produced under mild reduction routes. Furthermore, this controlled reduction process also establishes a direct relationship of Figure 6. Capacitance performance against structural characteristics of GEG 1-T , GEG 3-T , GEG 4-T and GEG 6-T series samples. a) Specific capacitance (at 1.0 A g À1 ) variation against SSA. b, c) SSA and capacitance variation against total oxygen concentration (of XPS data). d-f) Capacitance (at 1.0 A g À1 ) variation with respect to C-C, C-O and C=O group amount derived from C 1s XPS spectra. g, h) Capacitance (at 1.0 A g À1 ) variation with respect to C-O and C=O group amount derived from O 1s XPS spectra. The dotted arrows are guide lines for the data variation. i) Areal capacitance, that is, BET surface area normalized specific capacitance (at 1.0 A g À1 ), variation against SSA of a wide range of RGO and NPCs with relatively large SSA values.
Energy Environ. Mater. 2023, 6, e12637 9 of 12 a key role of oxygen residuals promoted capacitance values -this is the first report to reveal such dependence and clarifies the large discrepancies in SCs capacitance values reported in the literature. Another noteworthy structure-guided relationship between specific surface area and capacitance variation is established in highly graphitized RGO samples with oxygen residues of less than 2.5 AE 1 at% or nearly dominated by EDLC with the following relationship; capacitance ffi0.12 times of specific surface area (C ffi 0.12 9 SSA) compared to C ffi 0.10 9 SSA in a typical highly porous carbons. Such a relationship has not been revealed in the literature nor established in RGO structures of different reduction degrees due to dominantly varied pseudo-capacitance contributions over surface area specific EDLC. Interestingly, both the specific gravimetric capacitance and specific area normalized areal capacitance values in the RGO or EG and GEG materials of dominant mesoporosity are superior to typical mesoporous carbon materials produced via silica or other molecular template structures under intense pyrolysis processing routes, thus suggesting having a potential role for compact energy conversion and storage applications.
Synthesis of GO samples: GO 1 -GO 7 samples of different degrees of oxidation were prepared by Hummer's and modified Hummer's methods. [16,30] GO 3 sample was synthesized as follows; graphite powder, 10 g was stirred with cold concentrated H 2 SO 4 (230 mL at 0°C). KMnO 4 (30 g) was added to the suspension slowly to prevent a rapid rise in the temperature (less than 20°C). The solution underwent a color change at this point from black to a very dark green. After removal of the ice-bath, the mixture was stirred at room temperature for another 2 h. Deionized water (230 mL) was slowly added to the reaction vessel to keep the temperature under 98°C. The diluted suspension was stirred for an additional 15 min and further diluted with DI water (1.4 L) before adding H 2 O 2 (100 mL). Upon addition, vigorous foaming occurred and the solution turned brown. The mixture was stirred for 2 h at room temperature and left overnight. The product settled at the bottom was separated from the excess liquid by decantation followed by centrifugation. The product was washed by centrifugation until the pH reached neutral. Freeze-dried to obtain a final product. GO 1 , GO 2 and GO 4 were prepared under similar conditions but with a change in graphite:KMnO 4 weight ratio, for example, in 1:1, 1:2 and 1:4 g g À1 , respectively. GO 6 sample was synthesized as follows; graphite powder (2.0 g) was added to 9:1 mixture of concentrated H 2 SO 4 (45.0 mL) and H 3 PO 4 (5.0 mL) under stirring at cold (%0°C). KMnO 4 (12.0 g) was then added slowly to prevent the sudden temperature rise, not more than 10°C. The solution became a very dark green. The reaction mixture was then heated to 50°C and left for 3 h under stirring. At this point, the mixture turned into a brown paste. Deionized water (240.0 mL) was added very slowly, again to prevent a sudden temperature rise of highly exothermic reaction (water to a concentrated acid) to the mixture, which turned to a dark brown color. The reaction mixture was left to stir for several minutes. In the final step, H 2 O 2 (16.0 mL) was slowly added to the solution causing vigorous foaming and a color change to a bright yellow. The solution was stirred for another 30 min at a warm state (50°C) and left to settle overnight at room temperature. The solid product was separated from the excess liquid by centrifugation followed by decantation. The product was then washed in dilute HCl (3.4%, 750 mL) to remove any remaining salts, followed by further washing in DI water until the washings were pH neutral. The GO-6 sample was obtained after being freeze-dried for about 7 days. GO 5 was prepared under similar conditions as GO 6 but with a change in graphite:KMnO 4 weight ratio of 1:5 g g À1 and reaction (graphite + acid + KMnO 4 mixture) heating time set to 6-8 h under stirred conditions at 50°C. GO 7 was prepared under similar conditions as GO 6 but with a change in reaction (graphite + acid + KMnO 4 mixture) heating temperature and heating time; the mixture was stirred for 8 h at 50°C and left overnight at room temperature.
Synthesis of EG 1 to EG 7 samples: Thermal exfoliation has been carried out at 300°C (unless specified) in a preheated vertical tube furnace. 1.0 g of GO sample was charged into a glass tube of 1.5 inch diameter and 12 inch long under ambient air. The tube was then covered with a paper towel in ambient air and placed in a furnace. From this point, the exfoliation can be seen in about 2-5 min, some of the samples have exfoliated in multiple steps. Thus, the sample tube was still left in a furnace for additional 2 min for completing the process. After this, the sample tube was removed from the furnace and left on the bench to cool down. Exfoliated samples were collected for further characterization. No further chemical treatment was performed. EG 3@T , EG 4@T and EG 6@T samples were produced via thermal-shock exfoliation described above using 1.0 g of GO 3 , GO 4 and GO 6 precursors in a single batch and under different preheated temperatures of 160, 200, 250, 300, 400, 600, 800 and 1000°C. At 160 and 200°C exfoliation took more than 10 and up to 30 min, whereas at or above 800°C the process was completed rapidly in about a minute or less. For example, the GO 4 sample exfoliated at 600°C is named as EG 4@600 .
Synthesis of GEG samples: The as-exfoliated EG samples were annealed at 1000°C for 6 h, in a horizontal tube furnace, after heating it from room temperature at a rate of 5°C min À1 , and under nitrogen gas flow. The furnace tube was closed with a gas feed through end seals and sample area was thoroughly purged with nitrogen. The nitrogen flow was maintained throughout the reaction. The graphitized EG 1 to EG 7 , and EG 3@T , EG 4@T and EG 6@T samples are named as GEG 1 to GEG 7 , and GEG 3@T , GEG 4@T and GEG 6@T . GEG 1-T , GEG 3-T , GEG 4-T and GEG 6-T were produced via annealing EG 1 , EG 3 , EG 4 and EG 6 at different temperatures between 400 and 1100 C for 6 h, in a horizontal tube furnace, after heating it from room temperature at a rate of 5°C min À1 , and under nitrogen gas flow, as described above for GEG samples. The graphitized samples obtained at different graphitization temperatures, for example, EG 6 graphitized at 600°C for 6 h is named as GEG  .
Characterization of samples: Powder X-Ray diffraction (PXRD, Cu Ka radiation, Thermo Scientific) was carried out in the scan range of 2h = (2-60)°and step size of 0.01°. Raman spectra were recorded with a 950 microscope using a 514.5 nm laser on a Renishaw inVia spectrometer. X-ray photoemission spectroscopy (XPS, Al-K-alpha, Thermo Scientific) data, scanning electron microscopy (SEM, Jeol) and transmission electron microscopy (TEM, Jeol) measurements were carried out on the samples supported on a carbon tape or a carbon-coated copper TEM grid. N 2 adsorption-desorption isotherms in the pressure range of vacuum to 0.995 of relative pressure were measured at 77 K using a Quantachrome Autosorb-iQC. The specific surface area was determined from the N 2 isotherm, according to the Brunauer-Emmett-Teller (BET) method. QSDFT (quenched solid density functional theory, available within the Quantachrome ASiQwin isotherm analysis software) method with slit/cylindrical pores (≤50 nm) applied to obtain a pore-size distribution and cumulative pore volume. The total pore volume was estimated from the amount of N 2 adsorbed at a relative pressure, P P 0 À1 of %0.995. The samples were initially degassed at 180°C for up to 24 h prior to the actual adsorption isotherm measurements. Fabrication of supercapacitor electrodes: Working electrodes were prepared via direct dry powder packing in a nickel foam as follows; 1.00 mg of EG materials (with 0.01 mg accuracy) powder directly spread uniformly onto a 1 cm 9 1 cm area of precleaned and hand-flattened 1 cm 9 3 cm nickel foam strips and sealed into a pouch-like cell with another piece of nickel foam and compressed at 1.0-1.5 ton of pressure for 1 min using a hydraulic press. Electrochemical tests of electrodes: All electrochemical tests were carried out using an Autolab (Metrohm) electrochemical workstation, by a three-electrode method. Prefabricated EG and GEG based nickel foam strips were used as working electrodes with active samples area of 1 cm 9 1 cm, and Pt (1 cm 9 1 cm) and Ag/AgCl/saturated KCl as counter and reference electrode, respectively, in 1.0 M KOH electrolyte at room temperature. Before actual measurements, the electrode was subjected to several CV cycles at a voltage-sweep rate of 100 mV s À1 until a stable CV was obtained. Then, the actual CV tests were conducted at 50 mV s À1 in a fixed potential range of À0.8 to 0.0 V. The charge-Energy Environ. Mater. 2023, 6, e12637 discharge curves with lower and upper and cut-off potentials of À0.8 and 0.0 V, respectively, were recorded at discrete applied current densities between (0.5 and 10.0) A g À1 . Electrochemical impedance spectroscopy (EIS) tests were performed at open circuit potential with a sinusoidal signal in a frequency range from 100 kHz to 10 mHz at an amplitude of 10 mV. The specific gravimetric capacitance (F g À1 ) was calculated on the basis of charge-discharge curves at different current densities using the following equation: C = (I 9 Δt)/(m 9 ΔV), where C is the galvanostatic specific capacitance (F g À1 ), I is the discharge current (A), m is the mass of the electrode material (g), Dt is the discharge time (s), DV is the operating discharge potential window (V).