Edge-Plane Exposed N-Doped Carbon Nanofibers Toward Fast K-Ion Adsorption/Diffusion Kinetics for K-Ion Capacitors

Abstract

Open AccessCCS ChemistryRESEARCH ARTICLE1 Aug 2020Edge-Plane Exposed N-Doped Carbon Nanofibers Toward Fast K-Ion Adsorption/Diffusion Kinetics for K-Ion Capacitors Zheng Yi, Song Jiang, Yong Qian, Jie Tian, Ning Lin and Yitai Qian Zheng Yi Department of Applied Chemistry, Hefei National Laboratory for Physical Science at Microscale, University of Science and Technology of China, Hefei, Anhui 230026 Google Scholar More articles by this author , Song Jiang Department of Applied Chemistry, Hefei National Laboratory for Physical Science at Microscale, University of Science and Technology of China, Hefei, Anhui 230026 Google Scholar More articles by this author , Yong Qian Department of Applied Chemistry, Hefei National Laboratory for Physical Science at Microscale, University of Science and Technology of China, Hefei, Anhui 230026 Google Scholar More articles by this author , Jie Tian Experimental Center of Engineering and Material Science, University of Science and Technology of China, Hefei, Anhui 230026. Google Scholar More articles by this author , Ning Lin *Corresponding author(s): E-mail Address: [email protected] E-mail Address: [email protected] Department of Applied Chemistry, Hefei National Laboratory for Physical Science at Microscale, University of Science and Technology of China, Hefei, Anhui 230026 Google Scholar More articles by this author and Yitai Qian *Corresponding author(s): E-mail Address: [email protected] E-mail Address: [email protected] Department of Applied Chemistry, Hefei National Laboratory for Physical Science at Microscale, University of Science and Technology of China, Hefei, Anhui 230026 Google Scholar More articles by this author https://doi.org/10.31635/ccschem.020.201900103 SectionsSupplemental MaterialAboutAbstractPDF ToolsAdd to favoritesTrack Citations ShareFacebookTwitterLinked InEmail Sluggish kinetics severely limit the development of potassium-ion hybrid capacitors (PIHCs). Exposing active sites is recognized as an ideal strategy to resolve this issue, but the corresponding material design is challenging. Herein, carbon nanofibers with abundant, exposed edge-plane active sites due to (002) orientation adjustment were developed by a molten salt-assisted procedure. Importantly, due to the radial (002) orientation with more active edge-plane sites to adsorb K and shorten the K diffusion distance, the obtained carbon nanofibers harvest improved K adsorption/diffusion kinetics. Meanwhile, theoretical calculations indicate that the synchronically introduced N-doped defects can also lower the diffusion barrier and enhance K adsorption kinetics. Ex situ characterizations and electrochemical studies prove the improved kinetics that significantly improve the K storage properties of the obtained carbon nanofibers. Hence, a high cycling capacity of 252.8 mAh g−1 at 100 mA g−1 after 500 cycles and rate capacity of 181.5 mAh g−1 at 1000 mA g−1 after 1200 cycles have been achieved. Remarkably, the as-developed PIHCs deliver an energy density of 170 Wh kg−1 over 1–4 V, along with capacity retention of 81.6% at 2000 mA g−1 after 10,000 cycles. Download figure Download PowerPoint Introduction Currently, universal energy storage/conversion devices are lithium-ion batteries (LIBs) because of their high energy density.1–4 However, the power density of commercial LIBs is still deficient, which retards faster development in the power battery field.5,6 Moreover, the cycling stability of commercial LIBs must be enhanced.7 In comparison, supercapacitors (SCs) have recently attracted increasing attention, because of their high power and long lifespan.8,9 However, the energy density is relatively poor. To take into account both energy and power density, new energy storage device architecture, assembled by battery-type anodes and capacitor-type cathodes, has been developed to showcase both high power density and energy density.10–12 Until now, various kinds of battery-SC devices, such as lithium-ion hybrid capacitors (LIHCs), sodium-ion hybrid capacitors (SIHCs), potassium-ion hybrid capacitors (PIHCs), and zinc-ion hybrid capacitors (ZIHCs), have been developed.2,11,13 Among them, PIHCs are attracting significant attention because of the suitable potential and natural abundance of potassium, which have been reported to exhibit high energy density at high power density and excellent cycling life. Therefore, PIHCs are expected as a candidate to replace potassium-ion batteries (PIBs).14–16 To target high energy density PIHCs, anode performance is the key factor affecting practical performance.17 Therefore, meaningful effort has been made to enhance the electrochemical properties of the anodes for PIBs and PIHCs. Until now, many kinds of anodes, such as carbon-based materials, alloy-type anode materials, metal oxides, sulfides, and selenides, have been widely developed for PIB or PIHC anodes.18–28 Based on the high K storage capacity for reversible formation of KCx intercalation compounds and adsorption/desorption of K ions on the surface/interface, carbon-based anodes have long been selected as the ideal choice. For example, a highly graphitic carbon nanocage was reported with excellent cyclability and superior potassiation capacity of 175 mAh g−1 at 35 C.29 N-Doped hierarchical porous hollow carbon spheres were fabricated for application in high-performance PIBs or PIHCs anodes. This carbon-based anode coupled with an activated carbon cathode demonstrated an exceptionally high energy density of 114.2 Wh kg−1, power density of 8203 W kg−1, and long-life cycling capability.17 However, the most serious challenges facing carbon-based anodes are sluggish kinetics, poor rate, and bad cycling performance, due to the large ionic size of K+ (1.38 Å in radius) that induces dramatic volume changes.30–33 For instance, Ji et al.34 showed that a graphite anode could deliver a capacity of 263 mAh g−1 at C/10, but maintained only 80 mAh g−1 at 1 C. The poor rate performance has hindered the fast development of PIBs. Therefore, it is imperative to improve the rate performance and K storage kinetics of carbon-based anode materials. The introduction of doping and defects can increase the K ions active sites and adsorption kinetics. As a representative case, Guo et al.35 reported that Sulfur/Oxygen codoping favors the adsorption of K and reduces the structural deformation of hard carbon microspheres during potassiation/depotassiation, thus endowing the materials with a high capacity of 226.6 mAh g−1 over 100 cycles. Moreover, rationally adjusting and controlling the microstructure of the materials is also effective for kinetics improvement. Generally, the edge-plane is theoretically predicted to be more electroactive than the basal plane.36,37 As shown in Scheme 1, the basic plane carbon nanofiber has poor active sites and a long diffusion distance, resulting in sluggish kinetics and poor rate capability. However, for edge-plane carbon, rich active sites and short diffusion distance can endow the nanofiber with improved kinetics and rate capability. Hence, combining the electroactivities of edge-plane-rich sites and doped defects in carbonaceous materials can be a promising strategy for fast K-ion storage performance. Scheme 1 | Schematic of the K insertion model in basic plane and edge-plane carbon nanofibers. Download figure Download PowerPoint Herein, by employing a molten salt-assisted route, a one-dimensional (1D) N-doped carbon nanofiber was designed with radial (002) orientation perpendicular to the fiber axis, which is expected to expose more active edge-plane sites and enhance K adsorption kinetics. Those exposed edge-plane active sites can facilitate the potassium-ion intercalation/deintercalation and shorten the ion diffusion distance. The N-doped defects were also introduced to provide extra active sites for adsorbing K ions, reduce the diffusion barrier, and enhance the K adsorption kinetics. Moreover, the well constructed 1D framework can enhance the specific surface area and active sites and shorten the K diffusion distance. With the functionally integrated metrics, the as-obtained carbon nanofibers delivered high capacity and good rate capability compared with the counter samples. Moreover, as a battery-type anode for PIHCs, the prepotassiated carbon nanofibers could also show good rate capability over 10,000 cycles with high energy/power density. Experimental Methods Experimental Methods are available in the Supporting Information. Results and Discussion N-Doped carbon nanofibers (denoted as CFs) with radial (002) orientation were prepared by a controllably molten salt-supported route within a sealed stainless-steel autoclave. Figure 1a and Supporting Information Figure S1 illustrate the synthetic process of CFs and carbon spindles (denoted as CSs). ZnCl2 (m.p. 290 °C) was employed as the molten salt reaction medium, which can homogenize the local autogenic heat to create a homogeneous system in the reactor and to disperse the products and maintain a uniform reaction and growth environment.38–40 Ferrocene was used as the carbon source, which can be pyrolyzed at 600 °C to produce the target products. Figure 1 | (a) Schematic representation of the CFs. (b) XRD patterns, (c) Raman spectra of the CFs and CSs. (d) XPS spectra for N 1 s of the CFs. (e) Pore volume distribution of CFs and CSs. Download figure Download PowerPoint Figure 1b comparably presents the XRD patterns of as-prepared CFs and CSs. The broad diffraction peaks of CFs at 25.0° and CSs at 25.9° could be assigned to the (002) lattice planes of the carbon. Based on the Bragg equation, the average interspaces of CFs and CSs can be calculated to be 3.6 and 3.44 Å, respectively. These results agree well with the measurements from high-resolution transmission electron microscopy (HRTEM), as discussed further. The enhanced carbon interspaces can be ascribed to the introduction of N-doping defects. The Raman spectra (Figure 1c) exhibit two peaks located at 1339 and 1596 cm−1, which are assigned to the disorder/defect-induced band (D band) and in-plane vibrational band (G band) of the sp2 carbon, respectively. The relative intensity ratio (ID/IG) is developed to measure the disorder degree and structural defects. The ID/IG of CFs is 0.78, higher than that of CSs (0.73), which can be ascribed to the higher N-doping defect concentration. The integral intensity of ID/IG is also related to the in-plane crystallite sizes (La) of the CFs and CSs, which are calculated to be 24.6 and 26.3 nm [by employing the equation, La = (2.4 * 10−10) λlaser4 (ID/IG)−1], respectively.41 Also, the G band in CFs (1596 cm−1) was shifted to 1593 cm−1 (CSs), which may also be attributed to the enhanced carrier concentration induced by N-doping.42 X-Ray photoelectron spectroscopy (XPS) was performed to exhibit the doping states of the samples. The survey spectra of CFs ( Supporting Information Figure S2A) include C, N, and O elements. The N content is about 1 wt %. In contrast, there is no N content detected in the CS sample ( Supporting Information Figure S3). The high-resolution C1s spectra ( Supporting Information Figure S2B) of the CFs mainly exhibit three peaks, including C–C, C–N, and C=O bonding, respectively. It can be observed that C–C bonding is much higher than C–N or C=O. Figure 1d is the high-resolution N 1s spectrum of the CFs, which can be deconvoluted as pyrindinic-N (N-6), pyrrolic-N (N-5), and graphitic-N (N-Q), respectively.43,44 The relative ratios of N-6, N-5, and N-Q were calculated to be 28.2%, 26.8%, and 45.0%, respectively. Previous studies have proved that N-6 and N-5 are highly active sites for attracting and storing K ions.42,45 The N2 adsorption curves ( Supporting Information Figure S4) show that the Brunauer–Emmett–Teller (BET) specific surface areas of the synthesized CFs and CSs are 328 and 217 m2 g−1, respectively. Compared with the CSs, the enhanced BET specific surface areas of CFs can be ascribed to the decreased size and increased defects due to the introduction of N-doped sites. The pore size distributions of both CFs and CSs range from 1.6 to 150 nm (Figure 1e). The total pore volumes of the CFs and CSs are 0.39 and 0.16 m3 g−1. In general, the enhanced pore volumes and BET specific surface areas may increase the K-ion storage sites.46 Figure 2a and Supporting Information Figure S5 are the low-resolution SEM and TEM images of the CSs, respectively. It can be seen that direct pyrolysis of ferrocene in molten ZnCl2 gives a spindle-like carbon submicrostructure mixed with rods and a relatively smooth surface. The average sizes of the CSs are about 1 μm ( Supporting Information Figure S6A). From the TEM image (Figure 2b), it can be seen that the obtained CSs are solid and without any pores. The HRTEM image of the CSs is presented in Figure 2c. It can be seen that the lattice fringe is about 0.34 nm, which is radially oriented perpendicular to surface. From the EDS mapping ( Supporting Information Figure S7), the carbon element is uniformly distributed. Figure 2 | (a) SEM image, (b) TEM image, and (c) HRTEM image of the CSs. (d) SEM image, (e) HAADF-STEM image, and C, N element mapping, (f) TEM images, (g) HRTEM image, and (h) typical SAED pattern of the CFs. (i) TEM images, (j) HRTEM image, and (k) SAED pattern of the commercial CNTs. Download figure Download PowerPoint Notably, with the addition of NH4H2PO4, the obtained sample evolved to be one dimensionally rope-like, as shown in Figure 2d. The average size of the CFs is about 120 nm in diameter ( Supporting Information Figure S6B). Due to the formation of nanofiber from mixed spindle and rods (CSs), the yield was notably improved (∼50%). From the EDX mapping pictures (Figure 2e), both the carbon and nitrogen elements were uniformly distributed on all CFs. The existence of N suggests that some of the C elements were displaced by N elements, and thus, the formation of N-doped defects in CFs, which agrees well with the XPS results (Figure 1d). The low-resolution TEM image (Figure 2f) further shows that the obtained CFs are nanofiber-like. Moreover, the nanofiber is intertwined with many small fibers to form the rope-like morphology with extra space and pores ( Supporting Information Figure S8). It is reasonable to speculate that the NH4H2PO4 (m.p. 180 °C) additive plays an important role in the formation of N-doped carbon nanofibers. In the fabrication process, the NH4H2PO4 is resolved to produce NH3, H2O, and phosphorus-based compounds (NH4H2PO4→NH4PO3 + H2O; NH4H2PO4 →H3PO4+ NH3). As carbon sources, the ferrocene was pyrolyzed according to the following equation: C10H10Fe + H2O → 10 C+ FeO + 6H2.47 The evolved FeO acts as the catalyst to promote pyrolysis of ferrocene and catalyze the formation of carbon nanofibers. The derived NH3 acts as a N source to generate the N-doped defects in CFs. To present the orientation of the obtained products, HRTEM images and typical selected area electron diffraction (SAED) patterns are provided. As shown in Figure 2g, the lattice fringe of the CFs is about 0.36 nm, which is radially oriented perpendicular to the 1D direction of the nanofibers. Compared with the CSs (Figure 2c), there are some mesopores in the CFs (cycled in white), which endow the CFs with enhanced BET specific surface area ( Supporting Information Figure S4) and increased active sites. Moreover, the SAED patterns further confirm the orientation information of the CFs.48 As presented in Figure 2h, the diffraction ring of the (002) plane is evolved to two symmetrical spots, indicating that the crystallization direction of the CFs is unidirectional. Both the HRTEM and SAED patterns suggest the formation of 1D carbon nanofibers with radial (002) orientation. In contrast, commercial multiwalled carbon nanotubes (CNTs; Figure 2i) show lattice fringes of about 0.34 nm, which are parallel to the 1D direction (Figures 2j and 2k). Compared with the CNTs with axially oriented arrangement, the evolved crystallization direction of the CFs with radial (002) orientation can expose more active edge-plane sites for K-ion storage and shorten the K diffusion distance. The K-ion storage properties of these as-obtained carbon-based materials were evaluated by employing potassium metal as the counter and reference electrode. Figure 3a and Supporting Information Figure S9 reveal the K storage properties of the CFs. The CV curves were obtained at 0.1 mV s−1 ( Supporting Information Figure S9). A reduction peak located at 0.01 V may be attributed to the K insertion of the layer of CFs. The oxidation peak at 0.7 V can be assigned to the K extraction reaction. The charge/discharge curves (Figure 3a) are in agreement with the results of the CV curves. The initial potassiation and depotassiation capacities of the CFs are 746.1 and 315.4 mAh g−1, respectively. The first coulombic efficiency is 42.3%. This low first coulombic efficiency may be ascribed to the formation of a solid electrolyte interface (SEI) film and the irreversible capacity in the porous region of the active materials.49 After 60 cycles, the reversible capacity of the CFs is 273.9 mAh g−1 at a current density of 100 mA g−1 (Figure 3b). In contrast, the capacity of the CSs sample is 212 mAh g−1 after 60 cycles, while the CNTs deliver a reversible capacity of only 118.8 mAh g−1 after 30 cycles. Compared with the CSs, the enhanced reversible capacity of the CFs could be ascribed to the N-doped defects, 1D structure, and BET specific surface area. These advantages are in favor of the potassium diffusion kinetics ( Supporting Information Figure S10). Compared with the CNTs, the enhanced capacity is ascribed to the radially oriented (002) construction, which can also increase the active edge-plane sites for K storage. The synergistic effects of both advantages endow the CFs with improved reversible capacity. Figure 3 | The K storage performance of the CFs, CSs, and CNTs in half-cell PIBs. (a) Change/discharge curves of the CFs at 50 mA g−1 during the first two cycles. (b and c) Comparison of the (b) cycling stability at 100 mA g−1 and (c) rate capability from 50 to 2000 mA g−1 of the CFs, CSs, and CNTs. (d and e) Cycling performance of the CFs at (d) 100 mA g−1 for 500 cycles and (e) 1000 mA g−1 for 1200 cycles. (f) Comparison of the log(i) − log(v) curves of CFs and CNTs. (g) Capacitive contribution and total capacity at 1.0 mV s−1 for CFs. (h) Comparison of the normalized capacitive contribution ratio of CFs, CSs, and CNTs from 0.2 to 1.0 mV s−1. Download figure Download PowerPoint Figure 3c compares the rate capacities of the CFs, CSs, and CNTs with increasing current densities from 50 to 2000 mA g−1. It is shown that CFs sample exhibits the best rate capability. At 50 mA g−1, the CFs can deliver an average capacity of 344 mAh g−1, which is much higher than that of the CSs (257 mAh g−1) and CNTs (220.7 mAh g−1), respectively. The reversible capacities of the CFs are maintained at 174 and 159 mAh g−1 at 1000 and 2000 mA g−1, respectively. To evaluate the high rate and long-life performance of the CFs, the half-cell was cycled at 100 and 1000 mA g−1 for 500 and 1200 cycles, respectively. As presented in Figure 3d, a reversible capacity of 252.8 mAh g−1 was maintained after 500 cycles at 100 mA g−1 with high coulombic efficiency. The capacity of CFs is competitive with the previously reported carbon-based anodes for PIBs ( Supporting Information Table S1). Furthermore, even at a high current density of 1000 mA g−1 for 1200 cycles, a discharge capacity of 181.5 mAh g−1 was also maintained (Figure 3e), suggesting the excellent cycling and rate performance of the CFs. To study the good rate performance of the CFs, the redox pseudocapacitance-like contribution is comparatively studied, as presented in Figures 3f and 3g and Supporting Information Figures S11 and S12. Supporting Information Figures S11A, B and S12A are the CV curves of the CFs, CNTs, and CSs tested at ever-increasing scan rates from 0.2 to 1.0 mV s−1, respectively. According to the relationship between current (i) and scan rate (v): i = a * vb, the b-value can be calculated with logarithmic mathematical manipulation.50 After calculations, the b-values of the CFs and CNTs at peak one are determined to be 0.95 and 0.83, respectively (Figure 3f), which suggests that the K-ion storage is mainly dominated by the capacitive capacity contribution. It can be seen that the b-value of CFs is slightly higher than that of the CNTs. Moreover, according to the equation: i(V) = k1v+k2v1/2, the capacitive contribution among the total capacity can be further quantified by separating the current response (i) into capacitive effects (k1v) and diffusion-controlled reactions (k2v1/2) at a fixed potential (V).51 Figure 3g and Supporting Information Figure S12B, S11E show the typical voltage profile at 1.0 mV s−1 of CFs, CSs, and CNTs, respectively. Figure 3h compares the capacitive contribution from 0.2 to 1.0 mV s−1 between the CFs and CNTs. It is observed that the capacitive capacity contribution of the CFs is higher than that of the CSs and CNTs, which further suggests the exposed active edge plane sites and N-doping defects can boost the K-ion storage. To explore why the CFs deliver good cycling and rate capability, the ex situ Raman spectrum, HRTEM, and SEM images are provided and discussed. Figure 4a comparably shows the rate discharge curves of the CFs and CNTs for the first cycle at current densities of 0.5, 1, and 2 A g−1, respectively. Figures 4b and 4c are the corresponding Raman spectra at various current densities after first discharge. It is obvious that the rate capacity of CNTs is poorer than that of the CFs, especially at 1 and 2 A g−1. The Raman spectrum further confirms the enhanced rate capability of the CFs. For CFs (Figure 4b), the intensity ratios of ID/IG at ever-increasing current densities from 0.1 to 2 A g−1 have no obvious differences, which suggests that the insertion of K ions at low and high current density is similar. This result confirms the good rate capability of the CFs, which agrees with the results in the top half of Figure 4a. However, for the CNTs (Figure 4c), the ID/IG value at 0.1 A g−1 is higher than that at 0.5, 1, and 2 A g−1, suggesting unsatisfied K insertion into CNTs at high current density and poor rate capability.52 Figure 4 | (a) The rate discharge curves of CFs and CNTs for the first cycle at current densities of 0.5, 1, and 2 A g−1, respectively (top half, CFs; half bottom, CNTs). (b and c) Raman spectra of (b) CFs and (c) CNTs after first discharge at 0.1, 0.5, 1.0, and 2.0 g−1. (d–i) HRTEM images of the (d–f) CFs and (g–i) CNTs at different stages, (d and g) after first discharge, (e and h) after first recharge, and (f and i) after 20 cycles. The inset of (f and i) is the corresponding SAED image. (j) The SEM image and (k) EDS element mapping of the CFs after cycling. (l and m) Schematic representation of the K insertion into (l) CNTs with axially oriented (002) carbon construction, and (m) CFs with radially oriented (002) carbon construction. Download figure Download PowerPoint Figures 4d–i are the HRTEM images of the CFs and CNTs during different charge and discharge stages. After the first discharge, the (002) lattice of the CFs enlarged to be 0.44 nm (Figure 4d) from the initial 0.36 nm (Figure 2h), suggesting the formation of KCx compounds. As for the CNTs, the (002) lattice size increased to be 0.366 nm (Figure 4h) from the initial 0.34 nm (Figure 2k). Compared with the CNTs, the proportion of enlarged (002) lattice is more obvious, suggesting more K ions insertion and higher capacity. After the recharge process (Figures 4e and 4h), the lattice fringes of the CFs and CNTs returned to about 0.39 and 0.342 nm, respectively, implying reversible K intercalation/deintercalation cycles. After 20 cycles, the lattice fringe of the CFs was well constructed and perpendicular to the 1D direction (Figure 4f), which was also confirmed by the SAED images ( Supporting Information Figure S13 and inset of Figure 4f). The well-maintained structure during the charge/discharge process suggests good cycling performance of the CFs. However, for the CNTs after 20 cycles, some of the lattice fringes (cycled in white) are disordered (Figure 4i), which may be attributed to the repeated intercalation/deintercalation that results in structure damage and fading capacity. The TEM images during different charge/discharge stages are given, as exhibited in Supporting Information Figure S14. It is obvious that the microstructure of the CFs is well-maintained during the discharge and recharge process. Furthermore, the SEM images after 50 cycles are also provided. As presented in Supporting Information Figure S15A, the surface of the CSs after cycling becomes coarse and crinkly, which may be due to the volume change during repeated potassiation and depotassiation. For the CFs (Figure 4j), there is no obvious morphology change after 50 cycles. Moreover, all the elements, C, N, and K, are uniformly distributed after repeated cycling (Figure 4k). All results confirm the good structure stability of the CFs. Above all, the outstanding cycling and rate capacity of the CFs can be mainly attributed to the different (002) carbon orientation compared with that of the CNTs, as the schematic shows in Figures 4l and 4m. CFs with radially oriented (002) carbon construction, compared with the CNTs with axial (002) orientation, can induce more active edge-plane sites, which favors K ions insertion and increases the reversible capacity. The orthogonal orientation could also shorten the K diffusion distance, thus enhancing the rate performance of the CFs. To disclose the effects of the N-doped defects, K adsorption abilities on graphene with ideal, defective and N-5/N-6-doped sites were simulated. The relative adsorption energy between these graphene-based layers and one K atom was also calculated. In contrast, the adsorption energy of flawless graphene was calculated to be –0.5 eV (Figure 5a), while the graphene with defective sites was calculated to be –2.5 eV (Figure 5b). Moreover, when the N-5/N-6-doped sites were further introduced within the defective graphene, the adsorption energy between the sample and K atom was calculated to be −3.2 eV (Figure 5c). It is obvious that the adsorption energy of the sample with N-5 and N-6 sites is higher, suggesting stronger attractions to the K atom. Figure 5 | Theoretical simulations of N-doped defects on the K adsorption abilities. (a–c) Side and top views of K atoms absorbed in the (a) ideal, (b) defective, and (c) N-5/N-6-doped carbon structure. (d–f) Side views of electron density differences with the same level of isosurfaces of K absorbed in the (d) ideal, (e) defective, and (f) N-5/N-6-doped carbon structure. Download figure Download PowerPoint The N-doped defects not only improve the adsorption energies but also decrease the activation barrier. As presented in Figures 5d–f and Supporting Information Figure S16, the N-doped defects change the electronic structure of the sample. In the ideal graphene (Figure 5d), a net gain of electronic charge between the K atoms and the graphene is observed, suggesting charge transfer from the K atom to the nearest neighboring C atoms. However, for the N-5/N-6-doped carbon, the charge transfers to the bonding carbon atoms within the N-doping and accumulates around the N-doped sites. All results suggest that the N-5/N-6-doped sites in graphene provide a fast electron-transfer pathway, which is in favor of enhancing the K storage performance.53 Above all, the N-doping defects can efficiently enhance the K adsorption kinetics and decrease the K diffusion barrier, thus theoretically boosting the K storage performance. The overall performance (especially the good rate ability) of CFs suggests practical application in PIHCs. Commercial CMK3 acts as the capacitive-type cathode, the prepotassiated CFs are employed as the battery-type anode, and the dual-carbon PIHC device is fabricated. The rate and cycling performance of commercial CMK3 in half-cell PIBs are presented in Supporting Information Figures S17 and S18. Figure 6a displays the charge/discharge curves under different current densities of 0.1–2.0 A g−1. It is noted that the K intercalation/deinte