Shipeng Wang, Ziyan Liu, Qiangchun Liu, Baojun Wang, Wei Wei, Hao Wu, Zijie Xu, Shikuo Li, Fangzhi Huang, and Hui Zhang, Promoting the microwave absorption performance of hierarchical CF@NiO/Ni composites via phase and morphology evolution, Int. J. Miner. Metall. Mater., 30(2023), No. 3, pp.494-503. https://dx.doi.org/10.1007/s12613-022-2524-2
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With the rapid development of electromagnetic technology, electromagnetic waves (EMWs) have been widely used in satellite communications, antenna measurement, and sensing, which are convenient for our daily lives and industrial manufacturing. However, high-energy electromagnetic irradiation and unnecessary electromagnetic interference have caused severe electromagnetic pollution, which not only affects the normal operation of electronic equipment but also seriously endangers human health [1–5]. Accordingly, significant efforts have been made to design and prepare strong microwave absorbing materials (MAMs) to tackle the above issues. In general, high-efficiency MAMs are expected to simultaneously satisfy the demands of high absorption ability, wide absorption band, lightweight, and less thickness [6–9].
Depending on the drastic magnetic loss, traditional metal-based materials, such as ferrite and metal powder, have been developed to absorb EMWs [10–12]. Although they possess fascinating microwave absorption (MA) performance in particular frequencies, the high density and poor chemical stability severely restrict their practical utilization as lightweight and high-efficiency absorbers. Nanostructured carbon materials, such as graphene, carbon nanotube, and carbon fiber (CF), have gradually become ideal MAMs due to their superior dielectric properties and ultra-low density [13–16]. Among them, CF possesses a low percolation threshold and excellent dielectric loss due to the one-dimensional (1D) structure and high conductivity, which are sought after in advanced MAMs [17–18]. π electrons can migrate in the whole CF matrix in the form of free electrons and generate induced dipole moments and polarization current under the action of the EM field, thus consuming EMWs. However, pure CF employed as an absorbent has two drawbacks. First, the high electrical conductivity would greatly impair impedance matching, leading to strong reflection rather than absorption. Second, the single dielectric loss is not conducive to broadband absorption [19–20]. In light of these issues, combining CF with a magnetic metal or metal oxide may be an effective solution for improving impedance matching and magnetic loss [21–22]. For example, Chen et al. [23] prepared FeCo/CoFe2O4/CNF composites using the electrospinning method, and the obtained composites displayed an excellent MA performance with an absorption peak of −52.3 dB. Bandaru et al. [24] fabricated Fe3O4/cotton-derived CF composites, which displayed a minimum reflection loss (RLmin) of −56.8 dB. Zhao et al. [25] reported Co3O4/N-doped carbon/short CF that showed an optimum MA with an RLmin value of −29 dB and an efficient absorption bandwidth (EAB) of 5.44 GHz. However, most of the reported magnetic elements on CFs are irregular or agglomerated, which resulted in higher filler loading (usually higher than 20wt%) in the MA performance. This drawback restricts the further applications of lightweight requirements in aerospace and aviation.
Constructing hierarchical nanoarrays on CF may significantly preserve the lightweight feature due to their large surface area and special three-dimensional (3D) micro/nano structure. In this study, hierarchical CF@NiO/Ni composites were successfully prepared through the in situ
CF (Huangyu Electromagnetic Technology Co., Ltd., China) was first soaked in acetone to remove impurities and then immersed in an HNO3 solution for surface functionalization.
Typically, 1.0 mmol nickel chloride hexahydrate (NiCl2·6H2O) and 1.0 mmol hexamethylenetetramine (HMTA) were dissolved in 30 mL deionized water, and then 0.1 g CF was dispersed into the above solution. After sonication for 30 min, the resulting suspension mixture was transferred into a Teflon-lined autoclave and heated at 90°C for 10 h. Finally, light-green products were repeatedly washed with distilled water and dried at 80°C.
Hierarchical CF@NiO/Ni composites were prepared through an annealing process. In detail, the as-prepared CF@Ni(OH)2 precursors were calcined under the nitrogen gas flow at the targeted temperatures (300, 400, 500, and 600°C) for 2 h. Based on the composition from the later characterization and the pyrolysis temperature, the as-prepared samples were denoted as CF@NiO-300, CF@NiO/Ni-400, CF@NiO/Ni-500, and CF@Ni-600.
The crystal structure of the samples was identified via X-ray diffraction (XRD, Rigaku Smartlab Beijing Co., LTd., China). The morphologies and nanostructures were investigated via field emission scanning electron microscopy (FE-SEM, Zeiss, Sigma 500, Germany) and transmission electron microscopy (TEM, FEI Tecnai TF20, USA). The element distribution was conducted with an energy-dispersive X-ray spectrometer (EDS, Oxford Ultim Max 40, England) attached to the SEM. The graphitization degree of carbon was measured by a Raman spectrometer (inVia Reflex, England). X-ray photoelectron spectroscopy (ESCALAB 250, USA) was adopted to perform elemental analysis and determine the composition of the material. The composites (3wt%) were mixed with molten paraffin and pressed into toroidal specimens with an outer diameter of 7.00 mm and an inner diameter of 3.04 mm, and their electromagnetic parameters were measured using a vector network analyzer (Agilent E8363B, USA) in 2–18 GHz. Monostatic RCS simulations were performed on the commercial Computer Simulation Technology (CST) Studio Suite 2020. A 180 mm × 180 mm × 2 mm perfect electric conductor (PEC) was used as a metal plate to reflect the EMWs.
Fig. 1(a) depicts the preparation procedure of hierarchical CF@NiO/Ni composites. First, the hierarchical CF@Ni(OH)2 precursors are fabricated through a hydrothermal method. During the hydrothermal process, the HMTA is decomposed to generate an alkaline environment. Then, Ni2+ and OH− in the solution system combine to form layered metal hydroxides (Fig. S1(a)), which are vertically anchored on the CF’s surface. Second, the obtained CF@Ni(OH)2 precursors are directly annealed under an N2 atmosphere. The Ni(OH)2 can be converted into NiO or metallic Ni depending on the annealing temperature. Meanwhile, nanosheets form porous structures due to high temperature dehydration decomposition. The XRD patterns (Fig. 1(b)) show a broad peak at around 25.5°, which corresponds to the (002) plane of the CF substrate [26]. There are three typical characteristic peaks at the 2θ values of 37.26°, 43.28°, and 62.72° (denoted with ♣) in CF@NiO-300 correspond to the (101), (012), and (110) planes of NiO (JCPDS No. 44-1159). When the annealing temperature is 400°C, besides the NiO diffraction peaks, two discernable small diffraction peaks appear at around 44.50° and 51.85° (denoted with ♦), which correspond to the (111) and (200) crystal planes of Ni, respectively (JCPDS No. 04-0850). The phase transition of NiO into Ni is demonstrated due to the reduction of carbon at high temperatures. For CF@NiO/Ni-500, the intensity of the Ni diffraction peak is reinforced, implying that more Ni is produced in the composite. Finally, when extending the annealing temperature to 600°C, the diffraction peaks belong to NiO disappeared, and only Ni diffraction peaks can be observed, demonstrating the total conversion. The XRD results indicate that the phase structure of the CF@NiO/Ni composites can be modulated by adjusting the annealing temperature.
Fig. 2 presents the morphologies of the hierarchical CF@NiO/Ni composites. As shown in Fig. 2(a), the CF@NiO-300 displays uniform leaf-like morphology, and the vertical NiO nanosheets have a flat and smooth surface which is similar to Ni(OH)2 precursors (Fig. S1(b)). Moreover, these nanosheets are cross-linked, forming a stable 3D microporous structure, which allows incident microwaves to be reflected multiple times and thus consumed. At annealing temperatures of 400 and 500°C, the sheet-like morphology of NiO/Ni can still be preserved but become loose and porous (Fig. 2(b) and (c)). As shown in Fig. 2(c), the edges of the sheets of CF@NiO/Ni-500 become slightly curly, resulting from the conversion of NiO during the annealing process. However, when the temperature continued to increase, the small Ni nanocrystals agglomerated and formed irregular nanoparticles, leading to the destruction of the nanosheet for CF@Ni-600 (Fig. 2(d)). The TEM image also shows the porous structure of the NiO/Ni, and the discrete Ni nanoparticles are deposited on the porous NiO flakes (Fig. 2(e)). In addition, the high-resolution (HR-TEM image (Fig. 2(f)) exhibits two types of lattice fringes with interlayer distances of 0.24 and 0.20 nm, referring to the (101) crystal plane of NiO and the (111) crystal plane of metal Ni, respectively [27]. The HR-TEM results confirm the multi-heterojunction interface between NiO and Ni, which is in accordance with the former XRD results and the morphology revolution corresponding to the phase modulation. The element mapping in Fig. S2 indicates the coexistence and homogeneous distribution of the Ni and O elements on the CF surface.
Furthermore, Raman spectroscopy was used to illustrate the graphitization degree of C in the CF and CF@NiO/Ni-500 composite. In Fig. 3(a), two broad peaks at ~1340 and 1580 cm−1 arise from the disordered carbon (D band) and graphitic structure (G band). In particular, the CF@NiO/Ni-500 has a prominent peak at ~485 cm−1, which is assigned to the first-order longitudinal optical (1LO) mode of Ni–O [28–29]. The intensity ratios of D band and G band (ID/IG) are 0.8 and 0.93 for CF and CF@NiO/Ni-500, respectively. This finding reveals that the original sp2 hybridization system on the annealed CF surface is disrupted to produce more defect sites. These can act as polarization centers to provoke dipole polarization and thus consume electromagnetic energy [30].
The survey XPS spectrum in Fig. 3(b) proves the presence of the Ni, C, and O elements in CF@NiO/Ni-500, which is consistent with the EDS results. The electron interaction between the heterostructures was investigated by analyzing high-resolution Ni 2p spectra. For CF@NiO-300, the detailed Ni 2p spectrum is deconvoluted into two characteristic peaks at 855.5 and 872.6 eV, which are ascribed to Ni 2p3/2 and Ni 2p1/2, respectively. The peaks located at 860.9 and 879.3 eV (denoted as “Sat.”) are the satellites of the Ni 2p3/2 and Ni 2p1/2, respectively [31–32]. In particular, there is no Ni0 peak in the CF@NiO-300, confirming the absence of Ni metal, which is in accordance with the XRD result. Meanwhile, Ni 2p in CF@NiO/Ni-500 shows a new split peak at approximately 852.7 eV, corresponding to the metallic Ni characteristics, demonstrating the coexistence of metallic and oxidized states. Compared with CF@NiO-300 (853.4 and 855.5 eV for Ni), the binding energy of Ni2+ in CF@NiO/Ni-500 appears slightly positive shifts (ca. 0.3 eV for Ni2+ and 0.26 eV for Ni3+), indicating the intense electronic interaction and more charge aggregation/redistribution at the NiO–Ni contact interface (Fig. 3(c)). Differences in work functions have been reported in interfacial charges’ redistribution and accumulation, resulting in interfacial polarization [33–34]. In general, the work function of metals is higher than that of the corresponding oxide semiconductors. In the NiO/Ni system, the NiO semiconductor energy band is bent to form a Schottky barrier at the heterogeneous interface [35]. The Schottky barrier causes a high interfacial contact resistance, which hinders electron hopping and enriches positive and negative charges at the opposite parties of the interface (Fig. 3(d)). When subjected to periodic electromagnetic fields, positive and negative charges dissipate electromagnetic energy through a strong polarization behavior. Due to the high electron mobility of CF, a similar polarization phenomenon occurs at the Ni–C interface.
To verify the potential feasibility of the system in improving the MA performance, the frequency-dependent electromagnetic parameters (including real parts (ε′) and imaginary parts (ε″) of the complex permittivity and real parts (μ′) and imaginary parts (μ″) of the relative complex permeability) of CF and CF@NiO/Ni composites were measured. Due to the high electrical conductivity and high aspect ratio of CF, a conductive network can be formed even at low filler loadings [36]. As observed in Fig. S3, the ε″ values of the pure CF are too high, which is detrimental to the microwave attenuation owing to the poor impedance matching [37–38]. Nonetheless, after coating the porous NiO/Ni nanosheet arrays, the ε′ and ε″ of the composites decrease (Fig. 4(a) and (b)). This phenomenon can be interpreted with the high resistivity of NiO/Ni nanosheet arrays adequately preventing the conductive network’s formation and reducing the CF’s space charge polarization [39]. Furthermore, the ε′ value dramatically increases from 4.6 (CF@NiO-300) to 10.3 (CF@NiO/Ni-400), 12.4 (CF@NiO/Ni-500), and 16.6 (CF@ Ni-600) at 2.0 GHz, while the ε″ dramatically increases from 0.08 (CF@NiO-300) to 6.0 (CF@NiO/Ni-400), 7.7 (CF@NiO/Ni-400) and 13.8 (CF@Ni-400) at 2.0 GHz, suggesting that the increment metallic Ni content can strengthen the dielectric properties. The ε′ curve of CF@NiO/Ni-500 exhibits a fluctuation at approximately 9 GHz, which is probably related to the interfacial polarization caused by the NiO–Ni contact interface and porous structure [40]. To further explain the dielectric behavior of CF@NiO/Ni composites, the Debye model is analyzed based on the measured relative complex permittivity. Under the action of the periodic electric field, the electric dipole will move along with the changes in the electric field. When the electric field frequency increases to a certain degree, the change in the dipole will gradually fail to keep up with the change in the periodic frequency of the electric field, thus causing the hysteresis phenomenon and finally reaching the limit. This process can cause the dissipation of electromagnetic energy at high frequencies and even ultra-high frequencies. Thus far, the Debye equation is recognized as a practical theory for understanding polarization loss processes [41]:
| (ε′−εs+ε∞2)2+(ε″ | (1) |
where εs represents the permittivity at the electrostatic field and ε∞ represents the permittivity at the high-frequency limit.
According to Eq. (1), the semicircle in the ε′ versus ε″ curve is usually called the Cole–Cole semicircle showing a Debye relaxation process. The CF@NiO/Ni-400 and CF@NiO/Ni-500 composites have three semicircles (Fig. S4), which are numerically superior to those in CF@NiO-300 (one) and CF@Ni-600 (two), implying that the former has more polarization relaxation processes. The increased polarization behavior may originate from two aspects: First, the NiO/Ni nanosheet arrays immobilized on the CF surface provide abundant Ni-C contact sites. Second, the NiO-to-Ni conversion creates several Schottky contact interfaces.
The μ′ values of the samples fluctuate between 0.81 and 1.05, and the μ″ values vary between ‒0.07 and 0.3 (Fig. 4(c) and (d)). Moreover, there are several apparent fluctuations appearing at some special frequencies, which are likely to be related to the ferromagnetic resonance and eddy current effect. If the eddy current loss is the primary source of the magnetic loss, the eddy current loss C0 (
Quantitative parameters such as RL can be used to assess MA performance. The RL of the CF@NiO/Ni composites is calculated as follows [42–44]:
| \mathrm{R}\mathrm{L}=20 \cdot \mathrm{l}\mathrm{g}\left|\frac{{Z}_{\mathrm{i}\mathrm{n}}-{Z}_{0}}{{Z}_{\mathrm{i}\mathrm{n}}+{Z}_{0}}\right| | (2) |
| {Z}_{\mathrm{i}\mathrm{n}}={Z}_{0}\sqrt{{\mu }_{\mathrm{r}}/{\varepsilon }_{\mathrm{r}}}\mathrm{t}\mathrm{a}\mathrm{n}\mathrm{h}\left[\mathrm{j}\left(2 \text{π} fd/c\right)\sqrt{{\mu }_{\mathrm{r}}{\varepsilon }_{\mathrm{r}}}\right] | (3) |
where Z0 and Zin are the impedance of air and the absorber, and f, d, and c represent the frequency of EMWs, thickness, and velocity of light, respectively. εr and μr are the complex permittivity (εr = ε′ – jε″) and complex permeability (μr = μ ′ – jμ ″), respectively.
Fig. 5(a)–(h) shows the MA performance of specimens containing 3wt% composites with various thicknesses. For pure CF, the RLmin value is only −5.74 dB (Fig. S6), and such a weak MA property is attributed to its poor impedance matching. As displayed in Fig. 5(a) and (e), the CF@NiO-300 merely holds a negligible absorption performance, which can be explained by the insufficient low permittivity to effectively dissipate EMWs. After adequately adjusting their phase structure, the composites significantly exhibit an enhanced MA. For CF@NiO/Ni-400, the optimal absorption peak of −25.85 dB is significantly more potent than those of CF and CF@NiO-300, and its EAB can reach 6.40 GHz with a matching thickness of 2.25 mm. When the temperature increases to 500°C, the CF@NiO/Ni-500 composite exhibits a powerful electromagnetic absorption. It has an optimal absorption peak value of about −43.92 dB at 17.72 GHz. When the matching thickness is 1.80 mm, its broadest EAB is 5.64 GHz from 12.36 to 18 GHz, covering the whole Ku band. These outstanding achievements make the CF@NiO/Ni-500 composite a promising candidate for EM absorption. However, as the annealing temperature further rises, the absorption intensity and EAB decrease. As shown in Fig. 5(d) and (h), CF@Ni-600 shows inferior performance to CF@NiO/Ni-500, and its RLmin value is only −11.78 dB, and the EAB is 4.08 GHz. Together, the composites’ phase transitions and unique porous structures likely regulate their diverse MA performance. To clearly and intuitively display the MA performance of the samples, the property parameters are listed in Table 1. These data demonstrate that successfully modulating the phase structure and optimizing the porous structure have great potential in preparing advanced carbonaceous absorbing materials.
| Sample | RL / dB | Frequency / GHz | Thickness / mm | EAB / GHz |
| CF | −5.74 | 18.00 | 1.50 | 0 |
| CF@NiO-300 | −2.41 | 17.72 | 2.60 | 0 |
| CF@NiO/Ni-400 | −25.85 | 7.44 | 2.25 | 6.40 |
| CF@NiO/Ni-500 | −43.92 | 17.72 | 1.80 | 5.64 |
| CF@Ni-600 | −11.78 | 17.68 | 1.80 | 4.08 |
Generally, poor impedance matching causes intense reflection at the front surface, whereas low attenuation capability results in weak loss capacity. Therefore, proper impedance matching characteristics and strong EMW attenuation are prerequisites for developing an ideal absorber. Thus, the normalized impedance matching ratio (Zr = |Zin/Z0|) and the attenuation coefficient (α) of the CF@NiO/Ni composites are calculated to provide a basis for the further elucidation of the enhanced MA performance [45–46].
| {Z_{\rm{r}}} = \left| {\frac{{{Z_{{\rm{in }}}}}}{{{Z_0}}}} \right| = \left| {\sqrt {{\mu _{\rm{r}}}/{\varepsilon _{\rm{r}}}} \tanh \left[ {{\rm{j}}(2 \text{π} fd/c)\sqrt {{\mu _{\rm{r}}}{\varepsilon _{\rm{r}}}} } \right]} \right| | (4) |
| \begin{aligned} & \alpha = \frac{{\sqrt 2 \text{π} f}}{c} \times \\ & \quad \sqrt {\left( {{\mu ^{\prime \prime }}{\varepsilon ^{\prime \prime }} - {\mu ^\prime }{\varepsilon ^\prime }} \right) + \sqrt {{{\left( {{\mu ^{\prime \prime }}{\varepsilon ^{\prime \prime }} - {\mu ^\prime }{\varepsilon ^\prime }} \right)}^2} + {{\left( {{\mu ^{\prime \prime }}{\varepsilon ^\prime } + {\mu ^\prime }{\varepsilon ^{\prime \prime }}} \right)}^2}} } \end{aligned} | (5) |
The Zr and α curves of all the samples are shown in Fig. 6(a) and (b). From Fig. 6(a), CF@NiO/Ni-500 exhibits a satisfactory impedance matching characteristic, implying that most incident EMWs penetrate the inside of the CF@NiO/Ni-500 matrix and thereby are attenuated. Although CF@Ni-600 possesses the highest α (Fig. 6(b)), the unsatisfactory Zr produced a poor MA performance. Therefore, the enhanced MA capabilities of the CF@NiO/Ni-500 composite can be attributed to the high attenuation coefficient and good impedance matching.
In addition, electromagnetic energy can be consumed by the “geometric effect.” This special loss is called quarter-wavelength cancellation (nλ/4), which can be described as
| {t}_{\mathrm{m}}=\frac{n\lambda }{4}=\frac{nc}{{4f}_{\mathrm{m}}\sqrt{\left|{\varepsilon }_{\mathrm{r}}\right|\left|{\mu }_{\mathrm{r}}\right|}} \; (n=1,3, 5,\dots ) | (6) |
where tm and fm are the thickness and frequency corresponding to the RLmin peaks, λ is the wave-length of microwave, |εr| and |μr| are the modulus of εr and μr, and n represents the positive odd number [47–48]. Fig. S7 shows the RL and tm versus fm curves of the CF@NiO/Ni-500 composites. All the experimental thicknesses (red pentacle) are precisely located on the simulated tm curves (green line), which demonstrates that the quarter-wavelength cancellation theory has an important reference for analyzing the enhancement of the CF@NiO/Ni-500 MA performance properties.
The structural model of the CF@NiO/Ni composites and the possible MA mechanism are shown in Fig. 7(a). The enhancement of the MA performance can be summarized as follows: (1) The metal oxide array coating and its phase structure transformation optimize the impedance matching characteristics. This result is ascribed to the high resistivity of NiO/Ni nanosheet arrays, which appropriately prevent the establishment of a conductive network and reduce the space-charge polarization of CF. (2) Several heterogeneous interfaces are generated during the phase transition process, and the induced charge accumulation and redistribution at these interfaces promote the interfacial polarization loss. (3) The effective cooperation of dielectric and magnetic losses endows CF@NiO/Ni composites with an excellent MA behavior. (4) The unique 3D porous configuration of NiO/Ni arrays can provide multiple reflecting and scattering abilities, resulting in more incident EMWs being consumed.
Recently, a series of objective and precise criteria, including specific reflection loss SRLt (RL/layer thickness), SRLft (RL/(filler loading × layer thickness)), and SRLBft (RL × EAB/(filler loading × layer thickness)), have been proposed to further comprehensively evaluate the MA performance of MAMs [49–51]. Based on these methods, the |SRLt|, |SRLft|, and |SRLBft| of CF@NiO/Ni composites and other representative CF-based MAMs are calculated, and the results are shown in Fig. 7(b)–(d). It can be seen that the CF@NiO/Ni-500 exhibits superior performance to its counterparts, implying its great prospect as a high-efficiency MAM. The detailed MA parameters are listed in Table S1.
In practical applications, EMWs are obliquely incident on the coating at certain angles. Therefore, in aircraft stealth design, EMWs at different angles of incidence need to be counteracted [52]. A CST program is used to simulate RCS with characteristic sample coatings and incidence angles from −90° to 90° (Fig. 8(a)). As shown in Fig. 8(b), the RCS values of the PEC and CF are high, indicating their strong reflection and low absorption characteristic. By contrast, the RCS values are significantly reducing after coating CF@NiO/Ni, indicating that the porous hierarchical CF@NiO/Ni structure positively impacts MA. Similarly, in Fig. S8(a)–(d), the model coated with CF@NiO/Ni-400 has the strongest RCS reduction in all directions, which is attributed to its broad absorption bandwidth.
In summary, porous NiO/Ni nanosheet arrays are successfully constructed on the CF surface using the hydrothermal method and subsequent annealing treatment. The phase and morphology of the coating can be tuned by adjusting the annealing temperature, which significantly affects the complex permittivity and impedance matching of hierarchical CF@NiO/Ni composites. The tailored composition and hierarchical structure synergistic effect endow the CF@NiO/Ni-500 with a superior MA performance. In particular, the CF@NiO/Ni-500 achieves an RLmin value of −43.92 dB and a broad EAB of 5.64 GHz (covering the whole Ku band) at an ultra-low loading of 3wt%. The simulated RCS results reveal that the strong electromagnetic scattering can be effectively suppressed by employing CF@NiO/Ni composites on the PEC. We hope that this work could inspire the design of novel CF-based composites for MA via phase and morphology evolution.
This work was financially supported by the National Natural Science Foundation of China (Nos. 51872002 and 52172174), and the key research and development projects in Anhui province, China (No. 202004a07020026). We also gratefully acknowledge the support of Joint Laboratory of Electromagnetic Material Structure Design and Advanced Stealth Technology.
The authors declare no competing financial interest.
The online version contains supplementary material available at https://doi.org/10.1007/s12613-022-2524-2.
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| Sample | RL / dB | Frequency / GHz | Thickness / mm | EAB / GHz |
| CF | −5.74 | 18.00 | 1.50 | 0 |
| CF@NiO-300 | −2.41 | 17.72 | 2.60 | 0 |
| CF@NiO/Ni-400 | −25.85 | 7.44 | 2.25 | 6.40 |
| CF@NiO/Ni-500 | −43.92 | 17.72 | 1.80 | 5.64 |
| CF@Ni-600 | −11.78 | 17.68 | 1.80 | 4.08 |
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