MOF-derived yolk–shell Ni/C architectures assembled with Ni@C core–shell nanoparticles for lightweight microwave absorbents

Xiaolei Wang*a, Qiyao Genga, Guimei Shia, Yajing Zhangb and Da Li*c
aSchool of Environmental and Chemical Engineering, Shenyang University of Technology, Shenyang 110870, PR China. E-mail: xlwang@alum.imr.ac.cn
bCollege of Chemical Engineering, Shenyang University of Chemical Technology, Shenyang 110142, PR China
cShenyang National Laboratory for Materials Science, Institute of Metal Research, and International Centre for Materials Physics, Chinese Academy of Sciences, Shenyang 110016, PR China. E-mail: dali@imr.ac.cn

Received 24th August 2020 , Accepted 14th September 2020

First published on 17th September 2020


Yolk–shell Ni/C microspheres composed of Ni@C core–shell nanoparticles were successfully fabricated by decomposing a Ni-based metal–organic framework (Ni-MOF) at 500 °C and 600 °C. The Ni-MOF with a yolk–shell structure was prepared by a solvothermal method with an appropriate molar ratio of Ni(NO3)2·6H2O to C9H6O6 in the presence of PVP. The degree of crystallization of Ni and C was improved by increasing the pyrolysis temperature, which resulted in enhanced complex permittivity and optimized impedance matching of Ni/C microspheres for damping microwave. Meanwhile, the attenuation coefficient of Ni/C microspheres increased with the increment in pyrolysis temperature. The yolk–shell Ni/C microspheres obtained at 600 °C exhibited the optimal reflection loss (RL) reaching −39 dB with a bandwidth of 3.8 GHz (RL < −10 dB) at a thin matching thickness of 1.8 mm. The integrated bandwidth can achieve 12.3 GHz covering Ku-band (12–18 GHz), X-band (8–12 GHz), and most of C-band (5.7–8 GHz) with an appropriate thickness of 1.4–3.9 mm. Such excellent microwave absorption performance can be attributed to the synergistic effect of the magnetic and dielectric losses of Ni/C microspheres due to natural resonance, dipolar polarization and multiple interfacial polarizations at a unique yolk–shell interface, achieving the optimization of impedance matching and microwave attenuation. This work demonstrates that Ni/C microspheres with a desirable yolk–shell structure are potential candidates for the application in microwave absorption field.


Introduction

In recent years, numerous efforts have been made for solving the electromagnetic interference (EMI) issues due to the harmful effects on electronic devices and information security in military and civil fields. Microwave absorption materials (MAMs), which are distinct from the conventional electromagnetic shielding materials based on the reflection principle, can transform the unwanted electromagnetic energy into thermal energy through loss mechanism to eliminate EM pollution.1–4 Up to date, the design and fabrication of MAMs such as magnetic loss-based ferrites, soft ferromagnetic metals or alloys,5–7 dielectric loss-based metal oxides or sulfides,8,9 carbon materials,10 MXenes11 and magnetic/dielectric loss-based composites12,13 have been extensively progressed by constructing various components and microstructures, which can significantly tune the EM parameters and the corresponding microwave absorption performance. Magnetic/dielectric composites are expected to be promising candidates due to dielectric and magnetic losses of the components as well as interfacial polarization loss originating from the heterogeneous interface.14,15 Moreover, lightweight-featured magnetic/dielectric composites are an urgent requirement for high-performance microwave absorbents.

Porous yolk–shell structures with a distinct core@void@shell configuration have been extensively pursued due to the merits in reducing the absorbent density, modulating the electromagnetic parameters, improving the intrinsic impedance matching and enhancing multiple reflection and scattering attenuation of incident EM waves.16–18 Yolk–shell structures such as CoNi@TiO2,19 Ni@SnO2,20 ZnFe2O4@RGO@TiO2,21 FeCo@C,22 and Fe3O4@CuSiO3 (ref. 23) exhibited excellent microwave absorption performance. Compared with the oxide shell, the carbon-shell is favorable for lightweight, chemical stability and enhanced dielectric loss ability. Thus, the carbon-based yolk–shell composites should be more desirable for developing high-efficient microwave absorbents.24,25 More recently, the synthesis strategy based on the thermal decomposition of MOFs has emerged as an efficient method for preparing metal/carbon composites, which show excellent chemical homogeneity due to the coordination of metal ions and organic ligands on the atomic scale.26,27 The complex permittivity of composites can be effectively regulated via adjusting the pyrolysis temperature. Moreover, the magnetic nanoparticles can be homogeneously dispersed in the carbon matrix without aggregation and oxidation.28,29 For example, hollow Ni/C microspheres exhibited the minimum RL value of −57.25 dB with a bandwidth of 5.1 GHz at a thickness of 1.8 mm.30 Porous Ni@C composites have an RL value of −23.4 dB with an effective bandwidth of 4.68 GHz at a thickness of 1.9 mm.31 Yolk–shell Ni@C@ZnO microspheres showed the optimal RL of −55.8 dB and bandwidth around 3.5 GHz at a thickness of 2.5 mm.2 Although the introduction of ZnO brings extra interfacial polarization, the ZnO inevitably decreases the magnetization of composites and weakens the magnetic loss and the impedance matching to some extent. From a magnetic point of view, yolk–shell structured Ni/C microspheres composed of core–shelled Ni@C nanoparticles are favorable for good impedance matching, magnetic loss and thus high EM absorption performance.

In this work, yolk–shell Ni-MOFs with uniform particle size distribution were synthesized via a solvothermal route at a temperature of 180 °C for 10 h. The as-prepared Ni-MOFs were heated at different temperatures under N2 flow for 1 h to carbonize the organic framework. As a result, yolk–shell Ni/C microspheres composed of core–shell Ni@C nanoparticles were formed by spontaneous in situ carbon-thermal reduction reaction. These synthesis processes are illustrated in Scheme 1. The resultant yolk–shell Ni/C microspheres heated at 600 °C showed the optimal RL value of −39 dB at a thickness of 1.8 mm and bandwidth with RL values less than −10 dB reaching 3.8 GHz. The optimal bandwidth is 12.3 GHz covering 5.7–18 GHz by changing the absorbent layer thickness from 1.4 to 3.9 mm. The excellent microwave absorption performance could be attributed to the interfacial polarization loss of the unique yolk–shell microstructure and the synergistic effects of magnetic loss of Ni and dielectric loss of C.


image file: d0ce01242d-s1.tif
Scheme 1 Synthetic route of yolk–shell Ni/C microspheres.

Experimental

Synthesis of Ni-MOFs

All chemical reagents of analytical grade were purchased from Sinopharm Chemical Reagent Co. Ltd., Shenyang, China and used without further purification. In a typical preparation process according to the previous report with modification,32 0.04 mmol of PVP (K30), 1 mmol of C9H6O6 and 1.5 mmol of Ni(NO3)2·6H2O were in turn dissolved in a 30 mL of mixture solution (Vdistilled water[thin space (1/6-em)]:[thin space (1/6-em)]Vethanol[thin space (1/6-em)]:[thin space (1/6-em)]VDMF = 1[thin space (1/6-em)]:[thin space (1/6-em)]1[thin space (1/6-em)]:[thin space (1/6-em)]1, DMF is dimethylformamide) under mechanical stirring for 1 h. The light green solution obtained was transferred into a 50 mL Teflon-lined stainless-steel autoclave. The autoclave was then sealed and maintained at 180 °C for 10 h. The product was collected by centrifugation and washed several times with distilled water and ethanol. Finally, the product was dried in a vacuum oven at 50 °C.

Synthesis of Ni/C composites

The Ni-MOF was first thermally treated at 200 °C for 30 min and then heated to treatment temperatures of 500 °C, 600 °C and 700 °C, respectively, at a heating rate of 1 °C min−1 and maintained for 1 h with high-purity N2 (99.99%) protection. The resultant Ni/C composites were denoted as NC1, NC2 and NC3, respectively, according to the treatment temperature.

Characterization

Power X-ray diffraction (XRD) data were recorded using an XRD-6000 X-ray diffractometer (Shimadzu) equipped with a Cu Kα radiation source (40.0 KV, 30.0 mA). Scanning electron microscopic (SEM) images were collected using a Quanta 200S (FEI). Transmission electron microscopic (TEM) images were obtained using a Tecnai F20 operating at an accelerating voltage of 200 kV. Raman spectra were recorded using a confocal Raman spectroscopy system (Renishaw, in Via) with 633 nm laser. Nitrogen absorption–desorption isotherms were measured at 77 K using a Tristar II 3020 sorption analyzer (Micromeritics), and the specific surface area was determined by the Brunauer–Emmett–Teller (BET) method. The magnetic hysteresis loops were recorded using a LakeShore 7404 (LakeShore, USA) vibrating sample magnetometer (VSM). For the microwave absorption measurements, mixtures composed of the Ni@C sample and paraffin with a weight ratio of 2[thin space (1/6-em)]:[thin space (1/6-em)]3 were pressed into cylinders with an outer diameter of 7 mm, an inner diameter of 3.04 mm, and a thickness of 2 mm. An Agilent N5230A vector network analyzer was applied to determine the electromagnetic parameters in the frequency range of 1–18 GHz.

Results and discussion

Morphology and growth mechanism of Ni-MOF

The crystalline structure and morphology of the as-synthesized Ni-MOF are shown in Fig. 1. Fig. 1(a) clearly shows that the Ni-MOFs have a uniform spherical morphology with an average diameter of 3.1 μm. The rough surface demonstrates that the microspheres were aggregated from primary nanoparticles. A broken microsphere (inset in Fig. 1(a)) indicates a yolk–shell feature of the Ni-MOF. The TEM image in Fig. 1(b) further illustrates the detailed microstructure of the yolk–shell Ni-MOF. It can be seen that the microsphere has a core@void@shell configuration consistent with the observation of the SEM image. The sizes of the shell thickness, void and core are ca. 0.3 μm, 0.4 μm and 2.3 μm, respectively. The XRD pattern shown in Fig. 1(c) exhibits a main peak at 13.0° indexed to the Ni-MOF, revealing the presence of organic components.30–32
image file: d0ce01242d-f1.tif
Fig. 1 (a) SEM image, (b) TEM image and (c) XRD pattern of the Ni-MOF.

It was found that the molar ratio of raw materials can significantly influence the morphology of Ni-MOFs. When the molar ratio of Ni(NO3)2·6H2O to C9H6O6 is in the range of 1.5[thin space (1/6-em)]:[thin space (1/6-em)]0.5 to 1.5[thin space (1/6-em)]:[thin space (1/6-em)]1, the Ni-MOFs are mainly composed of uniform yolk–shell microspheres (Fig. S1(a) and 1(a)). When the molar ratio was gradually reduced to 1.5[thin space (1/6-em)]:[thin space (1/6-em)]2, the yolk–shell structure of microspheres gradually disappeared and showed irregular aggregation (Fig. S1(b)). Extra addition of C9H6O6 can accelerate the nucleation of Ni-MOFs and grows through a liquid-diffusion process, resulting in the formation of aggregates. Moreover, the proper amount of PVP can effectively tailor the morphology of Ni-MOFs, as shown in Fig. S2, which is gradually transformed from irregular aggregation to microspheres with an increase in the amount PVP.

To illustrate the growth mechanism of yolk–shell Ni-MOFs, the reaction-time dependence of microstructural evolution was investigated. Fig. 2(a) and (b) shows that only solid microspheres can be obtained at a reaction time of 2 h. It is obvious that the yolk–shell microspheres were gradually formed and the surfaces of microspheres became coarse and granular when the time was prolonged to 8 h (Fig. 2(c)–(f)). Fig. S3 shows that the samples have similar XRD patterns due to the same compositions. Based on the previous reports,33–35 the formation of Ni-MOFs could be illustrated by the Ostwald ripening mechanism involving core dissolution and shell re-deposition, as shown in Fig. 2(g). In the beginning stage, Ni2+ ions can coordinate with C9H6O6 ligands to produce primary nanoparticles, which are preferable to self-assemble into loose solid microspheres for the reduction of Gibbs free energy of the system in the presence of surfactant PVP. Considering the temperature gradient along the radial direction of microspheres in the actual reaction system, the surface temperature of the solid microspheres is higher than that in the interior region, which induces the high activity and dissolution of surface nanoparticles. Until the crystalline size is larger than the critical size in the surface domain, crystallization and re-deposition of nanoparticles occur and gradually grow to form the outside shell. With the prolonged reaction time, crystallization and re-deposition processes of surface nanoparticles proceed in turn and yield the cavity between the core and the shell consuming interior microspheres through the “inside to outside” transformation process. Thus, the yolk–shell structures were formed. With the prolonged reaction time of 18 and 24 h, the core gradually disappeared and the hollow structures were finally formed, which further confirm the flexibility of Ostwald ripening mechanism (Fig. S4).


image file: d0ce01242d-f2.tif
Fig. 2 SEM and TEM images of Ni-MOFs synthesized at different reaction times: (a) and (b) 2 h, (c) and (d) 5 h, (e) and (f) 8 h, and (g) illustration of the evolution process of the yolk–shell structure.

Structure properties of Ni/C composites

The XRD patterns of the Ni/C composites shown in Fig. 3 present the diffraction peaks at 44.4°, 51.7° and 76.2°, respectively, which could be indexed to the (111), (200) and (220) crystal planes of face-centered cubic (FCC) Ni according to the Joint Committee on Powder Diffraction Standards (JCPDS) card (No. 89-7128, space group Fm[3 with combining macron]m(225)). The inset figure shows that the (111) diffraction peak for the NC1, NC2, and NC3 shifts gradually to lower 2θ angles with the increase in the heat-treatment temperature. The corresponding planar spacing distance d of Ni nanoparticles slightly increases from 2.03 Å for NC1 to 2.04 Å for NC3, which may be attributed to the improved crystallization of Ni nanoparticles at a relatively high pyrolysis temperature. The average crystalline size estimated using the Scherrer formula is ca. 15 nm for NC1, 17 nm for NC2 and 28 nm for NC3, revealing that the high pyrolysis temperature accelerates the Ni atom diffusion and Ni nanoparticle growth. The absence of carbon diffraction peaks could be attributed to a small amount and/or poor crystallization (Fig. 3).
image file: d0ce01242d-f3.tif
Fig. 3 XRD patterns of Ni/C composites.

The SEM and TEM images shown in Fig. 4 and S4 exhibit the morphology and microstructure of Ni/C composites. It is obvious that NC1 and NC2 obtained at 500 °C and 600 °C well inherit the morphology of Ni-MOFs and exhibit a yolk–shell feature (Fig. S5(a) and (b) and 4(a) and (b)). The higher pyrolysis temperature of 700 °C causes structural breakage and collapse and yields numerous Ni@C nanoparticles (Fig. S4(c) and (d)). Compared with Ni-MOFs, the particle size of yolk–shell Ni/C microspheres slightly decreases and the cavity between core and shell is obviously magnified. Elemental mapping analysis in Fig. 4(c)–(e) confirms the existence of Ni and C elements with homogeneous distribution.


image file: d0ce01242d-f4.tif
Fig. 4 (a) SEM image, (b) TEM image and (c) STEM Z-contrast image of NC2. Elemental mapping of (d) C and (e) Ni. Inset shows a SEM image of a broken microsphere.

The HRTEM image shown in Fig. 5a reveals that the yolk–shell Ni/C microspheres are constructed with core–shelled Ni@C nanoparticles with an average particle size of ca. 25 nm. Fig. 5(b) shows the Ni core characterized by a lattice distance of 0.2 nm corresponding to the (111) lattice plane of Ni, while the carbon shell is featured by onion-like graphite with a thickness of ca. 3 nm. The graphitization degree of carbon shell is characterized by the Raman spectrum. Fig. S6 shows two peaks around 1342 and 1587 cm−1 corresponding to the D and G bands of carbon, respectively. Generally, the D band is related to defects or disorder of the graphite structure and the G band is a feature of perfect graphite.36,37 The value of ID/IG is 1.12 for NC1, 1.09 for NC2 and 1.01 for NC, demonstrating the reduction in the disorder state in the carbon shells with the increase in pyrolysis temperature. The N2 adsorption–desorption isotherms and pore size distributions of the samples were measured to characterize the mesoporous features. The IV-type isotherms with a hysteresis loop at a relative pressure ranging from 0.5 to 1.0 can be observed for all of the annealed samples (Fig. S7), indicating a mesoporous nature according to the International Union of Pure and Applied Chemistry (IUPAC) classification. The BJH desorption average pore size and the BET special surface area are 15.5 nm and 43.4 m2 g−1 for NC1, 13.8 nm and 38.8 m2 g−1 for NC2, and 14.4 nm and 55.1 m2 g−1 for NC3, respectively. The mesoporous feature should be attributed to the shrinkage of yolk–shell frameworks for the release of organic species during the pyrolysis process.


image file: d0ce01242d-f5.tif
Fig. 5 HRTEM images of (a) Ni@C core–shell nanoparticles and (b) the Ni core.

Magnetic and electromagnetic properties of Ni/C composites

Magnetic hysteresis loops of the Ni/C microspheres were investigated using a VSM at room temperature. Fig. 6 shows that the saturation magnetization (Ms) slightly increases with the heating temperature from 45.8 emu g−1 for NC1 and 47.1 emu g−1 for NC2 to 49.0 emu g−1 for NC3. The Ms value is slightly lower than that of 55 emu g−1 for bulk Ni possibly due to the presence of the non-magnetic carbon shell and/or spin disorder on the surface of Ni nanoparticles.38–40
image file: d0ce01242d-f6.tif
Fig. 6 Magnetic hysteresis loops of Ni/C composites.

According to the electromagnetic theory, the microwave absorption performance is highly dependent on the relative complex permittivity (εr = ε′ − ′′), the relative complex permeability (μr = μ′ − ′′) and the impedance matching. The reflection loss (RL) can be calculated from the complex permittivity and permeability at a given frequency and thickness by the transmit-line theory expressed as follows:24,27

 
RL = 20[thin space (1/6-em)] log|(ZinZ0)/(Zin + Z0)| (1)
 
Zin = Z0(μr/εr)1/2 tanh[j(2πfd/c)(μrεr)1/2] (2)
where Z0 is the impedance of free space, Zin the input impedance of the microwave absorbent, f the microwave frequency, d the thickness of the composites and c the velocity of light. An RL value of −10 dB corresponds to 90% microwave absorption. The real parts (ε′ and μ′) and imaginary parts (ε′′ and μ′′) represent the storage and loss ability of electromagnetic energy, respectively. In Fig. 7(a) and (b), ε′ and ε′′ of all the samples show a gradual decline tendency with an obvious fluctuation in the whole frequency range, revealing multiple dielectric resonance behaviors. The εr values are proportional to the pyrolysis temperature. The ε′ and ε′′ values are 12.8–7.2 and 4.1–0.63 for NC1, 19.0–8.8 and 12.4–1.3 for NC2, and 23.4–12.4 and 13.5–3.8 for NC3, respectively. In general, the Debye dipolar relaxation process can be utilized to illustrate the dielectric loss mechanism, which is described using the following equation:41,42
 
image file: d0ce01242d-t1.tif(3)
where εs and ε represent the static permittivity and the permittivity at a high frequency limit, respectively. According to this equation, the curve of εversus ε′′ can be acquired and each Cole–Cole semicircle is considered as a Debye relaxation process. As shown in Fig. 8(a), three Cole–Cole semicircles can be observed in the samples, representing the contribution of Debye relaxation processes of interfacial polarization among the Ni–C and C–paraffin interfaces and permanent dipolar polarization in the defective or disordered domain in carbon.28,42 The enhancement of ε′ value with the increase in heating temperature could be attributed to the improvements in Ni crystallinity and carbon graphitization degree, enhancing the dielectric polarization ability. In the case of ε′′, the enhancement of ε′′ may be ascribed to the improvement in the electronic conductivity of Ni/C composites according to the free electron theory.41,42 Additionally, in view of microstructural variations, the Ni@C nanoparticles among NC3 can act as a bridge to connect the Ni/C microspheres, which is apt to form the conductive network and further facilitate the electron transport. Dielectric loss tangents (tan[thin space (1/6-em)]δε = ε′′/ε′), as shown in Fig. 8(c), show that the tan[thin space (1/6-em)]δε values of NC2 and NC3 are higher than those of NC1 in the whole frequency range, while NC3 exhibits more excellent dielectric loss ability in the high frequency range of 11–18 GHz.


image file: d0ce01242d-f7.tif
Fig. 7 Frequency dependence of the electromagnetic parameters of Ni/C composites: (a) real part and (b) imaginary part of complex permittivity; (c) real part and (d) imaginary part of complex permeability.

image file: d0ce01242d-f8.tif
Fig. 8 (a) Cole–Cole curve, (b) frequency dependence of μ′′(μ′)−2f−1, (c) dielectric loss tangent and (d) magnetic loss tangent of Ni/C composites.

Fig. 7(c) and (d) clearly presents that the μrf curves of the NI/C microspheres are similar due to their little varied magnetic properties (Fig. 6). There exist two obvious peaks at ca. 3.3 and 12.0 GHz in the μ′′–f curves, indicating multiple magnetic resonance behaviors. It is well known that the magnetic loss are generally originated from magnetic hysteresis, domain wall displacement, eddy current loss, natural resonance and exchange resonance.22,24,31 At a weak EM field, the magnetic hysteresis loss can be ignored. Domain wall displacement does not occur because the average particle size of Ni nanoparticles is ca. 25 nm, which is much smaller than the single domain size of 55 nm.43 If the magnetic loss is mainly induced from the eddy current loss, the value of μ′′(μ′)−2(f)−1 is independent of frequency and equals to a constant value of 2πμ0d2δ/3, where d is the thickness of the absorbent, δ the electrical conductivity, and μ0 the permeability of vacuum. However, the value of μ′′(μ′)−2(f)−1 in Fig. 8(b) shows a changeable tendency over the whole frequency range, indicating the absence of the eddy current loss. According to the Kittel natural resonance,44 the resonance frequency of spherical magnetic materials can be expressed as fr = γ|Heff|, where γ = 28 GHz T−1 is the gyromagnetic ratio and Heff = 2|Keff|/Ms is the effective anisotropy field of a cubic crystal with Ms the saturation magnetization and Keff the effective anisotropy constant. Taking account of the Ms of fcc-type bulk Ni (ca. 490 emu cm−3) and the magnetocrystalline anisotropy constant (about −5 × 104 erg cm−3),43 the fr value is estimated to be 0.57 GHz. Compared with bulk Ni, the effective anisotropy field of Ni spherical nanoparticles is highly affected by surface anisotropy for Ni nanoparticles and may significantly enhance their resonance frequency to 4.3 GHz for 13 nm Ni nanaoparticles45 and 5.5 GHz for Ni@C nanoparticles with the size range of 25–30 nm.46 Thus, the experimental resonance peak at 3.3 GHz is probably attributed to the natural resonance, while the resonance peak at around 12.0 GHz is ascribed to the exchange resonance originating from magnetic dipole interaction.47,48 Fig. 8(d) shows the dependence of magnetic loss tangents (tan[thin space (1/6-em)]δμ = μ′′/μ′) on frequency. Both the natural resonance and the exchange resonance dominate tan[thin space (1/6-em)]δμ.

The microwave absorption performance of Ni/C composites is shown in Fig. 9. Fig. 9(a)–(c) presents that the absorption bandwidth with an RL < −10 dB is 6.1 GHz (8.7–14.8 GHz) for NC1, 10.0 GHz (6.7–16.7 GHz) for NC2 and 5.1 GHz (8.6–13.7 GHz) for NC3 in the thickness range of 1–3 mm, respectively. The NC2 exhibits excellent broad-frequency absorption and the best RL value is −39 dB with a bandwidth of 3.8 GHz at a less thickness of 1.8 mm. Moreover, the integrated bandwidth of NC2 can achieve 12.3 GHz covering Ku-band (12–18 GHz), X-band (8–12), and most of C-band (5.7–8 GHz) with a changeable thickness of 1.4–3.9 mm, as shown in Fig. 9(d). Table S1 lists the microwave absorption performance of some reported absorbers. In comparison, the NC2 exhibits competitive EM absorption performance with the lightweight, less thickness and wide-frequency band features.


image file: d0ce01242d-f9.tif
Fig. 9 Microwave absorption properties of Ni/C composites: (a) NC1, (b) NC2, and (c) NC3. (d) Three-dimensional plot of RL for NC2.

It is well known that EM absorbents with excellent microwave absorption performance should simultaneously satisfy the impedance matching and microwave attenuation. The impedance matching degree can be elevated by Z = |Zin/Zo| based on the formula (2), while the microwave attenuation ability can be calculated using the following formula:7,12,45

 
image file: d0ce01242d-t2.tif(4)
The Z value close to 1 infers that most of the incident EM wave can transmit into microwave absorbents. Fig. 10(a)–(c) plot the dependence of Z value as a function of frequency. The Z peak gradually shifts to a low frequency with the increase in the layer thickness of the samples. The value of Z peak is at 1–1.25 in 9–12 GHz for NC1, at 0.8–1.2 in 6–18 GHz for NC2 and at below 0.6 in 9–13 GHz for NC3, which may be the reason for the optimal RL peak being lower than −10 dB in the corresponding frequency range (Fig. 9). It should be noted that NC2 has the widest frequency band with the optimal Z value close to 1. The α value (Fig. 10(d)) gradually increases in the order of NC1, NC2 and NC3. It demonstrates that the increase in heating temperature can effectively regulate the Z and α values. NC3 treated at a relatively high temperature of 700 °C results in the largest εr values and the corresponding largest α value. However, the Z value is below 0.6 far from 1.0 over the whole frequency, resulting in the relative poor microwave absorption. However, although NC1 treated at 500 °C has the smallest α value due to the smallest εr values, the Z value around 1–1.25 superior to that of NC3 induces the enhancement of corresponding microwave absorption. This indicated that high temperatures are favorable for microwave attenuation and low temperatures are beneficial for impedance matching. An appropriate heating temperature of 600 °C can induce moderate εr values and contribute to the optimized balance between Z and α. Thus, NC2 shows excellent microwave absorption.


image file: d0ce01242d-f10.tif
Fig. 10 Impedance matching of NC1 (a), NC2 (b) and NC3 (c) and their attenuation coefficient (d).

In our opinion, the enhanced heating temperature of the Ni-MOF can significantly tune the complex permittivity of the Ni/C composites discussed above, but seldomly affect μr of the composites because of the same ferromagnetic metallic Ni component. In order to illustrate the contribution of Ni nanoparticles to the microwave absorption performance, the metallic Ni in NC2 was selectively etched by HCl treatment under a hydrothermal process at 220 °C for 10 h. The XRD pattern in Fig. S8(a) demonstrates that the acid etching treatment can effectively decrease the Ni content, where the etched sample exhibits weak XRD peaks corresponding to Ni and clear XRD peaks for (002) graphite located at ca. 26°. Similar to the morphology of NC2, the etched sample shows a homogeneous spherical morphology with a yolk–shell configuration (Fig. S8(b)). Compared with NC2, the etched sample exhibits larger εr and several magnetic resonances (Fig. S9). The Z value becomes less than 0.4 and the microwave absorption becomes poor although the α value is enhanced (Fig. S10). Due to the similar morphology between NC2 and the etched NC2 samples, the decrease in microwave absorption performance should be ascribed to the reduction of ferromagnetic Ni nanoparticles in the composite, which yield poor impedance matching. On the basis of the above analysis, the microwave absorption mechanism of yolk–shell Ni/C microspheres can be proposed and illustrated in Scheme 2. Multiple interfacial polarization, dipolar polarization, conduction loss and magnetic loss from the cooperative effect of Ni and C may result in optimized impedance matching and microwave attenuation of the Ni/C composites. Meanwhile, the unique yolk–shell configuration can facilitate multiple-reflection and scattering of the incident EM wave in the interior cavity. The combination of dielectric/magnetic loss and geometrical effect endows the yolk–shell Ni/C microspheres with excellent microwave absorption performance.


image file: d0ce01242d-s2.tif
Scheme 2 Schematic of the absorption mechanism of yolk–shell Ni/C microspheres.

Conclusions

In summary, yolk–shell Ni/C microspheres composed of Ni@C nanoparticles have been successfully prepared by a simple in situ carbon-thermal reduction process of yolk–shell Ni-MOFs at 500 and 600 °C under N2 protection. The improvement of the crystallinity of yolk–shell Ni/C microspheres by increasing the pyrolysis temperature is beneficial for the enhancement of the complex permittivity. The yolk–shell Ni/C microspheres obtained at 600 °C show the best reflection loss (RL) of −39 dB with a bandwidth of 3.8 GHz at a less thickness of 1.8 mm. The excellent microwave absorption performance is attributed to the optimized balance of impedance matching and microwave attenuation derived from the cooperative effect of magnetic Ni and dielectric C with the unique yolk–shell microstructure. The resultant yolk–shell Ni/C microspheres are attractive candidates in microwave absorption fields.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work has been supported by National Natural Science Foundation of China (51601120 and 51971221), plan for promoting innovative talents of Education Department of Liaoning Province (LCR2018015), and Shenyang Youth Science and Technology Project (RC200444).

References

  1. X. Liang, Z. Man, B. Quan, J. Zheng, W. Gu, Z. Zhang and G. Ji, Nano-Micro Lett., 2020, 12, 102 CrossRef CAS.
  2. L. Wang, X. Yu, X. Li, J. Zhang, M. Wang and R. Che, Chem. Eng. J., 2020, 383, 123099–123109 CrossRef CAS.
  3. A. Iqbal, F. Shahzad, K. Hantanasirisakul, M. K. Kim, J. Kwon, J. Hong, H. Kim, D. Kim, Y. Gogotsi and C. M. Koo, Science, 2020, 369, 446–450 CAS.
  4. Y. Cheng, J. Zhu, Y. Seow, H. Zhao, Z. J. Xu and G. Ji, Nano-Micro Lett., 2020, 12, 125 CrossRef.
  5. V. G. Andreev, S. B. Menshova, A. N. Klimov, R. M. Vergazov, S. B. Bibikov and M. V. Prokofifiev, J. Magn. Magn. Mater., 2015, 394, 1–6 CrossRef CAS.
  6. X. Zhang, Y. Li, R. Liu, Y. Rao, H. Rong and G. Qin, ACS Appl. Mater. Interfaces, 2016, 8, 3494–3498 CrossRef CAS.
  7. X. Wang, G. Shi, F. N. Shi, G. Xu, Y. Qi, D. Li, Z. Zhang, Y. Zhang and H. You, RSC Adv., 2016, 6, 40844–40853 RSC.
  8. T. Xia, C. Zhang, N. A. Oyler and X. Chen, Adv. Mater., 2013, 25, 6905–6910 CrossRef CAS.
  9. M. Q. Ning, M. M. Lu, J. B. Li, Z. Chen, Y. K. Dou, C. Z. Wang, F. Rehman, M. S. Cao and H. B. Jin, Nanoscale, 2015, 7, 15734–15740 RSC.
  10. X. Liang, B. Quan, Z. Man, B. Cao, N. Li, C. Wang, G. Ji and T. Yu, ACS Appl. Mater. Interfaces, 2019, 11(33), 30228–30233 CrossRef CAS.
  11. B. Dai, B. Zhao, X. Xie, T. Su, B. Fan, R. Zhang and R. Yang, J. Mater. Chem. C, 2018, 6, 5690–5697 RSC.
  12. X. Wang, Q. Geng, G. Shi, G. Xu, J. Yu, Y. Guan, Y. Zhang and D. Li, J. Alloys Compd., 2019, 803, 818–825 CrossRef CAS.
  13. X. Bao, X. Wang, X. Zhou, G. Shi, G. Xu, J. Yu, Y. Guan, Y. Zhang, D. Li and C. Choi, J. Alloys Compd., 2018, 769, 512–520 CrossRef CAS.
  14. Y. L. Ren, H. Y. Wu, M. M. Lu, Y. J. Chen, C. L. Zhu, P. Gao, M. S. Cao, C. Y. Li and Q. Y. Ouyang, ACS Appl. Mater. Interfaces, 2012, 4, 6436–6442 CrossRef CAS.
  15. J. Jiang, D. Li, D. Geng, J. An, J. He, W. Liu and Z. Zhang, Nanoscale, 2014, 6, 3967–3971 RSC.
  16. R. Qiang, Y. Du, Y. Wang, N. Wang, C. Tian, J. Ma, P. Xu and X. Han, Carbon, 2016, 98, 599–606 CrossRef CAS.
  17. Y. N. Ko, S. H. Choi, S. B. Park and Y. C. Kang, Nanoscale, 2014, 6, 10511–10515 RSC.
  18. J. Liu, W. You, J. Yu, X. Liu, X. Zhang, J. Guo and R. Che, ACS Appl. Nano Mater., 2019, 2, 910–916 CrossRef CAS.
  19. Q. Liu, Q. Cao, H. Bi, C. Liang, K. Yuan, W. She, Y. Yang and R. Che, Adv. Mater., 2016, 28, 486–490 CrossRef CAS.
  20. B. Zhao, X. Guo, W. Zhao, J. Deng, G. Shao, B. Fan, Z. Bai and R. Zhang, ACS Appl. Mater. Interfaces, 2016, 8, 28917–28925 CrossRef CAS.
  21. J. Feng, Y. Wang, Y. Hou and L. Li, Inorg. Chem. Front., 2017, 4, 935–945 RSC.
  22. H. Li, S. Bao, Y. Li, Y. Huang, J. Chen, H. Zhao, Z. Jiang, Q. Kuang and Z. Xie, ACS Appl. Mater. Interfaces, 2018, 10, 28839–28849 CrossRef CAS.
  23. J. Xu, J. Liu, R. Che, C. Liang, M. Cao, Y. Li and Z. Liu, Nanoscale, 2014, 6, 5782–5790 RSC.
  24. X. Jian, X. Xiao, L. Deng, W. Tian, X. Wang, N. Mahmood and S. Dou, ACS Appl. Mater. Interfaces, 2018, 10, 9369–9378 CrossRef CAS.
  25. J. Feng, Y. Hou, Y. Wang and L. Li, ACS Appl. Mater. Interfaces, 2017, 9, 14103–14111 CrossRef CAS.
  26. W. Liu, S. Tan, Z. Yang and G. Ji, ACS Appl. Mater. Interfaces, 2018, 10, 31610–31622 CrossRef CAS.
  27. W. Liu, L. Liu, Z. Yang, J. Xu, Y. Hou and G. Ji, ACS Appl. Mater. Interfaces, 2018, 10, 8965–8975 CrossRef CAS.
  28. Y. Lü, Y. Wang, H. Li, Y. Lin, Z. Jiang, Z. Xie, Q. Kuang and L. Zheng, ACS Appl. Mater. Interfaces, 2015, 7, 13604–13611 CrossRef.
  29. R. Qiang, Y. Du, H. Zhao, Y. Wang, C. Tian, Z. Li, X. Han and P. Xu, J. Mater. Chem. A, 2015, 3, 13426–13434 RSC.
  30. Y. Qiu, Y. Lin, H. Yang, L. Wang, M. Wang and B. Wen, Chem. Eng. J., 2020, 383, 123207 CrossRef CAS.
  31. W. Liu, Q. Shao, G. Ji, X. Liang, Y. Cheng, B. Quan and Y. Du, Chem. Eng. J., 2017, 313, 734–744 CrossRef CAS.
  32. F. Zou, Y. M. Chen, K. W. Liu, Z. T. Yu, W. F. Liang, S. M. Bhaway, M. Gao and Y. Zhu, ACS Nano, 2016, 10, 377–386 CrossRef CAS.
  33. L. Shen, L. Yu, X. Y. Yu, X. Zhang and X. W. Lou, Angew. Chem., 2014, 126, 1–6 CrossRef.
  34. Z. Q. Li, W. C. Chen, F. L. Guo, L. E. Mo, L. H. Hu and S. Y. Dai, Sci. Rep., 2015, 5, 14178 CrossRef CAS.
  35. J. Liu, Y. Zhou, J. Wang, Y. Pana and D. Xue, Chem. Commun., 2011, 47, 10380–10382 RSC.
  36. X. Wang, X. Bao, X. Zhou and G. Shi, J. Alloys Compd., 2018, 764, 701–708 CrossRef CAS.
  37. X. Li, D. Du, C. Wang, H. Wang and Z. Xu, J. Mater. Chem. C, 2018, 6, 558–567 RSC.
  38. S. H. Wu and D. H. Chen, J. Colloid Interface Sci., 2003, 259, 282–286 CrossRef CAS.
  39. Z. Liu, S. Li, Y. Yang, S. Peng, Z. Hu and Y. Qian, Adv. Mater., 2003, 15, 1946–1948 CrossRef CAS.
  40. G. Tong, Q. Hu, W. Wu, W. Li, H. Qian and Y. Liang, J. Mater. Chem., 2012, 22, 17494–17504 RSC.
  41. Y. Cheng, W. Meng, Z. Y. Li, H. Q. Zhao, J. M. Cao, Y. W. Du and G. B. Ji, J. Mater. Chem. C, 2017, 5, 8981–8987 RSC.
  42. B. Quan, X. Liang, G. Ji, Y. Cheng, W. Liu, J. Ma, Y. Zhang, D. Li and G. Xu, J. Alloys Compd., 2017, 728, 1065–1075 CrossRef CAS.
  43. D. L. Leslie-Pelecky and R. D. Rieke, Chem. Mater., 1996, 8, 1770–1783 CrossRef CAS.
  44. J. Zou, Z. Wang, M. Yan and H. Bi, J. Phys. D: Appl. Phys., 2014, 47, 275001 CrossRef.
  45. G. Wang, X. Peng, L. Yu, G. Wan, S. Lin and Y. Qin, J. Mater. Chem. A, 2015, 3, 2734–2740 RSC.
  46. X. F. Zhang, X. L. Dong, H. Huang, Y. Y. Liu, W. N. Wang, X. G. Zhu, B. Lv, J. P. Lei and C. G. Lee, Appl. Phys. Lett., 2006, 89, 053115 CrossRef.
  47. A. Aharoni, J. Appl. Phys., 1991, 69, 7762 CrossRef.
  48. P. Toneguzzo, G. Viau, O. Acher, F. Fiévet-Vincent and F. Fiévet, Adv. Mater., 1998, 10, 1032 CrossRef CAS.

Footnote

Electronic supplementary information (ESI) available. See DOI: 10.1039/d0ce01242d

This journal is © The Royal Society of Chemistry 2020