I. INTRODUCTION

Recently, electromagnetic (EM) shielding materials have attracted wide attention due to their ability to protect the environment and sensitive devices from EM interference (EMI).1–10 Compared with metal materials which cannot be used at high temperatures, ceramics are more suitable for EM shielding materials because of their excellent structural stability, oxidation resistance, creep resistance, and high-temperature mechanical properties.11 EM shielding materials should have mobile charge carriers that can interact with the EM radiations.12 They should contain carbon materials and nano-sized SiC, which have high electrical conductivity.13–16 Therefore, many researches focus on EMI shielding effectiveness (SE) of Si–C–N based ceramics (Si3N4, SiC, C, SiO2) and composites.17 Usually, EMI SE is used to evaluate the capability of materials to shield the EM wave. As Eq. (1) shows, the total loss of the EMI shielding (SET) is composed of the reflection loss (SER), the absorption loss (SEA), and multiple-reflections SE loss (SEM). In addition, SEM is negligible only if SET > 15 dB,

$${\rm{S}}{{\rm{E}}_{\rm{T}}} = {\rm{S}}{{\rm{E}}_{\rm{R}}} + {\rm{S}}{{\rm{E}}_{\rm{A}}} + {\rm{S}}{{\rm{E}}_{\rm{M}}}\quad .$$
(1)

Table I summarizes the SE of several typical Si–C–N based ceramics. It indicates that ceramics which contain insulating components and conductive phase are promising materials to block the EM radiation and eliminate the EMI microwave. Their conductivity and EMI SE can be easily adjusted and controlled by chemical approaches. Among these approaches, polymer infiltration and pyrolysis (PIP) is an efficient way to improve the electrical conductivity of ceramics by increasing the infiltration cycle.

TABLE I The comparison of EM shielding properties of different ceramics.
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Silicon based polymer-derived ceramics (PDCs) have obtained more attention because of their characterization that contains carbon materials and nano-sized SiC after heat-treatment.23 Especially, SiC nanowires (NWs) and free carbons can be synthesized by pyrolysis of polysiloxane (PSO).22 Porous silicon nitride (Si3N4) is selected as a matrix for its excellent mechanical properties, good oxidation resistance, and thermal-chemical corrosion resistance, which can be used in a harsh environment. Adding SiC NWs and free carbons into porous Si3N4 by PIP is an efficient way to prepare EMI shielding materials. Therefore, fabricating Si3N4–SiOC ceramics by PIP can create SiC NWs and free carbons in the pores to reduce the porosity and increase the density of the matrix, eventually making the Si3N4–SiOC ceramics with excellent EM shielding and high mechanical properties.

Si3N4 can be fabricated by hot pressing and cold pressing followed by reactive sintering. All these technologies can only produce components with simple geometries. Compared with prior technologies, three-dimensional printing (3DP) is one special method of additive manufacturing. It can fabricate complex-shaped components without wasting materials. According to our previous researches, Si3N4 ceramics fabricated by 3DP and pressureless-sinter have an open porosity of 78%.24 In order to further improve the mechanical properties of porous Si3N4 ceramics, direct nitridation of silicon is a promising way for the synthesis of Si3N4 ceramics. Compared with the previous work, the porosity of Si3N4 can be reduced by nearly 10% through direct nitridation of silicon.

In this paper, Si3N4–SiOC ceramics are synthesized by 3DP of silicon powders and direct nitridation of silicon, followed by PIP. It enables the complex-shape fabrication and microstructure designing and enhances EM shielding properties and mechanical properties of Si3N4–SiOC.

II. EXPERIMENT

The Si powder was mixed with Lu2O3 and dextrin powder with a mass ratio of 83:7:10. The mixed powder was ball-milled for 48 h in deionized water, and then dried at 80 °C for 24 h. The as-received powder blend was crushed and passed through a mesh sieve. The powder was printed into specimens with dimensions 3 × 10.5 × 24 mm3 by 3DP, which was carried out in a 3D-printer (Spectrum Z510, Z Corporation, Rock Hill, South Carolina) using a water-based printer solution. During printing, the layer thickness was set to 0.0875 mm. Then the as-prepared specimens were decarburized at a temperature of 700 °C for 360 min. After this, the specimens were nitrided at 1500 °C for 120 min in the N2 atmosphere in a tube furnace. The Si3N4 ceramics with a porosity of 70% were obtained.

PSO and 2 wt% ferrocene [Fe(C5H5)2] were mixed for 30 min at 70 °C in the Ar atmosphere, and then a transparent orange liquid was obtained. The porous Si3N4 specimens were dipped into the obtained orange liquid for 30 min by a pressure infiltration process each time. The polymer in porous Si3N4 ceramics was cross-linked at 300 °C for 2 h and annealed at 1350 °C with a speed of 5 °C/min for 2 h with Ar gas as the protection atmosphere. The above processes were repeated several times until the desired content was reached. The obtained samples were designated as PIP-0, PIP-1, PIP-2, PIP-3, and PIP-4, respectively, corresponding to the infiltration cycles of 0, 1, 2, 3, and 4.

X-ray diffraction (XRD) was used for phase analysis (X’Pert Pro, Philips, Almelo, The Netherlands). INVIN Raman spectroscopy (RMS; Renishaw, Gloucestershire, United Kingdom) was used to characterize the graphitization degree of samples. The microstructure was observed by scanning electron microscopy (SEM; S-4700, Hitachi Corp., Tokyo, Japan) and transmission electron microscopy (TEM; G-20, FEI-Tecnai, Hillsboro, Oregon; at 300 kV). The scattering parameters (S-parameters, S11 and S21) of each sample were measured by a vector network analyzer (VNA; MS4644A, Anritsu, Kanagawa, Japan) using the waveguide method in the X-band. The flexural strength was measured by a 3-point bending method on a universal testing machine (CMT 4304, Sans, Shenzhen, China) at a cross-head speed of 0.5 mm/min with a loading span of 30 mm.

III. RESULTS AND DISCUSSION

A. The microstructure of Si3N4–SiOC ceramics

Figure 1(a) shows XRD of Si3N4 (PIP-0) and Si3N4–SiOC. It is clearly indicated that the composites after the direct nitridation of silicon contain a high content of α-Si3N4, low content of β-Si3N4, and small amount of Si. The surface of the silicon nitride product is covered by long whiskers with a diameter in the micrometer range as shown in Fig. 2(a); it is indicated that the nitridation of silicon is related to the reaction in the gas phase. These whiskers are grown via a vaporization condensation type mechanism. Silicon vapor which generated from the surface reacted with nitrogen gas and deposited as Si3N4 whiskers.25

FIG. 1
figure 1

(a) XRD patterns and (b) Raman spectra of Si3N4–SiOC ceramics.

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FIG. 2
figure 2

SEM micrographs of the fracture surface of Si3N4–SiOC with different cycle numbers of PIP: (a) PIP-0, (b) PIP-1, (c) PIP-2, (d) PIP-3, and (e) PIP-4.

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Si3N4–SiOC ceramics were successfully synthesized after the PIP process. The density and porosity of porous Si3N4 were 0.87 g/cm3 and 70%, respectively. The density of porous Si3N4 after the nitridation of silicon was higher than that of Si3N4 prepared by 3DP and pressureless sintering of silicon nitride. After PIP, the weight gains of samples PIP-1 and PIP-2 were 33.4 wt% and 41.6%, the densities of samples PIP-1 and PIP-2 were 1.48 g/cm3 and 1.72 g/cm3, and porosities of samples PIP-1 and PIP-2 were 48 and 30%. As the infiltration cycle increased from three to four times, the weight gain of the porous Si3N4–SiC ceramic increased from 47.3 to 50.1 wt% with the density increasing gradually from 1.98 to 2.09 g/cm3 and the total porosity decreasing slowly from 23 to 21%.

Figure 1(a) shows the XRD patterns of Si3N4–SiOC with dependence on the cycle number of PIP. The original porous Si3N4 ceramic was composed of α-Si3N4 and β-Si3N4. After infiltration, the crystalline peaks at 2θ of 23.4°, 27.1°, 33.7°, and 36.1° increased, referring to the diffraction from the (110), (200), (101), and (210) lattice planes of β-Si3N4, which suggested that some α-Si3N4 were transformed into β-Si3N4 during the process of PIP. In the meantime, all the samples showed the crystalline peaks at 35.6° and 60.0° after infiltration, which referred to the (111) and (220) lattice planes of β-SiC. With the increase of the infiltration cycle, the peak intensity of β-SiC increased gradually while the peak intensity of Si3N4 decreased, indicating that the content of SiC increased.

The Raman spectra of the samples are shown in Fig. 1(b). All the spectra of the samples showed the representative D band and G band, suggesting the existence of free carbons. The intensity ratio of the D band and G band, ID/IG, could be used to evaluate the carbon cluster size. The ID/IG ratio of all samples after PIP, as shown in Fig. 1(b), had no remarkable change. The content of free carbons increased when the cycle number of PIP increased. The 2D band at about 2700 cm−1 is the overtone of the D band, which was sensitive to the stacking orders of graphene layers along the c axis. Therefore, the existence of the 2D band indicated the high crystallization of carbons.

Figure 2 shows the SEM micrographs of the fracture surface of Si3N4–SiOC ceramics with different cycle numbers of PIP. As shown in Fig. 2(b), SiC NWs with a diameter smaller than 100 nm were formed after the PIP process. With the increase of the cycle number of PIP, more SiC NWs were introduced into the porous ceramics. The porosity, as well as the size, of pores decreased.

Figure 3 shows the TEM images of the sample PIP-3. The interplanar spacing was 0.67 nm as shown in Fig. 3(a), corresponding to the (100) plane of Si3N4, which indicated that Si3N4 whiskers had already formed before the PIP process. There was some turbostratic carbon in the ceramics as shown in Fig. 3(b). The TEM image [Fig. 3(c)] showed the existence of SiC NWs. The SiC NWs had a zigzag surface morphology. The selected area electron diffraction (SAED) pattern [Fig. 3(d)] showed three diffraction rings corresponding to the typical lattice planes of 3C–SiC. Besides diffraction rings, the diffraction spots gave two symmetrical sets of face-centered cubic lattices, suggesting that the zigzag part was composed of (111) twinning structures. The structures of twinned SiC NWs were similar to that in previous works.26,27

FIG. 3
figure 3

TEM images of Si3N4–SiOC ceramics: (a) Si3N4, (b) turbostratic carbon, (c) twinned SiC nanowires, and (d) corresponding SAED.

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B. EMI shielding properties of Si3N4–SiOC ceramics

Figure 4 presents direct-current electrical conductivity (σdc) and EMI properties of Si3N4–SiOC ceramics. σdc of porous Si3N4 was 7.6 × 10−6 S/m. After PIP, the sample PIP-1 showed σdc of 3.6 × 10−1 S/m, indicating the formation of SiC NWs and turbostratic carbon played an important role in electrical conductivity. SiC NWs and carbon had high carrier concentration close to the metallic.28 Therefore, the bridging and connectivity of SiC NWs in the pores and the existence of turbostratic carbon could dramatically improve the electrical conductivity of Si3N4–SiOC ceramics. σdc of Si3N4–SiOC increased from 3.6 × 10−1 S/m to 45.3 S/m with the cycle number of PIP increasing from 1 to 4, which was attributed to the increase of the content of SiC NWs and turbostratic carbon.

FIG. 4
figure 4

(a) Electrical conductivity, (b) SET, (c) SER, and (d) SEA of Si3N4–SiOC ceramics.

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As shown in Figs. 4(b)–4(d), SER decreased and SEA increased with the frequency going up, which conformed to the trend of Eqs. (2) and (3) correctly. SET, SEA, and SER of samples all increased with the increase of the cycle number of PIP. The average SET, SEA, and SER of the samples increased from 2 to 35 dB, 0.1 to 28 dB, and 1 to 7 dB, respectively. SEA and SER were closely associated with the frequency and the electrical conductivity according to Eqs. (2) and (3), respectively,29

$${\rm{S}}{{\rm{E}}_{\rm{R}}} = 39.5 + 10\>\log {\sigma \over {2\pi f\mu }}\quad ,$$
(2)
$${\rm{S}}{{\rm{E}}_{\rm{A}}} = 8.7d\sqrt {\pi f\mu \sigma } \quad ,$$
(3)

where f is the frequency, σ is the electric conductivity, μ is the permeability, and d is the specimen thickness.

The schematic illustration of the shielding mechanism in Si3N4–SiOC ceramics is shown in Fig. 5. The increase of SER could be attributed to the formation of turbostratic carbon which leads to a high reflection of incident. Twinned SiC NWs with stacking fault and numerous dipoles formed in the pores could result in extensive polarization and energy dissipation.16 SiC NWs possess a large surface area. A larger surface area of conductive material could give rise to interfacial polarization. Energy dissipation and interfacial polarization eventually lead to the increase of SEA. SiC and carbon as the semiconductor and conductor make a great contribution to the absorption loss. Materials with low SER and high SEA are candidates for EM absorption. The existence of Si3N4 and SiOC as microwave transparent materials could ensure the impedance match between free space and ceramics, which makes Si3N4–SiOC ceramics have a high SET of 35 dB with a lower SER of 7 dB and higher SEA of 28 dB. It is indicated that more than 99.9% of the EM energy was reflected and absorbed and only less than 0.1% of the EM energy was transmitted.

FIG. 5
figure 5

Schematic illustration of the shielding mechanism in Si3N4–SiOC ceramics.

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C. Mechanical properties of Si3N4–SiOC ceramics

The flexural strength of Si3N4–SiOC ceramics with different cycle numbers of PIP was measured. After the direct nitridation of silicon, porous Si3N4 showed a low flexural strength (13 MPa) due to its high porosity and high content of α-Si3N4. After the infiltration cycle, the porosity of the sample PIP-1 decreased which would promote the increase of the flexural strength. The flexural strength of the sample PIP-1 was 37 MPa. When the infiltration cycle number was four times, the flexural strength of the sample PIP-4 increased to 63 MPa. Therefore, the increase of the content of SiOC enhanced the flexural strength of Si3N4–SiOC ceramics to a certain degree.

Table II offers present achievements in the field of processing of near-net-shaped 3D ceramic-based parts with a required shape using the 3DP technology.24,30–32 3DP can be used to create geometries that are difficult or impossible to create by traditional means.33 The other advantage of 3DP is that there are no limitations on the material species as long as it can be formulated into a powder which has good flowability and low internal friction. The drawback of 3DP is that ceramic-based parts have considerable porosity and require postprocessing. Sintering and infiltration are two most common postprocessing steps as shown in Table II. These steps can decrease the porosity and increase the flexural strength of ceramics. Si3N4–SiOC ceramics fabricated by 3DP and direct nitridation of silicon combined with PIP not only have high flexural strength but also possess excellent EMI shielding properties.

TABLE II Ceramic materials prepared by 3DP and corresponding properties.
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Figure 6 presents the computer-aided design model and the corresponding printed part composed of Si3N4 ceramics after the direct nitridation of silicon. The silicon nitride product has a linear shrinkage of 3–5%. It enables us to form components with accurate geometry dimension. This approach can also be used to print lattice truss structures, 3D mesh structures, and even more complex shapes.34,35 What is more, it can also be further extended to develop other functional ceramic systems for electromagnetic wave (EMW) absorption and shielding applications.

FIG. 6
figure 6

(a) Computer-aided design model, (b) corresponding porous Si3N4 after the direct nitridation of silicon, and (c) Si3N4–SiOC ceramics with a complex shape.

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IV. CONCLUSION

In this study, PDC–SiOC which contained SiC NWs and free carbons was in situ formed in porous Si3N4 ceramics by the PIP process. With the increase of the infiltration cycle number, most of the pores were filled with SiC NWs and free carbons, which acted as the conductive phase and reinforced phase, resulting in the improvement of both shielding and mechanical properties. As the porosity went down from 70 to 21%, the average SET rose from 2 dB to 35 dB with a lower SER of 7 dB and higher SEA of 28 dB, and the flexural stress reached 63 Mpa. Therefore, it is demonstrated that the SiC NWs reinforced Si3N4–SiOC ceramics are a promising EM shielding and structural material. The method in this work can be further extended to prepare other functional ceramic systems with a complex shape.