Synergistic regulation of color and mechanical properties of silicon nitride ceramics via engineering hollow structures of Eu-enriched secondary phases

INTRODUCTION

The escalating demand for 5th generation mobile communication technology necessitates the development of enclosure materials for electronic products, including cell phones and smart wearable devices, to satisfy stringent properties standards^[1,2]. These include high thermal conductivity (k), with k > 20 W/(m·K), for effective equipment cooling; low dielectric loss (tanδ), with tanδ < 10^-3, for optimal signal transmission, high mechanical properties, with fracture toughness (K_IC) > 10 MPa·m^1/2 and high flexural strength (σ) > 700 MPa, to ensure sufficient damage tolerance. Such stringent requirements render plastic-based, metal-based and glass-based materials and traditional structure ceramics such as zirconia (ZrO₂) and alumina (Al₂O₃) unsuitable. Alternatively, silicon nitride (Si₃N₄) has emerged as a leading contender for use in 5G electronic products, due to its superior overall performance, as depicted in Supplementary Figure 1. However, the issue of single color (typically grey) coupled with control over mechanical properties remains a significant challenge for large-scale commercialization^[3,4]. Therefore, there is an increasing urgency to address both color regulation and improvement of mechanical performance.

In the realm of color regulation, rare earth metal ions (Re³⁺) exhibit the capacity to absorb light in a manner that is triggered by the visible light excitation of a 4f→4f electron transition^[5,6]. Notably, among these rare earth metal ions, Eu³⁺ ions have the ability to transmute ultraviolet (UV) radiation into a strong orange-red emission. They are also highly sensitive to their coordination environment due to the unique combination of intra-configurational 4f→4f transitions that are either magnetic or electric-dipole in nature^[7,8]. These attributes position Eu³⁺ as a promising candidate for use as a colorant in the creation of orange, red, or similar orange-red Si₃N₄ ceramics. However, conventional coloring strategies such as dissolving Eu³⁺ ions within lattices to color ceramics, which has been successfully employed in zirconia^[9-11], are not applicable to Si₃N₄ coloring. This is primarily due to the robust covalent nature of the Si-N bond in β-Si₃N₄^[12], coupled with the substantial radius difference between Si⁴⁺ (0.41 Å^[13]) and Eu³⁺ (1.06 Å^[14]). In this case, the design of Eu³⁺-doped second crystal phases or new local structures as chromophores is the alternative way. Besides, it is reported that the introduction of second phases or novel local structures can also act as a feasible and practical way for improving mechanical properties^[15-17].

The selection of sintering additives in the liquid phase sintering process of Si₃N₄ is critically important due to their significant influence on the formation of secondary phases and novel local structures. Compared to the conventional sintering additives [e.g., yttrium oxide-aluminum oxide (Y₂O₃-Al₂O₃)^[18], ytterbium oxide-aluminum oxide (Yb₂O₃-Al₂O₃)^[19], MgO^[20], yttrium aluminum garnet (YAG, Y₃Al₅O₁₂)^[21], etc.], we develop YAG-MgO as a sintering additive for its lower eutectic temperature (< 1,613 K^[21]) than YAG or MgO via the YAG-MgO-SiO₂-Si₃N₄ reaction. The lower eutectic temperature offers more time for α-Si₃N₄ dissolution in the eutectic liquid phase, and then will promote the growth of high aspect ratio β-Si₃N₄ grains^[22-24] and the formation of second crystal phases or novel local structures.

In this contribution, we propose using phase separation and crystallization in liquid phases supported by Eu₂O₃-YAG-MgO to engineer hollow structures for achieving the color change from yellow to orange-red, while ensuring excellent mechanical properties of Si₃N₄. The combined microstructural characterizations of scanning electron microscopy (SEM), high-resolution transmission electron microscopy (HRTEM), scanning transmission electron microscopy (STEM), energy dispersive X-ray spectroscopy (EDS), electron energy loss spectroscopy (EELS) and cathodoluminescence (CL) were utilized to elucidate the crystallographic characteristics and chemical composition of the hollow structures, including morphology, chemical composition, distribution and luminescent characteristic. Based on these characterizations, the formation mechanism of the hollow structure in Si₃N₄ ceramics, as well as its impact on their color and mechanical properties, was meticulously analyzed and discussed.

MATERIALS AND METHOD

RESULTS AND DISCUSSION

During the liquid-phase sintering process of Si₃N₄ ceramics, the assembly of Eu ions, attributed to their high field strength and larger radius, leads to the segregation of the glass phase^[31]. This phenomenon is depicted in Supplementary Figure 3. The liquid phase, enriched with Eu, encapsulates the pores, thereby forming a hollow structure at the grain boundary. Furthermore, the generation of these hollow structures within β grains transpires during the solution-reprecipitation process^[12], which can be divided into two stages: initially, the α to β phase transformation in the liquid phase is facilitated by the homogeneous mixing of sintering additives (YAG-MgO) and colorants (Eu₂O₃) between 1,450-1,800 °C; all transformations into β are completed within this stage. Subsequently, the second stage involves the solution of small β grains and their reprecipitation on larger β grains at elevated temperatures (> 1,800 °C), concurrently introducing hollow structures into the larger β grains [Figure 1A]. A more comprehensive depiction of the second step is illustrated in Figure 1B and C. Two or more adjacent, small β grains with identical orientations are susceptible to dissolution, merging, and growth into larger β grains at elevated temperatures. Concurrently, the gas and liquid phases within the grain boundary are compressed into the β grains at the onset of the merger, and then persistently expelled to the grain boundary as the merger progresses. However, when sintering or grain growth is impeded, a limited amount of gas and liquid phase remains in the larger β grains. The residual liquid phase primarily consists of heavier Eu-clusters due to their lower diffusion rate and follows the growth of the β grain lattice, subsequently exhibiting a hexagonal morphology akin to that of the β grain. In contrast, the residual gas is enveloped by the hexagonal liquid phase, forming a hollow structure. Figure 1D illustrates the coloring process for Si₃N₄ ceramics, wherein Eu ions within the hollow structures absorb blue-green light and emit red-orange light due to electronic transitions under photon excitation.

Figure 1. Schematic illustration. (A) The solution-reprecipitation process: firstly, α-Si₃N₄ powders, colorants and sintering additives are uniformly mixed; secondly, α-β phase transformation during liquid phase sintering; thirdly, the solution of small β grains and reprecipitation on large β grains, accompanying by the introduction of hollow structures into the large β grains. (B) The generation process of the hollow structures within grains. (C) Proofs for the formation process of the hollow structure in β grain under scanning transmission electron microscopy. (D) Light propagation path in Si₃N₄ ceramics (left) and the schematic illustration of the hollow structure (right).

Samples with various Eu₂O₃ contents (2, 4, 5, 6, 7, 8, and 9 wt%) were designated as SEu-2, SEu-4, SEu-5, SEu-6, SEu-7, SEu-8, and SEu-9, respectively. SEM revealed that only rod-shaped grains and grain boundary liquid phases with distinct contrasts could be observed in the samples SEu-4, SEu-5, and SEu-6 [Supplementary Figure 4], aligning with prior research^[32-34]. Interestingly, STEM detected the hollow structures ranging from ~5 to ~110 nm in both grains and grain boundaries, as depicted in Figure 2A-D and Supplementary Figures 5-7. Additionally, HRTEM analyses demonstrated that the hollow structure in β-grains exhibits a hexahedral morphology identical to that of the β-grains [Figure 2E-H and Supplementary Figure 7]. The corresponding Fast Fourier Transformation (FFT) patterns of region 1 [Figure 2E and Supplementary Figure 7G] and region 2 [Supplementary Figure 7H] suggest two hexagonal phases, but the lattice spacing, d_average = 0.648, calculated from the FFT1 pattern, aligns with the zone axis [001], which falls in that of β-Si₃N₄ phase^[35,36], whereas the lattice spacing, d_average = 0.654, calculated from FFT2 pattern, is slightly different from that of the β grain. These findings suggest that the hollow structure is epitaxially grown on original β grains, albeit with minor distortions to the original crystal structure. Notably, the corresponding FFT4 pattern of the hollow structures at grain boundaries displays a halo indicative of an amorphous nature [Figure 2H], thus preventing the detection of the hollow structure by XRD patterns [Supplementary Figure 8] due to its amorphous nature. The chemical composition of the hollow structures is illustrated in Figure 2I and Supplementary Figure 9. Elemental maps reveal that the core, characterized by its element-free and rounded morphology, was classified as a pore. Conversely, the shell, which demonstrated a pronounced contrast difference with the β grains due to its enrichment in Eu elements, was identified as an Eu-rich secondary phase.

Figure 2. The crystallographic feature and chemical information of the hollow structure. (A and B) STEM images of the sample SEu-5 (the orange and blue squares mark the grain section and grain boundary regions, respectively): The hollow structures were formed in the grains and the grain boundary phase; (C) STEM magnified image of the orange area; (D) STEM magnified image of the blue area; (E-H) HRTEM images corresponding to the region 1-4 (the inset shows the corresponding FFT pattern). (I) STEM-EDS elemental analysis of hollow structure in sample SEu-5.

The electron valence states and luminescence of Eu ions in Si₃N₄ ceramics were examined using XPS, STEM-EELS, and SEM-CL techniques, as depicted in Figure 3A-C and Supplementary Figure 10. The XPS analysis of the SEu-5 sample revealed that both Eu³⁺ and Eu²⁺ signals, identified as Eu³⁺:3d_5/2,3/2 and Eu²⁺:3d_5/2,3/2, respectively, were readily found [Supplementary Figure 10]. This suggests a reduction of some Eu³⁺ to Eu²⁺ in an N₂ atmosphere. The STEM-EELS results indicated that the grain boundary glass phase (region 1) and the hollow structure (region 2) predominantly consisted of EELS signal pairs at 1,135.8/1,164.0 eV, attributed to Eu^3+[37], but no EELS peaks were observed in the β grains (region 3). The SEM-CL spectra of the β-grain and grain boundary phases in the SEu-5 displayed a broad emission peak and four narrow emission lines [Figure 3C]. The broad emission peak at 530 nm corresponds to the 5d-4f transition of Eu²⁺. Additionally, several distinct lines at 593, 616, 656, and 697 nm are associated with the ⁵D₀→⁷F₁,⁵D₀→⁷F₂, ⁵D₀→⁷F₃ and ⁵D₀→⁷F₄ transitions of Eu^3+[38,39], respectively. These CL emission peaks likely arise from the Eu-rich hollow structure present in grain boundary and β-grain. Consequently, the Eu²⁺/Eu³⁺-doped hollow structures present within grain boundaries and β-grains serve as a primary chromophore in Si₃N₄ ceramics. Furthermore, the distribution density of these hollow structures could be regulated via control of Eu₂O₃ content for color regulation of Si₃N₄ ceramics. The distribution density is defined as the average number of hollow structures per unit area (nm^-2) in the context of STEM, where 70 regions were randomly selected in each sample at magnification photography for counting. As shown in Figure 3D-G, the distribution density of the hollow structures progressively increases with Eu₂O₃ content. A higher distribution density implies more chromophores, leading to increased absorption of blue-green light. Further details will be explored in Figure 4.

Figure 3. The valence states and luminescence of Eu ions. (A) STEM image of the sample SEu-5. (B) EELS spectra of the grain boundary phase (region 1), the hollow structure (region 2) and the β-Si₃N₄ grain (region 3) corresponding to the STEM image: The pairs of EELS signals at 1,135.8/1,164.0 eV attributed to the Eu³⁺. (C) SEM-CL spectra of the β grain (grey line) and the grain boundary phase (orange line): There is a broadband emission peak at 530 nm and four narrow emission lines at 593, 616, 656 and 697 nm. (D-F) The distribution density of hollow structures in the sample SEu-4 (D), SEu-5 (E), SEu-6 (F) and SEu-7 (G) (the X-axis is defined as the average number of hollow structures per unit area (nm^-2); the percentage on the Y-axis refers to the relative frequency distributed in various hollow structure density intervals; the green lines mean Gauss Amp of the distribution density of hollow structure): With the increases of Eu₂O₃ content, the distribution density of hollow structure increases.

Figure 4. Color of Si₃N₄ ceramics. (A) Digital photo of Si₃N₄ ceramics with varying Eu₂O₃ mass rations: The orange color deepens with increasing Eu₂O₃ content. (B) L*, a* and b* parameters of Eu-doped Si₃N₄ ceramics: As the Eu₂O₃ content increases, the value of L* decreases, while the values of a* and b* increase. (C) Emission spectra of all samples by monitoring 465 nm excitation: As the Eu₂O₃ content increases, the emission spectra tend to redshift but the intensity of light emission is weakened. (D) Reflectance spectra of Eu-doped samples (the inset shows the reflected light of the samples is located in the wavelength range from 556 to 602 nm). (E) Plot of (αhν)²vs. hν of Eu-doped Si₃N₄ ceramics (the insert shows the enlarged image of sample SEu-9). (F) Eg values for Si₃N₄ ceramics with different content of Eu₂O₃: An augmentation in the Eu₂O₃ content results in a reduction of Eg values, transitioning from 2.24 to 2.04.

The influence of hollow structures on the optical properties of Si₃N₄ ceramics was examined, as depicted in Figure 4. Yellow and orange Si₃N₄ ceramics were synthesized using varying Eu₂O₃ concentrations. As the Eu₂O₃ concentration escalated from 2 to 9 wt%, the color transitioned from yellow to orange-red, with a heightened chroma [Figure 4A]. Figure 4B presents that the L* value diminishes while a* and b* values rise in response to the diminished brightness and a color shift towards reddish-yellow. The emission spectra of all samples by monitoring 465 nm excitation are shown in Figure 4C, where the intense broadband emission peaks for samples SEu-2, SEu-4, SEu-5, SEu-6, SEu-7, SEu-8, and SEu-9 occurred at 587, 605, 609, 612, 622, 624 and 626 nm, respectively, corresponding to the 5d→4f transition of Eu^2+[40]. A relatively faint yet broad emission is discernible at 530 nm, which can also be attributed to Eu²⁺ ion emissions. A more pronounced narrow emission peak at 700 nm is evident later in the spectrum. The ⁵D₀→⁷F₄ transitions of Eu³⁺ ions account for this narrowband emission^[39]. Additionally, samples SEu-2, SEu-4, SEu-5, SEu-6, SEu-7, SEu-8, and SEu-9 absorb blue or green light and exhibit strong reflectance in the yellow or orange wavelengths at 587, 588, 589, 593, 599, 600, and 601 nm, respectively [Figure 4D]. The band gap of Eu-doped Si₃N₄ ceramics falls within the yellow and orange light region (Eg: 2.03-2.4), as depicted in Figure 4E, and an increase in Eu₂O₃ content, a decrease from 2.24 to 2.04 in Eg values correspondingly [Figure 4F]. In conjunction with the findings presented in Figure 3D-G and the subsequent analysis of the optical properties of Si₃N₄ ceramics varying in Eu₂O₃ content, it can be deduced that the orange hue of Si₃N₄ ceramics exhibits a positive correlation with the distribution density of hollow structures. Specifically, as the distribution density of these structures increases, there is a corresponding rise in chromophore content. This leads to an enhanced absorption of blue-green light and a consequent increase in reflection of yellow-red light, resulting in a deeper orange coloration of the samples. Consequently, the coloration of Si₃N₄ ceramics can be effectively modulated by manipulating the distribution density of the hollow structures.

The impact of the hollow structures on the mechanical properties of Si₃N₄ ceramics was examined, as depicted in Figure 5 and Supplementary Figure 11. The K_IC and σ initially increased, subsequently decreasing in line with the trend of relative density [Supplementary Figure 11]. The microstructures of the sample SEu-5, as revealed by SEM and TEM micrographs, exhibit elongated grains, grain pull-out, curved crack growth paths, and pronounced intergranular fracture [Supplementary Figures 12 and 13]. These attributes promote the utilization of crack energy and augment fracture resistance^[41-44]. Consequently, the sample SEu-5 demonstrates superior mechanical properties, boasting a flexural strength of 915.5 ± 43.1 MPa and a fracture toughness of 11.1 ± 0.3 MPa·m^1/2 [Supplementary Figure 11]. The strain distribution depicted in Figure 5A is derived from geometric phase analysis (GPA), with the corresponding strain mapping illustrated in Figure 5B. Around the hollow structure, symmetrical compression-tension strain pairs are induced due to lattice mismatch [Figure 5C], which presumably leads to significant local internal stresses and a stress shielding effect^[45], both of which contribute to enhanced friction resistance against motion or microcrack accumulation. Consequently, the strain couples associated with compressive-tensile forces surrounding hollow structures may present an additional avenue for improving the toughness of bulk ceramics. In this case, we evaluated the physical properties of reference ZrO₂ ceramics, Al₂O₃ ceramics, Si₃N₄ ceramics and the sample SEu-5. These properties included dielectric loss (tanδ), thermal conductivity (k), density (ρ), hardness (H), flexural strength (σ) and fracture toughness (K_IC) [Figure 5D and E, Supplementary Figure 1]. Our findings indicate that Si₃N₄ ceramics can compensate for the mechanical deficiencies of Al₂O₃ ceramics and the thermal and dielectric shortcomings of ZrO₂ ceramics, thereby exhibiting superior overall performance [Figure 5D]. In particular, the sample SEu-5 demonstrates superior physical properties compared to the referenced Si₃N₄ ceramics, with its tanδ, k, ρ, K_IC, σ, and H reaching 8.24 × 10^-4, 27 W/(m·K), 3.24 g·cm^-3, 11.1 MPa·m^1/2, 915.5 MPa, and 14.1 GPa, respectively [Figure 5E]. The tanδ and ρ of the SEu-5 are ~100 and ~2 times lower than those of ZrO₂ ceramics, reducing signal transmission absorbance and weight. The k of the SEu-5 is ~10 times higher than that of ZrO₂ ceramics, accelerating equipment cooling. The K_IC and σ of the SEu-5 are ~2 times that of Al₂O₃ ceramics, enhancing damage tolerance and yield improvement.

Figure 5. The strain distribution of the hollow structure and physical properties of Si₃N₄ ceramics. (A) HRTEM graph of the sample SEu-5 (the inset shows the corresponding FFT pattern). (B) Geometric Phase Analysis (GPA) mapping, derived from Digital Micrograph calculations, illustrates the overall strain field at a nanoscale. The color scale signifies the change in strain intensity, ranging from -30% (compressive) to +30% (tensile). (C) The profile of the white-dashed arrow in (B). (D) Physical performance of different ceramic materials with dielectric loss (tanδ), thermal conductivity (k), density (ρ), fracture toughness (K_IC), flexural strength (σ) and hardness (H) for referenced ZrO₂ ceramics, referenced Al₂O₃ ceramics and referenced Si₃N₄ ceramics plotted as a radar map: The performance of Si₃N₄ ceramics is outstanding. (E) Physical performance for referenced Si₃N₄ ceramics and the sample SEu-5: The sample SEu-5 exhibits better physical properties. The references for the physical properties of all ceramics in (D and E) are listed in the Supplementary Materials.

CONCLUSIONS

In summary, we introduce a microstructure design that involves the construction of Eu-doped hollow structures to synthesize Si₃N₄ ceramics. These ceramics exhibit a color transition from yellow to orange-red and possess superior mechanical properties. The hollow structure is characterized by its amorphous nature, with the core corresponding to the pore phase and the shell to the Eu-rich liquid phase. This structure functions as one of the dominant chromophores due to the 5d→4f transition of Eu²⁺, which is coupled with the ⁵D₀→⁷F_J transition of Eu³⁺ under photon excitation. Specially, the distribution density of this chromophore correlates positively with the depth of orange. Our sample, SEu-5, demonstrates optimized mechanical properties, with a flexural strength of 915.5 ± 43.1 MPa and a fracture toughness of 11.1 ± 0.3 MPa·m^1/2. Factors such as grain pull-out, bridging and deflection, pronounced intergranular fractures, and the stress field surrounding the hollow structure collectively enhance the mechanical properties by reducing crack energy consumption.