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Cost effective preparation of Si3N4 ceramics with improved thermal conductivity and mechanical properties
Journal of the European Ceramic Society xxx (xxxx) xxx–xxx
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Original Article
Cost effective preparation of Si3N4 ceramics with improved thermal conductivity and mechanical properties ⁎
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Yusen Duana,b, , Ning Liua, Jingxian Zhanga,c, , Hui Zhanga,c, Xiaoguang Lia a State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China b University of Chinese Academy of Sciences, Beijing, 100049, China c Suzhou Institute of SICCAS (Shanghai Institute of Ceramics, Chinese Academy of Sciences), Taicang, 215499, Jiangsu Province, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Silicon nitride Gas pressure sintering Thermal conductivity Nitridation process Mechanical properties
High-purity silicon powder is used as the starting material for cost-effective preparation of silicon nitride ceramics with both high thermal conductivity and excellent mechanical properties using RE2O3 (RE=Y, La or Er) and MgO as sintering additives. Nitridation is a key procedure that would affect the properties of green bodies and the sintered samples. The β: (α+β) ratio can be increased as the samples nitrided at 1450ºC and a large amount of long rod-like β-Si3N4 grains were developed in the samples. It was found that the addition of Er2O3-MgO could help to improve the mechanical properties of the sintered Si3N4 ceramics, the thermal conductivity, flexural strength and fracture toughness of the sample were 90 W/(m∙K), 953 ± 28.3 MPa and 10.64 ± 0.61 MPa·m1/2, respectively. The RE3+ species with larger ionic radius tended to increase the oxygen of nitrided samples and decrease N/O ratio (triangle grain boundary) of sintered samples.
1. Introduction Silicon nitride (Si3N4) ceramics exhibit a combination of excellent properties, such as outstanding mechanical properties, good resistance to high temperature creep and chemical attack behavior etc., have found a wide range of applications [1,2]. Calculation of theoretical thermal conductivity of Si3N4 has been conducted by several researchers. Haggerty et al. reported that the theoretical thermal conductivity of Si3N4 is as high as that of AlN (320 W/(m∙K)) [3], Hirosaki et al. showed that the intrinsic thermal conductivity of β-Si3N4 in a and c direction are 170 and 450 W/(m∙K), respectively [4]. Due to the combination of high thermal conductivity, excellent mechanical behavior, low thermal expansion coefficient etc. properties, Si3N4 ceramics have been considered as the potential substrates for high power electronic devices. However, the thermal conductivity of polycrystalline Si3N4 ceramic is much lower than the intrinsic value of the single crystal. As usual, the thermal conductivity of Si3N4 ceramic is mainly determined by the phonon mean free path. Lattice defects, secondary phases, grain boundary films and porosity all affect the phonon mean free path to some extent [5]. Thus, how to reduce the influence of above factors is the key issue for increasing the thermal conductivity of Si3N4 ceramics. As reported by literatures, there are two main factors
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influencing the thermal conductivity: 1) the oxygen content. Oxygen can dope into the Si3N4 lattice and form defects ((O↔ON+VSi)) to decrease the phonon mean free path [6,7]. Selecting higher purity powders as raw materials, proper sintering additives with low O affinity and less aids content are possible routes to avoid introducing much impurities [8,9]. Another way is prolonging the holding time to accelerate the O inside the lattice to diffuse out. 2) reducing of the amorphous grain boundary phases. This can be realized by prolonging holding time and decreasing the cooling rate. During this process, coarser Si3N4 grains would be developed and thus the grain boundary film thickness would be reduced, which will also contribute to the increase in thermal conductivity. A slow cooling rate will be beneficial for the crystallization of the amorphous phase [10]. Zhou et al. used high purity Si powder as raw material to prepare Si3N4 ceramics through sintered reaction-bonded silicon nitrides (SRBSN) technology, the thermal conductivity of the specimens reached 177 W/(m·K) after sintering at 1900 °C for 60 h followed by cooling at the rate of 0.2℃/min [11]. Besides, the alignment of Si3N4 grains is also an effective way to prepare high thermal conductivity ceramics along special directions because β-Si3N4 showed different thermal conductivity in different direction. Zhu et al. applied strong magnetic field on Si3N4 slurries and using sintering additives with low oxygen content (Y2O3–MgSiN2) to achieve high thermal conductivity ceramics (176 W/(m·K)) [12,13].
Corresponding authors. E-mail addresses: [email protected] (Y. Duan), [email protected] (J. Zhang).
https://doi.org/10.1016/j.jeurceramsoc.2019.10.003 Received 8 April 2019; Received in revised form 28 September 2019; Accepted 1 October 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
Please cite this article as: Yusen Duan, et al., Journal of the European Ceramic Society, https://doi.org/10.1016/j.jeurceramsoc.2019.10.003
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Table 1 Nitridation results of samples RLaM, RYM and RErM. Samples
Nitridation (%)
β/(α + β) (%)
Bulk density (g/cm3)
Relative density (%ρth)
Linear shrinkage (%)
Open porosity (%)
RLaM RYM RErM
96.27 94.83 97.31
100 92.4 87.4
2.30 2.20 2.39
69.97 67.81 71.99
2.59 2.58 3.93
29.52 32.13 28.18
using the Archimedes method in distilled water. Specimens for thermal conductivity measurement were machined into disk-like shape (10 mm in diameter and 2 mm in thickness). Thermal diffusivity (α) measurements were performed at room temperature on a laser-flash apparatus (LFA427 Nano ash, NETZSCH Instruments Co. Ltd, Selb, Germany). A constant value of specific heat, 0.68 J·g−1 K−1, was used in this work. Thermal conductivity (κ) can be calculated from the measured bulk density, thermal diffusivity and specific heat capacity using the following equation:
However, these methods usually require long holding time at high temperature, which will lead to the excessive grain growth and decrease the flexural strength. In literature, after sintering at 1900℃ with short holding (3 h), the thermal conductivity of the sample was not so high (107 W/(m·K)) [10]. Therefore, the preparation of Si3N4 ceramics with both high thermal conductivity and high mechanical properties is still a challenge. It was found that the content of β-Si3N4 in the starting material would influence the thermal conductivity properties of the final sintered samples. Moulson [14] found that the α and β phases were formed through “high oxygen potential” and “low oxygen potential” modifications and later they described α-Si3N4 as a defect structure with oxygen replacing nitrogen at some sites and silicon vacancies formed for maintaining electrical neutrality. This would lead to the phonon scattering and reduce greatly the thermal conductivity. However, if βSi3N4 powder is directly used as the raw material, higher temperatures, longer-sintering times, or both are needed to prepare high performance ceramics [15]. In literature, it was confirmed that the α phase can be transformed to β-Si3N4 phase completely after nitridation process [16]. Therefore, using SRBSN method with high β phase content after nitriding may be one possible route for preparing high thermal conductivity Si3N4. In addition, in order to accelerate the α to β phase transformation and improve the mechanical properties of the ceramics after sintering, Y2O3 and Er2O3 with small ion radius were used as sintering additives based on the report in literature. In this study, the RE2O3 (RE = Y, La or Er) and MgO were used as sintering additives, and the nitridation process was adjusted to increase β-Si3N4 content in the green body. The effects of sintering additives and nitridation process on the sintering behaviors, microstructure, oxygen content, mechanical strength and thermal conductivity were investigated systemically.
κ = ρcp α The phase compositions were identified by X-ray diffraction (XRD, D8 Advance, Bruker, Germany) using Cu Kα. The microstructure was characterized by scanning electron microscopy (SEM; Magellan 400, FEI Co., Hillsboro, OR, USA) and transmission electron microscopy (TEM; Tecnai G2 F20, FEI Co., Hillsboro, OR, USA). For TEM samples, thin foils were prepared by cutting, polishing, dimpling, and followed by argon-ion-milling until perforation. To measure flexural strength via the three-point bending test (Model 5566, Instron Co., High Wycombe, UK), the specimens were machined into rectangle bars with the dimension of 3 × 4×36 mm. Fracture toughness was measured by singleedge notched beam method, and a minimum of 6 rectangular bars with dimensions of 3 mm × 6 mm × 30 mm were tested. 3. Results and discussion 3.1. Nitridation of Si compacts The nitridation results of the reaction-bonded silicon nitride (RBSN) samples are listed in Table 1. All the samples show more than 94% nitridation, and there is no Si peak observed by XRD (in Fig. 1). Thus, the nitridation is considered to be nearly complete. As expected, the phase content ratios β: (α + β) of all the samples were 100%, 92.4% and 87.4%, respectively. However, the ratios were usually lower than 15% as the Si compacts were nitrided at 1400℃ for 8 h in our previous experiments. This is because that the formation of α-Si3N4 phase was favored at reaction temperature below the melting point of silicon [17]. As reported by Becher [18,19] and Kitayama [20], the phase transformation rate showed negative correlation with radius of rare earth cations. In their work, Si3N4 powder was used as raw material and the transformation initiated at or above 1500℃. However, the opposed results were observed in this study by using high purity Si powder as starting material. The reaction started from the surface of particles, forming a layer of Si3N4 surrounding a reactant core, and continued through nitrogen diffusing inwards and silicon diffusing outwards through the shell to the reaction sites [21], the reaction rate was quite slowly and the α-Si3N4 was the main product. As the nitridation temperature higher than the melting point of silicon (1410℃), β-Si3N4 was formed by the gas diffusion inside the β-Si3N4 solid phase and reacted with liquid Si. At 1450℃, all samples show linear shrinkage higher than 2.5% compared to the shrinkages of samples reported in in literatures (less than 2%) [16,22]. The obtained samples had higher ratio of β/(α + β), higher open porosity and lower relative density, shown in Table 1. This phenomenon was consistent with the result in literature [23], the special rod-like structure of β-Si3N4 may occupy much more interspace than granular structure of α-Si3N4.
2. Experiment procedure The commercial high purity Si powder (99.999%, BET 6.18 m2/g, 0.88 wt% O, D50 = 1 μm, Haoxi Research Nanomaterials, Inc., Shanghai, China) was used in the present study. High purity Y2O3, La2O3 and Er2O3 (purity > 99%; Haoxi Research Nanomaterials, Inc., Shanghai, China) and MgO (purity > 99%; Qinhuangdao Yi Nuo Co, Ltd, China) were used as sintering additives. The composition of Si compacts with sintering additives was determined based on the final composition of the RBSN materials after full nitridation as Si3N4: RE2O3: MgO = 93: 2: 5 at molar ratio, designated as samples RYM, RLaM and RErM, respectively. The Si3N4, RE2O3 and MgO powders with defined composition were mixed in ethyl alcohol using a planetary mill for 4 h in a polyethylene jar with Si3N4 balls as the media, and then the obtained slurries were dried using rotary evaporator. After drying, the as-received Si powder showed a decreased D50 of 0.85 μm and an increased oxygen content of 1.14 wt%. The powders were sieved through 100-mesh screen and uniaxially die-pressed under 30 MPa and subsequently cold isostatically pressed at 200 MPa. The Si compacts were nitrided in graphite tube furnace with a mixed gas (N2:H2 = 95:5) at 1400℃ for 2 h and then further increased the temperature to 1450℃ for 6 h. The post-sintering was performed in graphite resistance furnace at 1900℃ for 2 h under a nitrogen pressure of 5 MPa with heating and cooling rates of 10℃/min. The bulk densities (ρ) of the as-sintered samples were measured 2
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Fig. 1. SEM images on the fracture surfaces of nitrided samples at 1450℃ for 2h: (a) RLaM, (b) RYM and (c) RErM. The figures in the upper-right corner are XRD patterns of corresponding samples.
As observed in Fig. 1, there are many rod-like particles in all SBSN samples. The needle-like Si3N4 grains are formed by gas-liquid reaction at the solid-liquid interface [24]. The reaction site is probably at the tip of the spike, thus extending its length in the fast-growing z-direction [25]. Moreover, the rod-like particles had obviously difference morphology with different sintering additives. Combination of XRD analysis results inserted into SEM micrograph for reference, Fig. 1, higher content β-Si3N4 phase in all samples can be observed. Especially, the RBSN sample RLaM had smaller grain length (less than 1 μm). On the contrary, the lengths of particles were in the range of 1 μm and 2 μm with also bigger particles in samples RYM and RErM, respectively. XRD pattern in Fig. 1a reveals that there was only β-Si3N4 phase in the nitrided sample RLaM. The segregation to and the adsorption of the La3+ on the N-rich Si3N4 grain surfaces was easier than Y3+ and Er3+ [19], therefore, less La3+ existed in glass phase and no crystalline phase contain La3+ was formed after cooling. The main crystalline phases in samples RYM and RErM were β- and α-Si3N4, while the peaks of β-Si3N4 were stronger than those of α-Si3N4. In addition to the Si3N4 phases, the crystal secondary phases were formed in the nitrided samples RYM and RErM. In sample RYM, both the YSiO2N and Y2Si3O3N4 were observed. The phase of ErSiO(N) was observed in sample RErM and cannot fit PDF cards, but the positions of diffraction peaks were similar to that of Y2Si3O3N4 [26]. The stoichiometric ratio of ErSiO(N) phase was still uncertain. Moreover, the liquid phase formed by RE2O3-Si3N4-MgOSiO2 system showed the formation of clusters, which was almost not referred in literatures. Especially, it was obvious that the liquid phase was unevenly distributed in sample RErM. Golla et al. also increased the nitridation temperature to 1450℃ for 2.5 h with Y2O3-MgO as sintering additives, but there were no obviously rod-like grains, liquid phases regions and the highest ratio of β: (α + β) was 63.2% [27]. As reported by Bhatt et al., the content of oxygen decreases with the increase in nitridation temperature [28]. As shown in Table 2, the
Table 2 Oxygen content of nitrided samples RLaM, RYM and RErM.
1400℃ 1450℃
RLaM
RYM
RErM
4.24 wt% 2.62 wt%
3.24 wt% 1.98 wt%
3.13 wt% 1.84 wt%
oxygen contents were 3.13 wt% and 1.84 wt% of samples RErM were nitrided at 1400℃ and 1450℃, respectively. The oxygen contents of nitrided samples RYM were similar to samples RErM. Besides, the samples RLaM had the highest oxygen contents at different heating temperatures. In Sample RLaM, the high O content indicated a low N/O ratio and a low glass transition temperature of La-Si-Mg-O-N secondary phase [29]. Thus, the secondary phase was probably in an amorphous state and undetectable by XRD. The high content of β-Si3N4 phase might help to increase the thermal conductivity and mechanical strength of final sintered Si3N4 ceramics because β-Si3N4 grains has low oxygen content.
3.2. Densification and thermal conductivity Table 3 shows the properties of samples after sintering at 1900℃ for 2 h. All samples have been well densified. The weight losses of all samples are less than 3% in the absence of powder bed, this may be caused by higher nitrogen pressure. Samples RLaM and RYM show higher shrinkage than RErM due to lower relative densities of the green bodies. Sample RErM had slightly higher thermal conductivity of about 90 W/(m·K) than the other two samples. The lower relative density of sample RLaM may result in more defects and lead to lower thermal diffusivity of 36.01 mm2·s−1. In addition, the even β-Si3N4 grain sizes of reacted samples may result in less content of grain boundaries after 3
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Table 3 Linear shrinkage, weight loss, density, thermal diffusivity and thermal conductivity for sintered SRBSNs. Samples
Linear shrinkage (%)
Weight loss (%)
Bulk density (g/ cm3)
Relative density (%ρth)
Open porosity (%)
Thermal diffusivity (mm2·s−1)
Thermal conductivity (W/ (m·K))
RLaM RYM RErM
11.82 11.4 10.97
2.61 1.91 2.73
3.25 3.23 3.29
98.64 99.6 99.06
0.32 0.44 0.25
36.01 37.07 40.3
79.6 81.4 90
Fig. 2. SEM micrographs of the fracture surface and polished surfaces of SRBSN samples at 1900 °C for 2h: (a, d) RLaM, (b, e)RYM and (c, f)RErM. The figures in the upper-right corner are XRD patterns of corresponding samples.
all samples after post sintering. A lot of small grains existed in submicron scale around larger grains. These fine grains would provide a big amount of interface and influence the phonon mean free path as well as thermal conductivity. From Fig. 2d-f, sample RLaM had smaller grain diameter than other two samples, but the grain lengths of all samples are similar. However, the aspect ratio increased with an
post-sintering. Thus, higher thermal conductivities of samples were obtained after very short soaking (2 h).
3.3. Microstructure and mechanical properties of SRBSNs As shown in Fig. 2, a similar bimodal microstructure developed for 4
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Fig. 3. TEM micrographs of multigrain junctions in (a, b) RErM. At the upper-right of each figure is the selected area electron diffraction (SAED) pattern of the region 1. Area (c) and site 2 (d) EDS results referred to Fig. 3b.
phase. The average N/O ratio in triangle areas of samples RLaM, RYM and RErM was 13.39 ± 1.78, 18.64 ± 0.48 and 18.79 ± 0.29%, respectively. The content was also measured by TEM with EDS. These results were consistent with oxygen contents of nitrided samples, thermal conductivity and fracture toughness of sintered samples. The higher N content can make the glass structure more highly cross-linked and tight with twofold-coordinated oxygen atom [23,31]. It will be more beneficial for the thermal conductivity improvement by reducing the SiO2 content through prolonging of sintering time or providing reduced atmosphere. However, this will also lead to the grain growth and the decrease in mechanical properties. Fig. 4 shows the elements distribution of sample RErM shown in Fig. 2d, it can be seen that the elements of Er, Mg and O distributed at the interface and triple point of Si3N4 grains. The compounds may be glassy MgSiO3, though it cannot be detected by XRD. Besides, the contents of sintering additives may be too much for the preparation of higher thermal conductivity Si3N4 ceramics due to the short soaking time. As shown in Fig. 5, samples RYM and RErM have higher fracture toughness of 10.41 and 10.6 MPa m1/2 than sample RLaM of 8.31 MPa m1/2. As mentioned above, samples RYM and RErM had similar grain sizes, and bigger than sample RLaM obviously. Meantime, the fracture toughness results were also the same trend, thus the fracture toughness may be determined by the grain sizes of the samples, crack deflection and particle pull-out would decrease the crack tip
increase in cationic radius as reported by Becher [18]. In this case, the majority of the Er3+ existed in the secondary phases, which would limit their adsorption at the grain surfaces as a result of it greater affinity for oxygen than for nitrogen [19,30]. This character may decrease content of lattice defects and improve thermal conductivity. Consequently, sample RErM shows higher thermal conductivity than the other two samples. According to the XRD patterns inserted in Fig. 2d-f, there is still only β-Si3N4 phase can be detected in the post-sintered sample RLaM, and the Y2Si3O3N4 and ErSiO(N) phases are located at the similar positions of sample RYM and RErM, respectively. However, the crystalinity of the secondary phase for RErM was higher than that in RYM. This might contribute to the difference in thermal conductivity. Compared with the results by Zhou et al. [10] with very slow cooling rate (0.2℃/min), the higher cooling rate (10℃/min) may influence the crystalization of the second phases and further affect the thermal conductivity of the samples. Selected area electron diffraction (SAED) analysis was used to characterize the region of red site 1 in Fig. 3a. The region1 can be referred to as the secondary phase with low crystallinity, consisted with the XRD and SEM analysis. Fig. 3c and d show elements content of the area and site 2 in Fig.3b. The content of Mg element was higher than that reported in literature [29], which might be due to the shorter sintering time. Besides, the N/O ratio in triangle area (site 2) was still too low, probably due to the very limited release of SiO2 from liquid
5
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Fig. 4. EDS mapping showing the elements distribution in the RErM specimen.
samples showed higher fracture toughness and lower flexural strength [33,34]. However, the sample RLaM had the lowest fracture toughness in this study and the similar results were reported by Horng et al [35]. In literatures, it is difficult to develop Si3N4 ceramics with high flexural strength, fracture toughness and together with high thermal conductivity. As reported by Kusano et al, the flexural strength and fracture toughness of Si3N4 sample were 801 ± 72 MPa and 7.3 ± 0.3 MPa·m1/2, respectively [36]. 4. Conclusions In this study, dense Si3N4 ceramics were prepared by gas pressure sintering at 1900 °C for 2 h with rare earth oxide (RE = Y, La or Er) and MgO as sintering additives. It was found that with the increase in nitridation temperature to 1450 °C, the ratios of β: (α + β) can be increased for all Si3N4 green bodies. After sintering, elongated grains with high aspect ratio can be developed with less second phase and high crystallinity. The oxygen contents of nitrided samples were consistent with average N/O ratio in triangle areas, thermal conductivity and fracture toughness of sintered samples. The thermal conductivity,
Fig. 5. Flexural strength and fracture toughness of sintered specimens.
stress [32] (i.e. self-toughness mechanism). Compared with sample RErM, samples RYM and RLaM had similar flexural strength and lower than 800 MPa. The results show that sample RErM had the highest flexural strength of 953 ± 28.3 MPa. The RE3+ species with larger ionic radius tend to give weaker grain boundary adhesion, and the 6
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flexural strength and fracture toughness were 90 W/(m∙K), 953 ± 28.3 MPa and 10.64 ± 0.61 MPa·m1/2 for sample RErM, respectively. Thus, increasing the β-Si3N4 content in green bodies can help to increase the thermal conductivity to some extent. The high strength, toughness and thermal conductivity might be due to the small radius of Er3+, which help to form the more viscous liquid phase and decrease lattice defects. It can be concluded that the Er2O3-MgO was an effective sintering additive system to prepare Si3N4 ceramics with excellent mechanical properties and high thermal conductivity through SRBSN method.
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This work was supported by the National Key Research and Development Program of China (2016YFB-0700300), National Natural Science Foundation of China (No.51572277,51702340), Shanghai Science and Technology Committee (17YF1428800, 17ZR1434800, 17dz2307000), and the State Key Laboratory of High Performance Ceramics and Superfine Microstructure of Shanghai Institute Ceramics, Chinese Academy of Sciences.
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Original Article
Cost effective preparation of Si3N4 ceramics with improved thermal conductivity and mechanical properties ⁎
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Yusen Duana,b, , Ning Liua, Jingxian Zhanga,c, , Hui Zhanga,c, Xiaoguang Lia a State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China b University of Chinese Academy of Sciences, Beijing, 100049, China c Suzhou Institute of SICCAS (Shanghai Institute of Ceramics, Chinese Academy of Sciences), Taicang, 215499, Jiangsu Province, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Silicon nitride Gas pressure sintering Thermal conductivity Nitridation process Mechanical properties
High-purity silicon powder is used as the starting material for cost-effective preparation of silicon nitride ceramics with both high thermal conductivity and excellent mechanical properties using RE2O3 (RE=Y, La or Er) and MgO as sintering additives. Nitridation is a key procedure that would affect the properties of green bodies and the sintered samples. The β: (α+β) ratio can be increased as the samples nitrided at 1450ºC and a large amount of long rod-like β-Si3N4 grains were developed in the samples. It was found that the addition of Er2O3-MgO could help to improve the mechanical properties of the sintered Si3N4 ceramics, the thermal conductivity, flexural strength and fracture toughness of the sample were 90 W/(m∙K), 953 ± 28.3 MPa and 10.64 ± 0.61 MPa·m1/2, respectively. The RE3+ species with larger ionic radius tended to increase the oxygen of nitrided samples and decrease N/O ratio (triangle grain boundary) of sintered samples.
1. Introduction Silicon nitride (Si3N4) ceramics exhibit a combination of excellent properties, such as outstanding mechanical properties, good resistance to high temperature creep and chemical attack behavior etc., have found a wide range of applications [1,2]. Calculation of theoretical thermal conductivity of Si3N4 has been conducted by several researchers. Haggerty et al. reported that the theoretical thermal conductivity of Si3N4 is as high as that of AlN (320 W/(m∙K)) [3], Hirosaki et al. showed that the intrinsic thermal conductivity of β-Si3N4 in a and c direction are 170 and 450 W/(m∙K), respectively [4]. Due to the combination of high thermal conductivity, excellent mechanical behavior, low thermal expansion coefficient etc. properties, Si3N4 ceramics have been considered as the potential substrates for high power electronic devices. However, the thermal conductivity of polycrystalline Si3N4 ceramic is much lower than the intrinsic value of the single crystal. As usual, the thermal conductivity of Si3N4 ceramic is mainly determined by the phonon mean free path. Lattice defects, secondary phases, grain boundary films and porosity all affect the phonon mean free path to some extent [5]. Thus, how to reduce the influence of above factors is the key issue for increasing the thermal conductivity of Si3N4 ceramics. As reported by literatures, there are two main factors
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influencing the thermal conductivity: 1) the oxygen content. Oxygen can dope into the Si3N4 lattice and form defects ((O↔ON+VSi)) to decrease the phonon mean free path [6,7]. Selecting higher purity powders as raw materials, proper sintering additives with low O affinity and less aids content are possible routes to avoid introducing much impurities [8,9]. Another way is prolonging the holding time to accelerate the O inside the lattice to diffuse out. 2) reducing of the amorphous grain boundary phases. This can be realized by prolonging holding time and decreasing the cooling rate. During this process, coarser Si3N4 grains would be developed and thus the grain boundary film thickness would be reduced, which will also contribute to the increase in thermal conductivity. A slow cooling rate will be beneficial for the crystallization of the amorphous phase [10]. Zhou et al. used high purity Si powder as raw material to prepare Si3N4 ceramics through sintered reaction-bonded silicon nitrides (SRBSN) technology, the thermal conductivity of the specimens reached 177 W/(m·K) after sintering at 1900 °C for 60 h followed by cooling at the rate of 0.2℃/min [11]. Besides, the alignment of Si3N4 grains is also an effective way to prepare high thermal conductivity ceramics along special directions because β-Si3N4 showed different thermal conductivity in different direction. Zhu et al. applied strong magnetic field on Si3N4 slurries and using sintering additives with low oxygen content (Y2O3–MgSiN2) to achieve high thermal conductivity ceramics (176 W/(m·K)) [12,13].
Corresponding authors. E-mail addresses: [email protected] (Y. Duan), [email protected] (J. Zhang).
https://doi.org/10.1016/j.jeurceramsoc.2019.10.003 Received 8 April 2019; Received in revised form 28 September 2019; Accepted 1 October 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
Please cite this article as: Yusen Duan, et al., Journal of the European Ceramic Society, https://doi.org/10.1016/j.jeurceramsoc.2019.10.003
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Table 1 Nitridation results of samples RLaM, RYM and RErM. Samples
Nitridation (%)
β/(α + β) (%)
Bulk density (g/cm3)
Relative density (%ρth)
Linear shrinkage (%)
Open porosity (%)
RLaM RYM RErM
96.27 94.83 97.31
100 92.4 87.4
2.30 2.20 2.39
69.97 67.81 71.99
2.59 2.58 3.93
29.52 32.13 28.18
using the Archimedes method in distilled water. Specimens for thermal conductivity measurement were machined into disk-like shape (10 mm in diameter and 2 mm in thickness). Thermal diffusivity (α) measurements were performed at room temperature on a laser-flash apparatus (LFA427 Nano ash, NETZSCH Instruments Co. Ltd, Selb, Germany). A constant value of specific heat, 0.68 J·g−1 K−1, was used in this work. Thermal conductivity (κ) can be calculated from the measured bulk density, thermal diffusivity and specific heat capacity using the following equation:
However, these methods usually require long holding time at high temperature, which will lead to the excessive grain growth and decrease the flexural strength. In literature, after sintering at 1900℃ with short holding (3 h), the thermal conductivity of the sample was not so high (107 W/(m·K)) [10]. Therefore, the preparation of Si3N4 ceramics with both high thermal conductivity and high mechanical properties is still a challenge. It was found that the content of β-Si3N4 in the starting material would influence the thermal conductivity properties of the final sintered samples. Moulson [14] found that the α and β phases were formed through “high oxygen potential” and “low oxygen potential” modifications and later they described α-Si3N4 as a defect structure with oxygen replacing nitrogen at some sites and silicon vacancies formed for maintaining electrical neutrality. This would lead to the phonon scattering and reduce greatly the thermal conductivity. However, if βSi3N4 powder is directly used as the raw material, higher temperatures, longer-sintering times, or both are needed to prepare high performance ceramics [15]. In literature, it was confirmed that the α phase can be transformed to β-Si3N4 phase completely after nitridation process [16]. Therefore, using SRBSN method with high β phase content after nitriding may be one possible route for preparing high thermal conductivity Si3N4. In addition, in order to accelerate the α to β phase transformation and improve the mechanical properties of the ceramics after sintering, Y2O3 and Er2O3 with small ion radius were used as sintering additives based on the report in literature. In this study, the RE2O3 (RE = Y, La or Er) and MgO were used as sintering additives, and the nitridation process was adjusted to increase β-Si3N4 content in the green body. The effects of sintering additives and nitridation process on the sintering behaviors, microstructure, oxygen content, mechanical strength and thermal conductivity were investigated systemically.
κ = ρcp α The phase compositions were identified by X-ray diffraction (XRD, D8 Advance, Bruker, Germany) using Cu Kα. The microstructure was characterized by scanning electron microscopy (SEM; Magellan 400, FEI Co., Hillsboro, OR, USA) and transmission electron microscopy (TEM; Tecnai G2 F20, FEI Co., Hillsboro, OR, USA). For TEM samples, thin foils were prepared by cutting, polishing, dimpling, and followed by argon-ion-milling until perforation. To measure flexural strength via the three-point bending test (Model 5566, Instron Co., High Wycombe, UK), the specimens were machined into rectangle bars with the dimension of 3 × 4×36 mm. Fracture toughness was measured by singleedge notched beam method, and a minimum of 6 rectangular bars with dimensions of 3 mm × 6 mm × 30 mm were tested. 3. Results and discussion 3.1. Nitridation of Si compacts The nitridation results of the reaction-bonded silicon nitride (RBSN) samples are listed in Table 1. All the samples show more than 94% nitridation, and there is no Si peak observed by XRD (in Fig. 1). Thus, the nitridation is considered to be nearly complete. As expected, the phase content ratios β: (α + β) of all the samples were 100%, 92.4% and 87.4%, respectively. However, the ratios were usually lower than 15% as the Si compacts were nitrided at 1400℃ for 8 h in our previous experiments. This is because that the formation of α-Si3N4 phase was favored at reaction temperature below the melting point of silicon [17]. As reported by Becher [18,19] and Kitayama [20], the phase transformation rate showed negative correlation with radius of rare earth cations. In their work, Si3N4 powder was used as raw material and the transformation initiated at or above 1500℃. However, the opposed results were observed in this study by using high purity Si powder as starting material. The reaction started from the surface of particles, forming a layer of Si3N4 surrounding a reactant core, and continued through nitrogen diffusing inwards and silicon diffusing outwards through the shell to the reaction sites [21], the reaction rate was quite slowly and the α-Si3N4 was the main product. As the nitridation temperature higher than the melting point of silicon (1410℃), β-Si3N4 was formed by the gas diffusion inside the β-Si3N4 solid phase and reacted with liquid Si. At 1450℃, all samples show linear shrinkage higher than 2.5% compared to the shrinkages of samples reported in in literatures (less than 2%) [16,22]. The obtained samples had higher ratio of β/(α + β), higher open porosity and lower relative density, shown in Table 1. This phenomenon was consistent with the result in literature [23], the special rod-like structure of β-Si3N4 may occupy much more interspace than granular structure of α-Si3N4.
2. Experiment procedure The commercial high purity Si powder (99.999%, BET 6.18 m2/g, 0.88 wt% O, D50 = 1 μm, Haoxi Research Nanomaterials, Inc., Shanghai, China) was used in the present study. High purity Y2O3, La2O3 and Er2O3 (purity > 99%; Haoxi Research Nanomaterials, Inc., Shanghai, China) and MgO (purity > 99%; Qinhuangdao Yi Nuo Co, Ltd, China) were used as sintering additives. The composition of Si compacts with sintering additives was determined based on the final composition of the RBSN materials after full nitridation as Si3N4: RE2O3: MgO = 93: 2: 5 at molar ratio, designated as samples RYM, RLaM and RErM, respectively. The Si3N4, RE2O3 and MgO powders with defined composition were mixed in ethyl alcohol using a planetary mill for 4 h in a polyethylene jar with Si3N4 balls as the media, and then the obtained slurries were dried using rotary evaporator. After drying, the as-received Si powder showed a decreased D50 of 0.85 μm and an increased oxygen content of 1.14 wt%. The powders were sieved through 100-mesh screen and uniaxially die-pressed under 30 MPa and subsequently cold isostatically pressed at 200 MPa. The Si compacts were nitrided in graphite tube furnace with a mixed gas (N2:H2 = 95:5) at 1400℃ for 2 h and then further increased the temperature to 1450℃ for 6 h. The post-sintering was performed in graphite resistance furnace at 1900℃ for 2 h under a nitrogen pressure of 5 MPa with heating and cooling rates of 10℃/min. The bulk densities (ρ) of the as-sintered samples were measured 2
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Fig. 1. SEM images on the fracture surfaces of nitrided samples at 1450℃ for 2h: (a) RLaM, (b) RYM and (c) RErM. The figures in the upper-right corner are XRD patterns of corresponding samples.
As observed in Fig. 1, there are many rod-like particles in all SBSN samples. The needle-like Si3N4 grains are formed by gas-liquid reaction at the solid-liquid interface [24]. The reaction site is probably at the tip of the spike, thus extending its length in the fast-growing z-direction [25]. Moreover, the rod-like particles had obviously difference morphology with different sintering additives. Combination of XRD analysis results inserted into SEM micrograph for reference, Fig. 1, higher content β-Si3N4 phase in all samples can be observed. Especially, the RBSN sample RLaM had smaller grain length (less than 1 μm). On the contrary, the lengths of particles were in the range of 1 μm and 2 μm with also bigger particles in samples RYM and RErM, respectively. XRD pattern in Fig. 1a reveals that there was only β-Si3N4 phase in the nitrided sample RLaM. The segregation to and the adsorption of the La3+ on the N-rich Si3N4 grain surfaces was easier than Y3+ and Er3+ [19], therefore, less La3+ existed in glass phase and no crystalline phase contain La3+ was formed after cooling. The main crystalline phases in samples RYM and RErM were β- and α-Si3N4, while the peaks of β-Si3N4 were stronger than those of α-Si3N4. In addition to the Si3N4 phases, the crystal secondary phases were formed in the nitrided samples RYM and RErM. In sample RYM, both the YSiO2N and Y2Si3O3N4 were observed. The phase of ErSiO(N) was observed in sample RErM and cannot fit PDF cards, but the positions of diffraction peaks were similar to that of Y2Si3O3N4 [26]. The stoichiometric ratio of ErSiO(N) phase was still uncertain. Moreover, the liquid phase formed by RE2O3-Si3N4-MgOSiO2 system showed the formation of clusters, which was almost not referred in literatures. Especially, it was obvious that the liquid phase was unevenly distributed in sample RErM. Golla et al. also increased the nitridation temperature to 1450℃ for 2.5 h with Y2O3-MgO as sintering additives, but there were no obviously rod-like grains, liquid phases regions and the highest ratio of β: (α + β) was 63.2% [27]. As reported by Bhatt et al., the content of oxygen decreases with the increase in nitridation temperature [28]. As shown in Table 2, the
Table 2 Oxygen content of nitrided samples RLaM, RYM and RErM.
1400℃ 1450℃
RLaM
RYM
RErM
4.24 wt% 2.62 wt%
3.24 wt% 1.98 wt%
3.13 wt% 1.84 wt%
oxygen contents were 3.13 wt% and 1.84 wt% of samples RErM were nitrided at 1400℃ and 1450℃, respectively. The oxygen contents of nitrided samples RYM were similar to samples RErM. Besides, the samples RLaM had the highest oxygen contents at different heating temperatures. In Sample RLaM, the high O content indicated a low N/O ratio and a low glass transition temperature of La-Si-Mg-O-N secondary phase [29]. Thus, the secondary phase was probably in an amorphous state and undetectable by XRD. The high content of β-Si3N4 phase might help to increase the thermal conductivity and mechanical strength of final sintered Si3N4 ceramics because β-Si3N4 grains has low oxygen content.
3.2. Densification and thermal conductivity Table 3 shows the properties of samples after sintering at 1900℃ for 2 h. All samples have been well densified. The weight losses of all samples are less than 3% in the absence of powder bed, this may be caused by higher nitrogen pressure. Samples RLaM and RYM show higher shrinkage than RErM due to lower relative densities of the green bodies. Sample RErM had slightly higher thermal conductivity of about 90 W/(m·K) than the other two samples. The lower relative density of sample RLaM may result in more defects and lead to lower thermal diffusivity of 36.01 mm2·s−1. In addition, the even β-Si3N4 grain sizes of reacted samples may result in less content of grain boundaries after 3
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Table 3 Linear shrinkage, weight loss, density, thermal diffusivity and thermal conductivity for sintered SRBSNs. Samples
Linear shrinkage (%)
Weight loss (%)
Bulk density (g/ cm3)
Relative density (%ρth)
Open porosity (%)
Thermal diffusivity (mm2·s−1)
Thermal conductivity (W/ (m·K))
RLaM RYM RErM
11.82 11.4 10.97
2.61 1.91 2.73
3.25 3.23 3.29
98.64 99.6 99.06
0.32 0.44 0.25
36.01 37.07 40.3
79.6 81.4 90
Fig. 2. SEM micrographs of the fracture surface and polished surfaces of SRBSN samples at 1900 °C for 2h: (a, d) RLaM, (b, e)RYM and (c, f)RErM. The figures in the upper-right corner are XRD patterns of corresponding samples.
all samples after post sintering. A lot of small grains existed in submicron scale around larger grains. These fine grains would provide a big amount of interface and influence the phonon mean free path as well as thermal conductivity. From Fig. 2d-f, sample RLaM had smaller grain diameter than other two samples, but the grain lengths of all samples are similar. However, the aspect ratio increased with an
post-sintering. Thus, higher thermal conductivities of samples were obtained after very short soaking (2 h).
3.3. Microstructure and mechanical properties of SRBSNs As shown in Fig. 2, a similar bimodal microstructure developed for 4
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Fig. 3. TEM micrographs of multigrain junctions in (a, b) RErM. At the upper-right of each figure is the selected area electron diffraction (SAED) pattern of the region 1. Area (c) and site 2 (d) EDS results referred to Fig. 3b.
phase. The average N/O ratio in triangle areas of samples RLaM, RYM and RErM was 13.39 ± 1.78, 18.64 ± 0.48 and 18.79 ± 0.29%, respectively. The content was also measured by TEM with EDS. These results were consistent with oxygen contents of nitrided samples, thermal conductivity and fracture toughness of sintered samples. The higher N content can make the glass structure more highly cross-linked and tight with twofold-coordinated oxygen atom [23,31]. It will be more beneficial for the thermal conductivity improvement by reducing the SiO2 content through prolonging of sintering time or providing reduced atmosphere. However, this will also lead to the grain growth and the decrease in mechanical properties. Fig. 4 shows the elements distribution of sample RErM shown in Fig. 2d, it can be seen that the elements of Er, Mg and O distributed at the interface and triple point of Si3N4 grains. The compounds may be glassy MgSiO3, though it cannot be detected by XRD. Besides, the contents of sintering additives may be too much for the preparation of higher thermal conductivity Si3N4 ceramics due to the short soaking time. As shown in Fig. 5, samples RYM and RErM have higher fracture toughness of 10.41 and 10.6 MPa m1/2 than sample RLaM of 8.31 MPa m1/2. As mentioned above, samples RYM and RErM had similar grain sizes, and bigger than sample RLaM obviously. Meantime, the fracture toughness results were also the same trend, thus the fracture toughness may be determined by the grain sizes of the samples, crack deflection and particle pull-out would decrease the crack tip
increase in cationic radius as reported by Becher [18]. In this case, the majority of the Er3+ existed in the secondary phases, which would limit their adsorption at the grain surfaces as a result of it greater affinity for oxygen than for nitrogen [19,30]. This character may decrease content of lattice defects and improve thermal conductivity. Consequently, sample RErM shows higher thermal conductivity than the other two samples. According to the XRD patterns inserted in Fig. 2d-f, there is still only β-Si3N4 phase can be detected in the post-sintered sample RLaM, and the Y2Si3O3N4 and ErSiO(N) phases are located at the similar positions of sample RYM and RErM, respectively. However, the crystalinity of the secondary phase for RErM was higher than that in RYM. This might contribute to the difference in thermal conductivity. Compared with the results by Zhou et al. [10] with very slow cooling rate (0.2℃/min), the higher cooling rate (10℃/min) may influence the crystalization of the second phases and further affect the thermal conductivity of the samples. Selected area electron diffraction (SAED) analysis was used to characterize the region of red site 1 in Fig. 3a. The region1 can be referred to as the secondary phase with low crystallinity, consisted with the XRD and SEM analysis. Fig. 3c and d show elements content of the area and site 2 in Fig.3b. The content of Mg element was higher than that reported in literature [29], which might be due to the shorter sintering time. Besides, the N/O ratio in triangle area (site 2) was still too low, probably due to the very limited release of SiO2 from liquid
5
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Fig. 4. EDS mapping showing the elements distribution in the RErM specimen.
samples showed higher fracture toughness and lower flexural strength [33,34]. However, the sample RLaM had the lowest fracture toughness in this study and the similar results were reported by Horng et al [35]. In literatures, it is difficult to develop Si3N4 ceramics with high flexural strength, fracture toughness and together with high thermal conductivity. As reported by Kusano et al, the flexural strength and fracture toughness of Si3N4 sample were 801 ± 72 MPa and 7.3 ± 0.3 MPa·m1/2, respectively [36]. 4. Conclusions In this study, dense Si3N4 ceramics were prepared by gas pressure sintering at 1900 °C for 2 h with rare earth oxide (RE = Y, La or Er) and MgO as sintering additives. It was found that with the increase in nitridation temperature to 1450 °C, the ratios of β: (α + β) can be increased for all Si3N4 green bodies. After sintering, elongated grains with high aspect ratio can be developed with less second phase and high crystallinity. The oxygen contents of nitrided samples were consistent with average N/O ratio in triangle areas, thermal conductivity and fracture toughness of sintered samples. The thermal conductivity,
Fig. 5. Flexural strength and fracture toughness of sintered specimens.
stress [32] (i.e. self-toughness mechanism). Compared with sample RErM, samples RYM and RLaM had similar flexural strength and lower than 800 MPa. The results show that sample RErM had the highest flexural strength of 953 ± 28.3 MPa. The RE3+ species with larger ionic radius tend to give weaker grain boundary adhesion, and the 6
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flexural strength and fracture toughness were 90 W/(m∙K), 953 ± 28.3 MPa and 10.64 ± 0.61 MPa·m1/2 for sample RErM, respectively. Thus, increasing the β-Si3N4 content in green bodies can help to increase the thermal conductivity to some extent. The high strength, toughness and thermal conductivity might be due to the small radius of Er3+, which help to form the more viscous liquid phase and decrease lattice defects. It can be concluded that the Er2O3-MgO was an effective sintering additive system to prepare Si3N4 ceramics with excellent mechanical properties and high thermal conductivity through SRBSN method.
[16]
[17]
[18]
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This work was supported by the National Key Research and Development Program of China (2016YFB-0700300), National Natural Science Foundation of China (No.51572277,51702340), Shanghai Science and Technology Committee (17YF1428800, 17ZR1434800, 17dz2307000), and the State Key Laboratory of High Performance Ceramics and Superfine Microstructure of Shanghai Institute Ceramics, Chinese Academy of Sciences.
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