- Email: [email protected]
Thermal shock resistance of rare-earth doped in-situ SiAlON reinforced h-BN matrix ceramics under vacuum thermal cycling
Ceramics International 45 (2019) 20121–20127
Contents lists available at ScienceDirect
Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Thermal shock resistance of rare-earth doped in-situ SiAlON reinforced h-BN matrix ceramics under vacuum thermal cycling
T
Zhuo Tiana,∗, YuPing Yanga, Yong Wangb, Haidong Wua, Wei Liua, Shanghua Wua,∗∗ a b
School of Electromechanical Engineering, Guangdong University of Technology, Guangzhou, 510006, Guangdong, China Dongguan South China Design Innovation Institute, Dongguan, 523808, Guangdong, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Thermal shock Bending strength Thermal expansion coefficient Residual strength
In-situ SiAlON reinforced BN-matrix ceramics were prepared by hot pressing sintering, and the effects of different rare earth oxides on the thermal shock resistance of the materials were investigated. The effects of rare earth oxides on the phase composition, microstructure, bending strength, thermal properties and thermal shock resistance of the composites were studied. The results show that the phase composition and bending strength of ceramics with different rare earth oxides had no obvious change. However, the influence on the thermal expansion coefficient of the material was notable. The thermal expansion coefficient of the ceramics with CeO2 increased by 24.6% compared with Sm2O3 in the test temperature range. After 50 cycles of thermal shock at Δt = 1150 °C, the residual strength of ceramics with CeO2 was down to 157.1 MPa, decreased by 40.6% compared with the one tested in room temperature. And the Sm2O3-added ceramics reduced by 34.7%–167.1 MPa after thermal shock. The decrease of the residual strength of ceramics is mainly caused by the internal stress generated by the mismatch between the growth of quartz and SiAlON phase in the matrix and the thermal expansion coefficient of the matrix. However, no macro cracks were observed on the surface of the samples after thermal shock.
1. Introduction Hall-effect thrusters (based on the discovery by Edwin Hall) are usually referred to as Hall thrusters or Hall-current thrusters. And it has benefited for considerable theoretical and experimental researches since the 1960s [1–3]. Ordinarily, Hall thrusters are often regarded as a moderate specific impulse (1,000s-3,000s) space propulsion technology [2,3]. Possessing high-specific thrust, excellent efficiency and high reliability, Hall thruster has been widely used for various space missions, such as satellite attitude control, orbit trimming and power compensator [4–10]. And with the emergence of higher-power Hall thrusters (∼100 KW, USA, 2017), using Hall thrusters as power sources for deep space exploration vehicle is also putting on the agenda. At the same time, Hall thrusters with higher power have also entered the design stage. However, higher-power also makes the Hall thrusters work under higher temperature. For example, when the discharge input power, discharge voltage and propellant mass flow rate are respectively 200 W, 250 V and 0.94 mg/s, the temperature of inner wall is 355.4 °C [11]. When the relevant operating parameters are adjusted to 2700 W, 300 V and 10 mg/s respectively, the temperature of inner wall is 504.8 °C
∗
[12]. When the Hall thruster used in related fields, it needs to face an irregular ignition operation. And this makes the discharge channel to be subjected to cyclic thermal shock. Thus, the organization and performance stability of channel material under high temperature become one of the most important factors to ensure the normal operation of the thruster. Because of the outstanding chemical stability, unbeatable electric insulation, excellent machinability and suitable secondary electronic emission factor [8], hexagonal boron nitride (h-BN) matrix composites are considered as the best candidate for manufacturing hall thruster channels. However, the fabrication of purified high-density h-BN is extremely difficult due to its crystal structure (B-N bond and B-N six membered ring), which usually leads to the degradation of resistance of plasma erosion. In order to obtain high performance h-BN materials, hot pressing (HP) sintering is usually carried out, and sintering additives are added to improve the sinterability of h-BN materials. In recent decades, several kinds of ceramic composites containing h-BN have been explored as the material of Hall Thruster channel wall, such as BN-AlN [13,14], BN-SiO2 [15,16], BN-ZrO2 [17–19], BN-Si3N4 [20] and BN-SiAlON [20].
Corresponding author. Corresponding author. E-mail addresses: [email protected] (Z. Tian), [email protected] (S. Wu).
∗∗
https://doi.org/10.1016/j.ceramint.2019.06.277 Received 23 April 2019; Received in revised form 21 June 2019; Accepted 26 June 2019 Available online 27 June 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
Nevertheless, most of the research focuses on the mechanical properties, microstructural characterization and erosion resistance performances [21–23], while the stabilities of organization and mechanical properties under thermal cycling at high temperature are rarely reported. Therefore, it is of positive significance to carry out the research on the related aspects for its application in Hall thrusters. In the present work, BN-SiAlON materials were fabricated by hot pressing sintering. And the organization stability under high temperature and thermal shock resistance of BN-SiAlON composite materials were investigated. To improve the sinterability of BN, SiO2 was added as the liquid phase sintering aid. The SiAlON phase was obtained by insitu reaction with AlN. The phase composition and the microstructure were studied using X-ray diffraction (XRD) and scanning electron microscope (SEM), and the mechanical properties were characterized using three-point bending method after thermal cycling. 2. Experimental procedure
rate of sample surface temperature's rise and fall under different power conditions was taken as the max of 6 test results. Fig. 1 shows the experimental results of the temperature rise and fall on the surface of the sample at the max state. The ambient temperature of the experimental cabin was 20 °C. Thermal shock tests were carried out at temperature from 770 °C to 1170 °C. The temperature differences (Δt) were selected as 750 °C, 850 °C, 950 °C, 1050 °C and 1150 °C, respectively. The material was subjected to thermal cycling experiments using a vacuum induction furnace. The vacuum induction furnace first heated to the preset temperature, then moved the sample to the heated area. The holding time for a single thermal cycle was 10 min. The cooling of the ceramic sample was natural cooling at the cold end of the vacuum furnace (20 °C). The thermal cycling times were 1, 5, 10, 20, 30 and 50 times, respectively. For each condition, at least six samples were tested and investigated. 2.3. Characterization
2.1. Powder mixture and sintering The raw materials used in this work were h-BN (1.31 μm, 99.5% purity, with impurities of B3+, O2−(B2O3), Na+, K+, Ca2+, Cl-, Advanced Technology & Materials Co. Ltd., Beijing, China), fused quartz/SiO2 (0.9 μm, 98.0% purity, with impurities of Ca2+, Mg2+, Na+, Cl-, Al3+, Fe3+, Guangyu Quartz Co. Ltd., Lianyungang, China) and AlN powders (1.5 μm, 99.5% purity, with impurities of Al3+, O2− (Al2O3), Na+, Liaoning Nitride Ltd., China). In order to promote the formation of SiAlON phase, rare earth oxides CeO2, Nd2O3 and Sm2O3 were added as sintering aids, respectively. The volume ratio of h-BN: SiO2: AlN in raw materials was 66.5:28.5:5. The addition of rare earth oxides was 3 vol% of the total raw materials. The powders was added to an anhydrous ethanol medium, and milled with ZrO2 balls for 12 h in plastic bottles. The mixed slurry was then evaporated at 90 °C in the rotary evaporator to get dried powder and finally sieved through 200 mesh to obtain a composite powder. Hot pressing sintering was carried out at 1800 °C for 1 h under 20 MPa pressure in 1 atm N2 atmosphere. 2.2. Experiment method The densities of specimens were measured by the Archimedes method. The temperature calibration samples with a designed dimension of 10 mm × 30 mm × 40 mm were polished by metallographic sand paper before test, and placed 4 inches away from plasma flow source. A temperature sensor was placed on the surface of the sample to test the temperature rise and fall time of the surface and the maximum temperature under different power conditions. The maximum experimental condition of plasma source was 1.3L/s as Ar gas flow rate, 40 V for discharge voltage, 750 A for discharge current. At the above condition, the plasma velocity was above 25000 m/s, unit density was 3.68 × 1024 ions/(m2 •s), and the plasma energy was 5.36 eV/ion. The
The density of the ceramic material was measured using the Archimedes principle. The specific heat capacity of the material was measured by differential scanning calorimetry (DSC 404, Netzsch, Germany). The thermal diffusivity of the material was measured using a laser thermal conductivity meter (LFA 427, Netzsch, Germany). The relative content of the sample phase after thermal shock was calculated by Rietveld analysis using XRD data. The residual flexural strength after thermal shock was tested via the three-point bending method using rectangular bars (3 mm × 4 mm × 36 mm) in a universal testing machine (Istron-5569, Instron, USA) with a span of 30 mm. Phase analysis after thermal shock was done by X-ray diffraction (XRD; Rigaku D/Max 2200VPC, Japan). And the microstructure was examined by scanning electron microscopy (SEM; SUPRA™ 55, ZEISS Co., German). 3. Results and discussion 3.1. Phase compositions Fig. 2 shows XRD patterns of h-BN-SiAlON ceramic composites fabricated by HP with different rare earth oxides. Clearly, it can be found that the phase compositions in all composites were basically alike, and the main phase was h-BN. No other phase containing [B] was found. This indicates that chemical reaction did not take place between h-BN and other phase during sintering process. In addition to the BN phase, SiAlON phase was generated during HP sintering. And the process was accompanied by the consumption of AlN and SiO2. The chemical reactions to generate SiAlON phase can be proposed as follow: SiO2 + AlN→Si3N4+Al2O3
(1)
Si3N4 + Al2O3 + AlN→Si2Al3O7N
(2)
Si3N4 + Al2O3 + AlN→Si5AlON7
(3)
Fig. 1. Thermocouples placement and temperature test result: a) thermocouples placement; b) temperature measurements. 20122
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
Fig. 4. Thermal expansion curve of ceramics. Fig. 2. XRD patterns of sintered ceramic materials with different rare earth oxides.
Si3N4 + Al2O3 + AlN→Si3Al3O3N5
(4)
With the introduction of different rare earth elements, the main crystal phase of the generated SiAlON phase did not change. From Fig. 3, it can be seen that the main crystal structure of the generated SiAlON phase was triclinic Si2Al3O7N, hexagonal Si5AlON7 and Si3Al3O3N5 (Fig. 3(a), (b), (c)). Unlike common rod-shaped crystals, the SiAlON particles generated in this experiment were mostly equiaxed crystals and agglomerated together. The generated of equiaxed crystals is mainly caused by the following factors. As a polycrystalline material, the growth of crystals is limited by the free growth space in addition to the interfacial energy. Studies have shown that the growth of grains will stop when they meet the grains of the same size during the growth process [24,25]. Meanwhile, the existence of refractory lamellar BN also leads to the inhibition of the growth of SiAlON rods. These generated SiAlON particles can act as a reinforcing phase dispersed in the
Fig. 3. TEM images of Generated phase: a) Si2Al3O7N; b) Si5AlON7; c) Si3Al3O3N5; d) Cristobalite.
matrix to strengthen and toughen the material. In addition, the transition from amorphous silica to cristobalite was also observed in the matrix. The size of the produced cristobalite ranged from 8 nm to 15 nm, as shown in Fig. 3(d).
3.2. Thermal properties of composite ceramics The results of the thermal expansion coefficient of BN-matrix composite ceramics between room temperature and 1200 °C are shown in Fig. 4. It can be seen from Fig. 4 that the thermal expansion curve of the ceramic material can be divided into three intervals. When the temperature is between 80 °C and 200 °C, the thermal expansion test of the composite ceramic is in the low temperature stabilization stage. When the temperature is between 200 °C and 960 °C, the composite ceramic is in a stable thermal expansion stage. However, when the temperature is higher than 960 °C, the thermal expansion curve of the composite ceramic shows a rapid rise. Previous studies have shown that the amorphous silica can be partly converted into cristobalite of 8–15 nm after sintering (Fig. 3(d)). As the temperature increases, when the energy obtained by the precipitated crystals reaches the driving force that required for their growth, the crystals begin to grow. And the higher the temperature, the faster growth of the crystal. Fig. 5 shows the results of internal phase analysis of BN-SiO2 ceramic after insulation at different temperatures in N2 atmosphere. It
Fig. 5. The interior XRD patterns of BN-SiO2 ceramic after insulation at different temperatures.
20123
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
can be seen that when the temperature is lower than 900 °C, there is only h-BN diffraction peak existed. As the test temperature is further increased, SiO2 crystals are precipitated on the surface of the sample. In the thermal expansion curve, the thermal expansion coefficient of the composite ceramic starts to increase at about 960 °C, which is higher than 900 °C. This is because the heating rate of the composite ceramic is 5 °C/min when measuring the coefficient of thermal expansion. When the surface phase of materials is analyzed at different temperatures, the heating rate is 75∼100 °C/s. The faster heating rate, the faster temperature changes of the sample surface. And the crystallization behavior will have a certain lag. This is because the thermal conductivity of ceramic materials decreases with the increase of ambient temperature (Sup.1). Therefore, as the temperature difference of the thermal vibration increases, the temperature homogenization time of the sample becomes longer, which means there is a temperature gradient inside the sample. And this results in the delayed growth of the internal crystal grains relative to those on the surface of the sample. Furthermore, the existence of the temperature gradient, on the one hand, leads to internal stress caused by the difference in thermal expansion coefficients between the phases on the same temperature surface. On the other hand, it also causes internal stress due to the difference in thermal expansion coefficient between the inside and outside of the sample. The accumulation of these two factors adversely affects the residual strength of the material. In addition, the relative intensity of SiAlON diffraction peak on the sample surface was increased (Δt = 1150 °C, Fig. 6) while this phenomenon was not found on the surface of the sample at Δt = 1050 °C. Since the content of AlN was much lower than that of silicon oxide, there was no residual AlN in the sintering process [26]. And the reaction temperature between the two phases was higher than 1600 °C. Therefore, in the XRD pattern after 1170 °C heat preservation, the appearance of the SiAlON diffraction peak was caused by the growth of the nanocrystal grains generated under high temperature. In short, in
Table 1 Relative content of each phase inside the sample after thermal shock (Δt = 1150 °C)a. Rare earth oxides
CeO2 Nd2O3 Sm2O3
Phase h-BN /wt%
SiAlON /wt%
SiO2b /wt%
72.23 70.74 70.76
8.37 8.02 7.89
19.40 21.24 21.35
a As the rare earth oxides have not been detected, the relative content of rare earth oxides in the products has not been calculated. b Including amorphous phase and crystalline phase.
Part III region, the rapid increase of the thermal expansion curve was caused by the interaction of the nano-quartz and the nano-SiAlON phase.The relative content of each phase on the surface of the sample after thermal shock is shown in Table 1. At the same time, it can be found that the addition of rare earth elements has a significant effect on the thermal expansion properties of the material. The thermal expansion coefficient of CeO2-added samples is 24.6% higher than that of Sm2O3-added samples. Studies have shown that the addition of rare earth oxides can promote the nucleation and growth of SiAlON phase [27]. The smaller the ion radius of rare earth elements, the lower the diffusion resistance in the liquid phase, the more obvious the promoting effect on the formation of SiAlON. Among the rare earth elements selected herein, Ce4+(87 p.m.) has the smallest ionic radius. Therefore, it is easiest to promote the nucleation and growth of SiAlON phase in the sintering process to get the highest relative content of SiAlON in the matrix. As the thermal expansion coefficient of SiAlON(3.0∼4.5 × 10-6/K) is higher than that of h-BN (0.7 × 10-6/K (⊥), 2.7 × 10-6/K (∥) and SiO2(0.54 × 10-6/K), the
Fig. 6. XRD analysis results of ceramics after thermal shock at different temperature: a) CeO2; b) Nd2O3; c) Sm2O3. 20124
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
Fig. 7. SEM images of sample surface after thermal shock (Δt = 750 °C): a) CeO2; b) Nd2O3; c) Sm2O3.
ceramic material with CeO2 has the highest thermal expansion coefficient. However, due to the relatively low content of SiAlON phase in the matrix, the coefficient of thermal expansion of the material is still at a low level. 3.3. Surface morphology after thermal shock The surface morphology of ceramics with different rare earth oxides after thermal shock are shown Fig. 7 and Fig. 8. No macro-cracks or micro-cracks were found on the surface of samples after thermal shock. This indicates that although there is a mismatch of thermal expansion coefficient between SiAlON (3.0∼4.5 × 10-6/K) and matrix materials (h-BN (0.7 × 10-6/K (⊥), 2.7 × 10-6/K (∥)) [8] and SiO2 (0.54 × 10-6/ K)). However, due to the presence of 1.7 (CeO2) to 2.0% (Sm2O3) pores in the matrix with a size between 40 and 480 nm (Sup.2), the compressive stress on the matrix material caused by the expansion of the SiAlON phase during thermal shock is partially released through the pores. Thus, the macro or micro cracks inside or on the surface of the ceramics caused by the excessive concentration of internal stress can be avoided. At the same time, it can be found in Fig. 9 that amorphous substances are formed on the surface of the sample. The analysis shows that the amorphous phase is mainly SiO2. This is because the SiO2 on the surface of the sample changes from solid to liquid phase when the holding temperature reaches the transition point of SiO2. Under the action of molecular vibration and surface tension, viscous silicon oxide gradually covered the surface of the sample. It can be seen that when Δt = 1150 °C, the surface of the sample after thermal shock is smoother than that of Δt = 750 °C. Meanwhile, as Δt increases, white particles gradually precipitate on the surface of the sample. The elemental composition is shown in Fig. 9. It is well known that crystal preferentially nucleate at defects or impurities. And the presence of impurities also causes a decrease in the phase transition temperature of the fused silica. Therefore, as the temperature increases, when the fused silica changes from a solid state to a viscous state, the impurities are simultaneously incorporated into the viscous fused silica. In the process of agglomeration of fused quartz in viscous state, some impurities will be encapsulated by fused quartz in viscous state, and impurity elements will be transferred at the same time. Then it precipitates from the liquid phase during the cooling process.
Since these impurities are introduced from the raw materials in the preparation stage, which is unavoidable. Therefore, precipitation of impurity elements in the liquid phase is inevitable on the surface of the sample. However, the bonding force between the glass phase and the matrix is very weak, and cracks are easily generated under the irradiation of secondary electrons, as shown in Fig. 9(b). It can be seen that there may be a weak bond between the precipitated glass phase and the matrix. The joint can be regarded as a defective area on the surface of the sample, where cracks are readily generated. When subjected to an applied load, the crack is preferentially generated and expanded in this region, which leads to the damage of the sample. The precipitation of the liquid phase on the surface of the sample will reduce the relative content of the fused silica inside the sample, weakening the bonding strength between the h-BN particles inside the sample, and adversely affecting the properties of the material after thermal shock. 3.4. Residual strength after thermal shock The residual strength of composites with different rare earth oxides after thermal shock are shown in Fig. 10. And Fig. 11 is the flexural strength of ceramics before thermal shock. Compared with the strength of 245.6 MPa without adding rare earth oxide, the flexural strength of the material increased by 7.7%–264.5 MPa when added 5 vol% CeO2. Also, when 5 vol% Sm2O3 added, the flexural strength of the material only increased by 4.2%–255.9 MPa. This shows that the introduction of rare earth elements does not significantly improve the flexural strength of the materials. It can be seen that the residual strength gradually decreased with the increase of thermal shock temperature difference and cycle times in the test temperature range. When Δt was lower than 1050 °C, the residual strength after thermal shock did not decrease significantly. For the ceramic material to which CeO2 was added, after 50 cycles, the residual strength was 211.0 MPa, decreased by 20.2%. It can be seen that the material exhibited good thermal shock resistance when Δt was lower than 1050 °C. When Δt was higher than 1050 °C, the mechanical properties of the material after thermal shock were rapidly reduced. When Δt up to 1150 °C, after 50 thermal cycles, the residual strength of CeO2-added ceramic was 157.1 MPa, reduced by 40.6%, which was the largest decline of the material residual strength. According to the relevant impact resistance factor, in the previous
Fig. 8. SEM images of sample surface after thermal shock (Δt = 1150 °C): a) CeO2; b) Nd2O3; c) Sm2O3. 20125
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
Fig. 9. Precipitate phase of ceramics on sample surface: a) Initial picture; b) after 5s; c) Energy spectrum of white particles.
undamaged material, the failure condition during the thermal shock was obtained by setting the maximum allowable safety stress as the bending strength. The thermal shock resistance parameter R is used to indicate the conditions of crack initiation failure. For the case of an infinite plate, it is symmetrically heated or cooled with a constant heat transfer coefficient, R is defined as:
R=
σf (1 − ν ) Eα
(5)
where σf is the bending strength, α is the thermal expansion coefficient, E is the Young modulus and v is the Poisson ratio. For ceramics with the same matrix, assuming that v and E are the same, the constant k can be defined as:
k=
(1 − ν ) E
(6) Fig. 11. Bending strength of ceramics at room temperature.
Then equation (5) can be simplified as:
R=k
σf α
(7)
The lower the value of R, the better thermal shock resistance of the samples. Referring to the flexural strength at the room temperature in
Fig. 10, it can be seen that the values of strength are very close. It is well known that the lower the coefficient of thermal expansion, the better the thermal shock resistance of the material. Thus, the ceramic system
Fig. 10. Residual strength of ceramics after thermal shock. 20126
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
with Sm2O3 has the best thermal shock resistance under this test condition. Furthermore, for the material system studied in this paper, the thermal shock resistance evaluation of the specimens still obeys the relevant rules even though there are micro-crack sources on the surface of the specimens. To summarize, the thermal shock process of ceramic samples in this paper is equivalent to “high temperature heat treatment”. When the ambient temperature is higher than 1000 °C, the fused quartz transformed into cristobalite, accompanying by the growth of nano-cristobalite and the formation of nano-sialon grains. Therefore, the thermal expansion curve of the material increases rapidly under the combined action of the two factors. During the thermal shock process, as the difference of thermal expansion coefficient between matrix material and grown quartz and SiAlON phase, the growing grains will produce compressive stress around them and then cause internal stress. The higher the internal stress of the material with high thermal expansion coefficient, the worse the thermal shock resistance of the material, while the residual strength of the material decreases more obviously after the thermal shock. Therefore, under the combined action of the internal and external factors of the samples, the residual strength of the ceramic decreases sharply with the increase of the thermal shock temperature difference and the cycle index. 4. Conclusions In this study, the effect of rare earth oxides on the thermal shock resistance of in-situ SiAlON phase-enhanced BN-matrix ceramics was investigated. The effect of microstructure and thermal properties of the composites have been studied in detail. Several key conclusions can be drawn as follows. 1) The introduction of rare earth oxides has little effect on the improvement of the flexural strength of the material. When CeO2 was added, the flexural strength of the material was 264.5 MPa, which was only increased by 7.7% compared to the unadded one of 245.6 MPa. 2) The influence of rare earth oxides on thermal expansion coefficient of the material is more obvious. The thermal expansion coefficient of the CeO2-added was increased by 24.6%, which was higher than that with Sm2O3. 3) Different rare earth oxides have significantly different effects on the thermal shock resistance of materials. After 50 cycles at Δt = 1150 °C, the ceramic material added with CeO2 had a residual strength of 157.1 MPa, which was reduced by 40.6% % compared with the room temperature strength. Under the same conditions, the residual strength of the sample added Sm2O3 decreased by 34.7%–167.1 MPa. However, no cracks due to thermal expansion mismatch were found on the surface after thermal shock. 4) The decrease in the residual strength of the ceramic material is mainly caused by the internal stress generated from the mismatch of the thermal expansion coefficient between the new generated and grown phases in the matrix material. Acknowledgements The authors would like to thank Laboratory for Materials in Extreme Environments in UCLA lead by prof. Nasr Ghoniem for testing the sputtering temperature of the sample. And this work was financially supported by Local Innovative and Research Team Project of Guangdong Pearl River Talents Program (Grant No.2017BT01C169), Science and Technology Major Project of Guangdong Province (Grant No.2016B090915002), Science and Technology Major Project of Guangdong Province (Grant No.2017B090911011), the China Postdoctoral Science Foundation (CPSF, Grant No. 2018M633013), Special Fund for the Cultivation of
Science and Technology Innovation of Guangdong University Students in 2018(Grant No. PDJHB0157). Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.ceramint.2019.06.277. References [1] M. Martinez-Sanchez, J.E. Pollard, Spacecraft electric propulsion-an over view, J. Propuls. Power 4 (1998) 688–699. [2] S. Mazouffre, Electric propulsion for satellites and spacecraft: established technologies and novel approaches, Plasma Sources Sci. Technol. 25 (3) (2016) 1–27. [3] J.P. Boeuf, Tutorial: physics and modeling of Hall thrusters, J. Appl. Phys. 121 (1) (2017) 1–23. [4] Y.H. Chiu, L.A. Brad, W. Kip, A.D. Rainer, George F. Karabadzhak, Passive Optical Diagnostic of Xe-Propelled Hall thrusters. I. Emission cross sections, J. Appl. Phys. 99 (2006) 1–10. [5] B.M. Nathan, G. Nicolas, A.C. Mark, Linear geometry Hall thruster with boron nitride and diamond walls, In: 27th Inter. Electr. Propul. Conf., Pasadena, CA, USA, IEPC-01-39. [6] N. Gascon, M. Dudeck, S. Barral, Wall material effects in stationary plasma thrusters. I. Parametric studies of an SPT-100, Phys. Plasmas 10 (2003) 4123–4136. [7] S. Barral, K. Makowski, Z. Peradzynski, Wall material effects in stationary plasma thrusters. II. Near-wall and in-wall conductivity, Phys. Plasmas 10 (2003) 4137–4152. [8] M. Britton, D. Waters, R. Messer, E. Sechkar, Sputtering Erosion Measurement on Boron Nitride as a Hall Thruster Material, Cleveland, NASA/TM-2002-211837. [9] H. Tahara, K. Imanaka, S. Yuge, Effects of channel wall material on thruster performance and plasma characteristics of Hall-effect thrusters, Vacuum 80 (2006) 1216–1222. [10] Y. Raitses, D. Staack, M. Keidar, N.J. Fisch, Electron-wall interaction in Hall thrusters, Phys. Plasmas 12 (5) (2005) 057104. [11] A.M. Rafael, D. Hoang, L.R. Walker Mitchell, Power deposition into the discharge channel of a Hall effect thruster, J. Propuls. Power 30 (1) (2014) 209–220. [12] A.M. Bohnert, Thermal characterization of a hall effect thruster, Thesis for the Degree of Master of Science in Astronautical Engineering (2008) 37–50. [13] T. Kusunose, N. Sakayanagi, T. Sekino, Y. Ando, Y. Ando, Fabrication and characterization of aluminum nitride/boron nitride nanocomposites by carbothermal reduction and nitridation of aluminum borate powders, J. Nanosci. Nanotechnol. 8 (2008) 5846–5853. [14] G.J. Zhang, H. Kita, N. Kondo, T. Ohji, Reactive hot-pressed aluminaboron nitride composites with Y2O3 sintering additive, J. Am. Ceram. Soc. 88 (8) (2005) 2246–2248. [15] H.Z. Zhai, H.N. Cai, X.Z. Yang, J.B. Li, G.F. Guo, C.B. Cao, Preparation and properties of BN-SiO2 composite ceramics, Key Eng. Mater. 336–338 (2007) 1426–1428. [16] Y. Garnier, V. Viel, J.F. Roussel, J. Bernard, Low-energy xenon ion sputtering of ceramics investigated for stationary plasma thrusters, J. Vac. Sci. Technol. A 17 (6) (1999) 3246–3255. [17] X.H. Zhang, R.B. Zhang, G.Q. Chen, W.B. Han, Microstructure, mechanical properties and thermal shock resistance of hot-pressed ZrO2(3Y)- BN composites, Mater. Sci. Eng. A-Struct. 497 (1–2) (2008) 195–199. [18] Y.L. Li, J.X. Zhang, G.J. Qiao, Z.H. Jin, Fabrication and properties of machinable 3Y-ZrO2/BN Nanocomposites, Mater. Sci. Eng. A-Struct. 397 (1–2) (2005) 35–40. [19] X.M. Duan, D.C. Jia, Q.C. Meng, Z.H. Yang, Y. Yu, Y. Zhou, D.R. Yu, Y.J. Ding, Study on the plasma erosion resistance of ZrO2p(3Y)/BN-SiO2 composite ceramics, Composites Part B 46 (2013) 130–134. [20] Y.L. Li, R.X. Li, J.X. Zhang, Enhanced mechanical properties of machinable Si3N4/ BN composites by spark plasma sintering, Mater. Sci. Eng. A-Struct. 483–484 (2008) 207–210. [21] M. Britton, D. Waters, R. Messer, E. Sechkar, Sputtering Erosion Measurement on Boron Nitride as a Hall Thruster Material. Cleveland, NASA, TM-2002-211837. [22] M. Tartz, T. Heyn, C. Bundesmann, H. Neumann, Measuring sputter yields of ceramic materials, 31st Inter. Electr. Propul. Conf. IEPC, Ann Arbor, Michigan, USA, 2009, p. 240. [23] S.A. Khartov, A.B. Nadiradze, I.I. Shkarban, Y.V. Zikeeva, SPT's high lifetime -some problems of solution, 29th Inter. Electr. Propul. Conf. IEPC, Princeton, USA, 2005, p. 62. [24] Z. Tian, X.M. Duan, Z.H. Yang, S.Q. Ye, D.C. Jia, Y. Zhou, Microstructure and erosion resistance of in-situ SiAlON reinforced BN-SiO2 composite ceramics, J. Wuhan Univ. Technol.-Materials Sci. Ed. 31 (2) (2016) 315–320. [25] X.M. Hou, C.S. Yue, A.K. Singh, M. Zhang, Morphological development and oxidation of elongated β-SiAlON material, Corros. Sci. 53 (6) (2011) 2051–2057. [26] Zhuo Tian, Dechang Jia, Duam Xiaoming, Zhihua Yang, Yu Zhou, Effects of AlN content on phase composition, microstructure and mechanical properties of BNbased composite ceramics, J. Chin. Ceram. Soc. 41 (12) (2013) 1603–1608. [27] M. Zeuner, S. Pagano, W. Schnick, Nitridosilicates and oxonitridosilicates: from ceramic materials to structural and functional diversity, Angew. Chem. Int. Ed. 50 (2011) 7754–7775.
20127
Contents lists available at ScienceDirect
Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Thermal shock resistance of rare-earth doped in-situ SiAlON reinforced h-BN matrix ceramics under vacuum thermal cycling
T
Zhuo Tiana,∗, YuPing Yanga, Yong Wangb, Haidong Wua, Wei Liua, Shanghua Wua,∗∗ a b
School of Electromechanical Engineering, Guangdong University of Technology, Guangzhou, 510006, Guangdong, China Dongguan South China Design Innovation Institute, Dongguan, 523808, Guangdong, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Thermal shock Bending strength Thermal expansion coefficient Residual strength
In-situ SiAlON reinforced BN-matrix ceramics were prepared by hot pressing sintering, and the effects of different rare earth oxides on the thermal shock resistance of the materials were investigated. The effects of rare earth oxides on the phase composition, microstructure, bending strength, thermal properties and thermal shock resistance of the composites were studied. The results show that the phase composition and bending strength of ceramics with different rare earth oxides had no obvious change. However, the influence on the thermal expansion coefficient of the material was notable. The thermal expansion coefficient of the ceramics with CeO2 increased by 24.6% compared with Sm2O3 in the test temperature range. After 50 cycles of thermal shock at Δt = 1150 °C, the residual strength of ceramics with CeO2 was down to 157.1 MPa, decreased by 40.6% compared with the one tested in room temperature. And the Sm2O3-added ceramics reduced by 34.7%–167.1 MPa after thermal shock. The decrease of the residual strength of ceramics is mainly caused by the internal stress generated by the mismatch between the growth of quartz and SiAlON phase in the matrix and the thermal expansion coefficient of the matrix. However, no macro cracks were observed on the surface of the samples after thermal shock.
1. Introduction Hall-effect thrusters (based on the discovery by Edwin Hall) are usually referred to as Hall thrusters or Hall-current thrusters. And it has benefited for considerable theoretical and experimental researches since the 1960s [1–3]. Ordinarily, Hall thrusters are often regarded as a moderate specific impulse (1,000s-3,000s) space propulsion technology [2,3]. Possessing high-specific thrust, excellent efficiency and high reliability, Hall thruster has been widely used for various space missions, such as satellite attitude control, orbit trimming and power compensator [4–10]. And with the emergence of higher-power Hall thrusters (∼100 KW, USA, 2017), using Hall thrusters as power sources for deep space exploration vehicle is also putting on the agenda. At the same time, Hall thrusters with higher power have also entered the design stage. However, higher-power also makes the Hall thrusters work under higher temperature. For example, when the discharge input power, discharge voltage and propellant mass flow rate are respectively 200 W, 250 V and 0.94 mg/s, the temperature of inner wall is 355.4 °C [11]. When the relevant operating parameters are adjusted to 2700 W, 300 V and 10 mg/s respectively, the temperature of inner wall is 504.8 °C
∗
[12]. When the Hall thruster used in related fields, it needs to face an irregular ignition operation. And this makes the discharge channel to be subjected to cyclic thermal shock. Thus, the organization and performance stability of channel material under high temperature become one of the most important factors to ensure the normal operation of the thruster. Because of the outstanding chemical stability, unbeatable electric insulation, excellent machinability and suitable secondary electronic emission factor [8], hexagonal boron nitride (h-BN) matrix composites are considered as the best candidate for manufacturing hall thruster channels. However, the fabrication of purified high-density h-BN is extremely difficult due to its crystal structure (B-N bond and B-N six membered ring), which usually leads to the degradation of resistance of plasma erosion. In order to obtain high performance h-BN materials, hot pressing (HP) sintering is usually carried out, and sintering additives are added to improve the sinterability of h-BN materials. In recent decades, several kinds of ceramic composites containing h-BN have been explored as the material of Hall Thruster channel wall, such as BN-AlN [13,14], BN-SiO2 [15,16], BN-ZrO2 [17–19], BN-Si3N4 [20] and BN-SiAlON [20].
Corresponding author. Corresponding author. E-mail addresses: [email protected] (Z. Tian), [email protected] (S. Wu).
∗∗
https://doi.org/10.1016/j.ceramint.2019.06.277 Received 23 April 2019; Received in revised form 21 June 2019; Accepted 26 June 2019 Available online 27 June 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
Nevertheless, most of the research focuses on the mechanical properties, microstructural characterization and erosion resistance performances [21–23], while the stabilities of organization and mechanical properties under thermal cycling at high temperature are rarely reported. Therefore, it is of positive significance to carry out the research on the related aspects for its application in Hall thrusters. In the present work, BN-SiAlON materials were fabricated by hot pressing sintering. And the organization stability under high temperature and thermal shock resistance of BN-SiAlON composite materials were investigated. To improve the sinterability of BN, SiO2 was added as the liquid phase sintering aid. The SiAlON phase was obtained by insitu reaction with AlN. The phase composition and the microstructure were studied using X-ray diffraction (XRD) and scanning electron microscope (SEM), and the mechanical properties were characterized using three-point bending method after thermal cycling. 2. Experimental procedure
rate of sample surface temperature's rise and fall under different power conditions was taken as the max of 6 test results. Fig. 1 shows the experimental results of the temperature rise and fall on the surface of the sample at the max state. The ambient temperature of the experimental cabin was 20 °C. Thermal shock tests were carried out at temperature from 770 °C to 1170 °C. The temperature differences (Δt) were selected as 750 °C, 850 °C, 950 °C, 1050 °C and 1150 °C, respectively. The material was subjected to thermal cycling experiments using a vacuum induction furnace. The vacuum induction furnace first heated to the preset temperature, then moved the sample to the heated area. The holding time for a single thermal cycle was 10 min. The cooling of the ceramic sample was natural cooling at the cold end of the vacuum furnace (20 °C). The thermal cycling times were 1, 5, 10, 20, 30 and 50 times, respectively. For each condition, at least six samples were tested and investigated. 2.3. Characterization
2.1. Powder mixture and sintering The raw materials used in this work were h-BN (1.31 μm, 99.5% purity, with impurities of B3+, O2−(B2O3), Na+, K+, Ca2+, Cl-, Advanced Technology & Materials Co. Ltd., Beijing, China), fused quartz/SiO2 (0.9 μm, 98.0% purity, with impurities of Ca2+, Mg2+, Na+, Cl-, Al3+, Fe3+, Guangyu Quartz Co. Ltd., Lianyungang, China) and AlN powders (1.5 μm, 99.5% purity, with impurities of Al3+, O2− (Al2O3), Na+, Liaoning Nitride Ltd., China). In order to promote the formation of SiAlON phase, rare earth oxides CeO2, Nd2O3 and Sm2O3 were added as sintering aids, respectively. The volume ratio of h-BN: SiO2: AlN in raw materials was 66.5:28.5:5. The addition of rare earth oxides was 3 vol% of the total raw materials. The powders was added to an anhydrous ethanol medium, and milled with ZrO2 balls for 12 h in plastic bottles. The mixed slurry was then evaporated at 90 °C in the rotary evaporator to get dried powder and finally sieved through 200 mesh to obtain a composite powder. Hot pressing sintering was carried out at 1800 °C for 1 h under 20 MPa pressure in 1 atm N2 atmosphere. 2.2. Experiment method The densities of specimens were measured by the Archimedes method. The temperature calibration samples with a designed dimension of 10 mm × 30 mm × 40 mm were polished by metallographic sand paper before test, and placed 4 inches away from plasma flow source. A temperature sensor was placed on the surface of the sample to test the temperature rise and fall time of the surface and the maximum temperature under different power conditions. The maximum experimental condition of plasma source was 1.3L/s as Ar gas flow rate, 40 V for discharge voltage, 750 A for discharge current. At the above condition, the plasma velocity was above 25000 m/s, unit density was 3.68 × 1024 ions/(m2 •s), and the plasma energy was 5.36 eV/ion. The
The density of the ceramic material was measured using the Archimedes principle. The specific heat capacity of the material was measured by differential scanning calorimetry (DSC 404, Netzsch, Germany). The thermal diffusivity of the material was measured using a laser thermal conductivity meter (LFA 427, Netzsch, Germany). The relative content of the sample phase after thermal shock was calculated by Rietveld analysis using XRD data. The residual flexural strength after thermal shock was tested via the three-point bending method using rectangular bars (3 mm × 4 mm × 36 mm) in a universal testing machine (Istron-5569, Instron, USA) with a span of 30 mm. Phase analysis after thermal shock was done by X-ray diffraction (XRD; Rigaku D/Max 2200VPC, Japan). And the microstructure was examined by scanning electron microscopy (SEM; SUPRA™ 55, ZEISS Co., German). 3. Results and discussion 3.1. Phase compositions Fig. 2 shows XRD patterns of h-BN-SiAlON ceramic composites fabricated by HP with different rare earth oxides. Clearly, it can be found that the phase compositions in all composites were basically alike, and the main phase was h-BN. No other phase containing [B] was found. This indicates that chemical reaction did not take place between h-BN and other phase during sintering process. In addition to the BN phase, SiAlON phase was generated during HP sintering. And the process was accompanied by the consumption of AlN and SiO2. The chemical reactions to generate SiAlON phase can be proposed as follow: SiO2 + AlN→Si3N4+Al2O3
(1)
Si3N4 + Al2O3 + AlN→Si2Al3O7N
(2)
Si3N4 + Al2O3 + AlN→Si5AlON7
(3)
Fig. 1. Thermocouples placement and temperature test result: a) thermocouples placement; b) temperature measurements. 20122
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
Fig. 4. Thermal expansion curve of ceramics. Fig. 2. XRD patterns of sintered ceramic materials with different rare earth oxides.
Si3N4 + Al2O3 + AlN→Si3Al3O3N5
(4)
With the introduction of different rare earth elements, the main crystal phase of the generated SiAlON phase did not change. From Fig. 3, it can be seen that the main crystal structure of the generated SiAlON phase was triclinic Si2Al3O7N, hexagonal Si5AlON7 and Si3Al3O3N5 (Fig. 3(a), (b), (c)). Unlike common rod-shaped crystals, the SiAlON particles generated in this experiment were mostly equiaxed crystals and agglomerated together. The generated of equiaxed crystals is mainly caused by the following factors. As a polycrystalline material, the growth of crystals is limited by the free growth space in addition to the interfacial energy. Studies have shown that the growth of grains will stop when they meet the grains of the same size during the growth process [24,25]. Meanwhile, the existence of refractory lamellar BN also leads to the inhibition of the growth of SiAlON rods. These generated SiAlON particles can act as a reinforcing phase dispersed in the
Fig. 3. TEM images of Generated phase: a) Si2Al3O7N; b) Si5AlON7; c) Si3Al3O3N5; d) Cristobalite.
matrix to strengthen and toughen the material. In addition, the transition from amorphous silica to cristobalite was also observed in the matrix. The size of the produced cristobalite ranged from 8 nm to 15 nm, as shown in Fig. 3(d).
3.2. Thermal properties of composite ceramics The results of the thermal expansion coefficient of BN-matrix composite ceramics between room temperature and 1200 °C are shown in Fig. 4. It can be seen from Fig. 4 that the thermal expansion curve of the ceramic material can be divided into three intervals. When the temperature is between 80 °C and 200 °C, the thermal expansion test of the composite ceramic is in the low temperature stabilization stage. When the temperature is between 200 °C and 960 °C, the composite ceramic is in a stable thermal expansion stage. However, when the temperature is higher than 960 °C, the thermal expansion curve of the composite ceramic shows a rapid rise. Previous studies have shown that the amorphous silica can be partly converted into cristobalite of 8–15 nm after sintering (Fig. 3(d)). As the temperature increases, when the energy obtained by the precipitated crystals reaches the driving force that required for their growth, the crystals begin to grow. And the higher the temperature, the faster growth of the crystal. Fig. 5 shows the results of internal phase analysis of BN-SiO2 ceramic after insulation at different temperatures in N2 atmosphere. It
Fig. 5. The interior XRD patterns of BN-SiO2 ceramic after insulation at different temperatures.
20123
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
can be seen that when the temperature is lower than 900 °C, there is only h-BN diffraction peak existed. As the test temperature is further increased, SiO2 crystals are precipitated on the surface of the sample. In the thermal expansion curve, the thermal expansion coefficient of the composite ceramic starts to increase at about 960 °C, which is higher than 900 °C. This is because the heating rate of the composite ceramic is 5 °C/min when measuring the coefficient of thermal expansion. When the surface phase of materials is analyzed at different temperatures, the heating rate is 75∼100 °C/s. The faster heating rate, the faster temperature changes of the sample surface. And the crystallization behavior will have a certain lag. This is because the thermal conductivity of ceramic materials decreases with the increase of ambient temperature (Sup.1). Therefore, as the temperature difference of the thermal vibration increases, the temperature homogenization time of the sample becomes longer, which means there is a temperature gradient inside the sample. And this results in the delayed growth of the internal crystal grains relative to those on the surface of the sample. Furthermore, the existence of the temperature gradient, on the one hand, leads to internal stress caused by the difference in thermal expansion coefficients between the phases on the same temperature surface. On the other hand, it also causes internal stress due to the difference in thermal expansion coefficient between the inside and outside of the sample. The accumulation of these two factors adversely affects the residual strength of the material. In addition, the relative intensity of SiAlON diffraction peak on the sample surface was increased (Δt = 1150 °C, Fig. 6) while this phenomenon was not found on the surface of the sample at Δt = 1050 °C. Since the content of AlN was much lower than that of silicon oxide, there was no residual AlN in the sintering process [26]. And the reaction temperature between the two phases was higher than 1600 °C. Therefore, in the XRD pattern after 1170 °C heat preservation, the appearance of the SiAlON diffraction peak was caused by the growth of the nanocrystal grains generated under high temperature. In short, in
Table 1 Relative content of each phase inside the sample after thermal shock (Δt = 1150 °C)a. Rare earth oxides
CeO2 Nd2O3 Sm2O3
Phase h-BN /wt%
SiAlON /wt%
SiO2b /wt%
72.23 70.74 70.76
8.37 8.02 7.89
19.40 21.24 21.35
a As the rare earth oxides have not been detected, the relative content of rare earth oxides in the products has not been calculated. b Including amorphous phase and crystalline phase.
Part III region, the rapid increase of the thermal expansion curve was caused by the interaction of the nano-quartz and the nano-SiAlON phase.The relative content of each phase on the surface of the sample after thermal shock is shown in Table 1. At the same time, it can be found that the addition of rare earth elements has a significant effect on the thermal expansion properties of the material. The thermal expansion coefficient of CeO2-added samples is 24.6% higher than that of Sm2O3-added samples. Studies have shown that the addition of rare earth oxides can promote the nucleation and growth of SiAlON phase [27]. The smaller the ion radius of rare earth elements, the lower the diffusion resistance in the liquid phase, the more obvious the promoting effect on the formation of SiAlON. Among the rare earth elements selected herein, Ce4+(87 p.m.) has the smallest ionic radius. Therefore, it is easiest to promote the nucleation and growth of SiAlON phase in the sintering process to get the highest relative content of SiAlON in the matrix. As the thermal expansion coefficient of SiAlON(3.0∼4.5 × 10-6/K) is higher than that of h-BN (0.7 × 10-6/K (⊥), 2.7 × 10-6/K (∥) and SiO2(0.54 × 10-6/K), the
Fig. 6. XRD analysis results of ceramics after thermal shock at different temperature: a) CeO2; b) Nd2O3; c) Sm2O3. 20124
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
Fig. 7. SEM images of sample surface after thermal shock (Δt = 750 °C): a) CeO2; b) Nd2O3; c) Sm2O3.
ceramic material with CeO2 has the highest thermal expansion coefficient. However, due to the relatively low content of SiAlON phase in the matrix, the coefficient of thermal expansion of the material is still at a low level. 3.3. Surface morphology after thermal shock The surface morphology of ceramics with different rare earth oxides after thermal shock are shown Fig. 7 and Fig. 8. No macro-cracks or micro-cracks were found on the surface of samples after thermal shock. This indicates that although there is a mismatch of thermal expansion coefficient between SiAlON (3.0∼4.5 × 10-6/K) and matrix materials (h-BN (0.7 × 10-6/K (⊥), 2.7 × 10-6/K (∥)) [8] and SiO2 (0.54 × 10-6/ K)). However, due to the presence of 1.7 (CeO2) to 2.0% (Sm2O3) pores in the matrix with a size between 40 and 480 nm (Sup.2), the compressive stress on the matrix material caused by the expansion of the SiAlON phase during thermal shock is partially released through the pores. Thus, the macro or micro cracks inside or on the surface of the ceramics caused by the excessive concentration of internal stress can be avoided. At the same time, it can be found in Fig. 9 that amorphous substances are formed on the surface of the sample. The analysis shows that the amorphous phase is mainly SiO2. This is because the SiO2 on the surface of the sample changes from solid to liquid phase when the holding temperature reaches the transition point of SiO2. Under the action of molecular vibration and surface tension, viscous silicon oxide gradually covered the surface of the sample. It can be seen that when Δt = 1150 °C, the surface of the sample after thermal shock is smoother than that of Δt = 750 °C. Meanwhile, as Δt increases, white particles gradually precipitate on the surface of the sample. The elemental composition is shown in Fig. 9. It is well known that crystal preferentially nucleate at defects or impurities. And the presence of impurities also causes a decrease in the phase transition temperature of the fused silica. Therefore, as the temperature increases, when the fused silica changes from a solid state to a viscous state, the impurities are simultaneously incorporated into the viscous fused silica. In the process of agglomeration of fused quartz in viscous state, some impurities will be encapsulated by fused quartz in viscous state, and impurity elements will be transferred at the same time. Then it precipitates from the liquid phase during the cooling process.
Since these impurities are introduced from the raw materials in the preparation stage, which is unavoidable. Therefore, precipitation of impurity elements in the liquid phase is inevitable on the surface of the sample. However, the bonding force between the glass phase and the matrix is very weak, and cracks are easily generated under the irradiation of secondary electrons, as shown in Fig. 9(b). It can be seen that there may be a weak bond between the precipitated glass phase and the matrix. The joint can be regarded as a defective area on the surface of the sample, where cracks are readily generated. When subjected to an applied load, the crack is preferentially generated and expanded in this region, which leads to the damage of the sample. The precipitation of the liquid phase on the surface of the sample will reduce the relative content of the fused silica inside the sample, weakening the bonding strength between the h-BN particles inside the sample, and adversely affecting the properties of the material after thermal shock. 3.4. Residual strength after thermal shock The residual strength of composites with different rare earth oxides after thermal shock are shown in Fig. 10. And Fig. 11 is the flexural strength of ceramics before thermal shock. Compared with the strength of 245.6 MPa without adding rare earth oxide, the flexural strength of the material increased by 7.7%–264.5 MPa when added 5 vol% CeO2. Also, when 5 vol% Sm2O3 added, the flexural strength of the material only increased by 4.2%–255.9 MPa. This shows that the introduction of rare earth elements does not significantly improve the flexural strength of the materials. It can be seen that the residual strength gradually decreased with the increase of thermal shock temperature difference and cycle times in the test temperature range. When Δt was lower than 1050 °C, the residual strength after thermal shock did not decrease significantly. For the ceramic material to which CeO2 was added, after 50 cycles, the residual strength was 211.0 MPa, decreased by 20.2%. It can be seen that the material exhibited good thermal shock resistance when Δt was lower than 1050 °C. When Δt was higher than 1050 °C, the mechanical properties of the material after thermal shock were rapidly reduced. When Δt up to 1150 °C, after 50 thermal cycles, the residual strength of CeO2-added ceramic was 157.1 MPa, reduced by 40.6%, which was the largest decline of the material residual strength. According to the relevant impact resistance factor, in the previous
Fig. 8. SEM images of sample surface after thermal shock (Δt = 1150 °C): a) CeO2; b) Nd2O3; c) Sm2O3. 20125
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
Fig. 9. Precipitate phase of ceramics on sample surface: a) Initial picture; b) after 5s; c) Energy spectrum of white particles.
undamaged material, the failure condition during the thermal shock was obtained by setting the maximum allowable safety stress as the bending strength. The thermal shock resistance parameter R is used to indicate the conditions of crack initiation failure. For the case of an infinite plate, it is symmetrically heated or cooled with a constant heat transfer coefficient, R is defined as:
R=
σf (1 − ν ) Eα
(5)
where σf is the bending strength, α is the thermal expansion coefficient, E is the Young modulus and v is the Poisson ratio. For ceramics with the same matrix, assuming that v and E are the same, the constant k can be defined as:
k=
(1 − ν ) E
(6) Fig. 11. Bending strength of ceramics at room temperature.
Then equation (5) can be simplified as:
R=k
σf α
(7)
The lower the value of R, the better thermal shock resistance of the samples. Referring to the flexural strength at the room temperature in
Fig. 10, it can be seen that the values of strength are very close. It is well known that the lower the coefficient of thermal expansion, the better the thermal shock resistance of the material. Thus, the ceramic system
Fig. 10. Residual strength of ceramics after thermal shock. 20126
Ceramics International 45 (2019) 20121–20127
Z. Tian, et al.
with Sm2O3 has the best thermal shock resistance under this test condition. Furthermore, for the material system studied in this paper, the thermal shock resistance evaluation of the specimens still obeys the relevant rules even though there are micro-crack sources on the surface of the specimens. To summarize, the thermal shock process of ceramic samples in this paper is equivalent to “high temperature heat treatment”. When the ambient temperature is higher than 1000 °C, the fused quartz transformed into cristobalite, accompanying by the growth of nano-cristobalite and the formation of nano-sialon grains. Therefore, the thermal expansion curve of the material increases rapidly under the combined action of the two factors. During the thermal shock process, as the difference of thermal expansion coefficient between matrix material and grown quartz and SiAlON phase, the growing grains will produce compressive stress around them and then cause internal stress. The higher the internal stress of the material with high thermal expansion coefficient, the worse the thermal shock resistance of the material, while the residual strength of the material decreases more obviously after the thermal shock. Therefore, under the combined action of the internal and external factors of the samples, the residual strength of the ceramic decreases sharply with the increase of the thermal shock temperature difference and the cycle index. 4. Conclusions In this study, the effect of rare earth oxides on the thermal shock resistance of in-situ SiAlON phase-enhanced BN-matrix ceramics was investigated. The effect of microstructure and thermal properties of the composites have been studied in detail. Several key conclusions can be drawn as follows. 1) The introduction of rare earth oxides has little effect on the improvement of the flexural strength of the material. When CeO2 was added, the flexural strength of the material was 264.5 MPa, which was only increased by 7.7% compared to the unadded one of 245.6 MPa. 2) The influence of rare earth oxides on thermal expansion coefficient of the material is more obvious. The thermal expansion coefficient of the CeO2-added was increased by 24.6%, which was higher than that with Sm2O3. 3) Different rare earth oxides have significantly different effects on the thermal shock resistance of materials. After 50 cycles at Δt = 1150 °C, the ceramic material added with CeO2 had a residual strength of 157.1 MPa, which was reduced by 40.6% % compared with the room temperature strength. Under the same conditions, the residual strength of the sample added Sm2O3 decreased by 34.7%–167.1 MPa. However, no cracks due to thermal expansion mismatch were found on the surface after thermal shock. 4) The decrease in the residual strength of the ceramic material is mainly caused by the internal stress generated from the mismatch of the thermal expansion coefficient between the new generated and grown phases in the matrix material. Acknowledgements The authors would like to thank Laboratory for Materials in Extreme Environments in UCLA lead by prof. Nasr Ghoniem for testing the sputtering temperature of the sample. And this work was financially supported by Local Innovative and Research Team Project of Guangdong Pearl River Talents Program (Grant No.2017BT01C169), Science and Technology Major Project of Guangdong Province (Grant No.2016B090915002), Science and Technology Major Project of Guangdong Province (Grant No.2017B090911011), the China Postdoctoral Science Foundation (CPSF, Grant No. 2018M633013), Special Fund for the Cultivation of
Science and Technology Innovation of Guangdong University Students in 2018(Grant No. PDJHB0157). Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.ceramint.2019.06.277. References [1] M. Martinez-Sanchez, J.E. Pollard, Spacecraft electric propulsion-an over view, J. Propuls. Power 4 (1998) 688–699. [2] S. Mazouffre, Electric propulsion for satellites and spacecraft: established technologies and novel approaches, Plasma Sources Sci. Technol. 25 (3) (2016) 1–27. [3] J.P. Boeuf, Tutorial: physics and modeling of Hall thrusters, J. Appl. Phys. 121 (1) (2017) 1–23. [4] Y.H. Chiu, L.A. Brad, W. Kip, A.D. Rainer, George F. Karabadzhak, Passive Optical Diagnostic of Xe-Propelled Hall thrusters. I. Emission cross sections, J. Appl. Phys. 99 (2006) 1–10. [5] B.M. Nathan, G. Nicolas, A.C. Mark, Linear geometry Hall thruster with boron nitride and diamond walls, In: 27th Inter. Electr. Propul. Conf., Pasadena, CA, USA, IEPC-01-39. [6] N. Gascon, M. Dudeck, S. Barral, Wall material effects in stationary plasma thrusters. I. Parametric studies of an SPT-100, Phys. Plasmas 10 (2003) 4123–4136. [7] S. Barral, K. Makowski, Z. Peradzynski, Wall material effects in stationary plasma thrusters. II. Near-wall and in-wall conductivity, Phys. Plasmas 10 (2003) 4137–4152. [8] M. Britton, D. Waters, R. Messer, E. Sechkar, Sputtering Erosion Measurement on Boron Nitride as a Hall Thruster Material, Cleveland, NASA/TM-2002-211837. [9] H. Tahara, K. Imanaka, S. Yuge, Effects of channel wall material on thruster performance and plasma characteristics of Hall-effect thrusters, Vacuum 80 (2006) 1216–1222. [10] Y. Raitses, D. Staack, M. Keidar, N.J. Fisch, Electron-wall interaction in Hall thrusters, Phys. Plasmas 12 (5) (2005) 057104. [11] A.M. Rafael, D. Hoang, L.R. Walker Mitchell, Power deposition into the discharge channel of a Hall effect thruster, J. Propuls. Power 30 (1) (2014) 209–220. [12] A.M. Bohnert, Thermal characterization of a hall effect thruster, Thesis for the Degree of Master of Science in Astronautical Engineering (2008) 37–50. [13] T. Kusunose, N. Sakayanagi, T. Sekino, Y. Ando, Y. Ando, Fabrication and characterization of aluminum nitride/boron nitride nanocomposites by carbothermal reduction and nitridation of aluminum borate powders, J. Nanosci. Nanotechnol. 8 (2008) 5846–5853. [14] G.J. Zhang, H. Kita, N. Kondo, T. Ohji, Reactive hot-pressed aluminaboron nitride composites with Y2O3 sintering additive, J. Am. Ceram. Soc. 88 (8) (2005) 2246–2248. [15] H.Z. Zhai, H.N. Cai, X.Z. Yang, J.B. Li, G.F. Guo, C.B. Cao, Preparation and properties of BN-SiO2 composite ceramics, Key Eng. Mater. 336–338 (2007) 1426–1428. [16] Y. Garnier, V. Viel, J.F. Roussel, J. Bernard, Low-energy xenon ion sputtering of ceramics investigated for stationary plasma thrusters, J. Vac. Sci. Technol. A 17 (6) (1999) 3246–3255. [17] X.H. Zhang, R.B. Zhang, G.Q. Chen, W.B. Han, Microstructure, mechanical properties and thermal shock resistance of hot-pressed ZrO2(3Y)- BN composites, Mater. Sci. Eng. A-Struct. 497 (1–2) (2008) 195–199. [18] Y.L. Li, J.X. Zhang, G.J. Qiao, Z.H. Jin, Fabrication and properties of machinable 3Y-ZrO2/BN Nanocomposites, Mater. Sci. Eng. A-Struct. 397 (1–2) (2005) 35–40. [19] X.M. Duan, D.C. Jia, Q.C. Meng, Z.H. Yang, Y. Yu, Y. Zhou, D.R. Yu, Y.J. Ding, Study on the plasma erosion resistance of ZrO2p(3Y)/BN-SiO2 composite ceramics, Composites Part B 46 (2013) 130–134. [20] Y.L. Li, R.X. Li, J.X. Zhang, Enhanced mechanical properties of machinable Si3N4/ BN composites by spark plasma sintering, Mater. Sci. Eng. A-Struct. 483–484 (2008) 207–210. [21] M. Britton, D. Waters, R. Messer, E. Sechkar, Sputtering Erosion Measurement on Boron Nitride as a Hall Thruster Material. Cleveland, NASA, TM-2002-211837. [22] M. Tartz, T. Heyn, C. Bundesmann, H. Neumann, Measuring sputter yields of ceramic materials, 31st Inter. Electr. Propul. Conf. IEPC, Ann Arbor, Michigan, USA, 2009, p. 240. [23] S.A. Khartov, A.B. Nadiradze, I.I. Shkarban, Y.V. Zikeeva, SPT's high lifetime -some problems of solution, 29th Inter. Electr. Propul. Conf. IEPC, Princeton, USA, 2005, p. 62. [24] Z. Tian, X.M. Duan, Z.H. Yang, S.Q. Ye, D.C. Jia, Y. Zhou, Microstructure and erosion resistance of in-situ SiAlON reinforced BN-SiO2 composite ceramics, J. Wuhan Univ. Technol.-Materials Sci. Ed. 31 (2) (2016) 315–320. [25] X.M. Hou, C.S. Yue, A.K. Singh, M. Zhang, Morphological development and oxidation of elongated β-SiAlON material, Corros. Sci. 53 (6) (2011) 2051–2057. [26] Zhuo Tian, Dechang Jia, Duam Xiaoming, Zhihua Yang, Yu Zhou, Effects of AlN content on phase composition, microstructure and mechanical properties of BNbased composite ceramics, J. Chin. Ceram. Soc. 41 (12) (2013) 1603–1608. [27] M. Zeuner, S. Pagano, W. Schnick, Nitridosilicates and oxonitridosilicates: from ceramic materials to structural and functional diversity, Angew. Chem. Int. Ed. 50 (2011) 7754–7775.
20127