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Refractory metal silicides reinforced by in-situ formed Nb2O5 fibers and mullite nanoclusters
Ceramics International 43 (2017) 16362–16370
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Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Refractory metal silicides reinforced by in-situ formed Nb2O5 fibers and mullite nanoclusters Jing-Lian Fana, a b
⁎,1
MARK
, Qiong Lua,1, Pei-Zhong Fengb, Wei Lia
State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China School of Materials Science and Engineering, China University of Mining and Technology, Xuzhou 221116, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Refractory metal silicides In-situ hot pressing Thermal shock resistance Bending strength
There is keen interest in the use of refractory metal silicides as structural materials or thermal barrier coatings for a high temperature environment. However, a long-standing problem for these materials is their poor thermal shock property. To address this challenge, Nb-Al-SiC elements were introduced into the MoSi2 matrix and consolidated by in-situ hot pressing. We find that this treatment leads to improved performance of MoSi2 composites in high temperature thermal shock resistance and bending strength. After in-situ HPing, the Nb, Al2O3 particles, and SiC nanoclusters were uniformly dispersed in the MoSi2 matrix and inhibited the movement of dislocation, resulting in a strengthening effect. During the thermal shock process, the fragmentized oxide layer present in the surface of the pure MoSi2 alloy disappeared completely, and a dense multi-component oxide layer was formed in-situ on the surface of the MoSi2 composites. The dense multi-component oxide layer was composed of SiO2 glass, fiber-structured Nb2O5, and nano-sized mullite phases. The fiber structured and nano-sized oxide phases play an important role in strengthening the oxide layer.
1. Introduction Refractory metal silicides (for example, MoSi2, WSi2, NbSi2, TaSi2, and TiSi2) are potential candidate materials for applications in aerospace [1], turbine engines [2], and very large-scale integration [3], as they offer high melting point, high thermal and electrical conductivity, and high strength and high creep properties at high temperatures. Among these silicides, MoSi2 [4,5] has received especially strong attention due to its relatively low density (6.24 g/cm3) and its thermodynamic compatibility with many other ceramic reinforcements. MoSi2 also has excellent isothermal oxidation resistance at high temperature (1000–1600 °C) [6] and has been widely used as an oxidation protection coating [7] for high-temperature metal [8] and composites materials [9] due to the formation of a continuous SiO2 oxide layer with low oxygen permeability at elevated temperature (10–13 g/cm s at 1200 °C, 10–11 g/cm s at 2200 °C) [10]. However, the use of MoSi2 is limited by its poor high temperature thermal shock property. As most high-temperature materials experience thermal shock, especially between room temperature and high temperature, the cycle thermal shock properties of MoSi2 are recognized as an important baseline of thermal stability at the operating temperature [11]. In the aerospace field, for example, turbine blades can reach 1200–1300 °C during take off [2] and are
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cooled to room temperature, requiring both excellent isothermal oxidation and thermal shock performance of the materials. So, it is urgent to find an effective method to improve the thermal shock property of MoSi2 alloy. The poor thermal shock resistance of MoSi2 is due to the initiation of micro-cracks in the SiO2 glassy film and their extensive propagation [12], which will gradually decrease the bending strength of MoSi2 materials with increased thermal shock cycle. This extensive creation of micro-cracks is due to the mismatch of the thermal conductivity and the coefficient of thermal expansion (CTE) between the oxide layer and the MoSi2 matrix [13]. Improving the high temperature thermal shock properties in air or oxygen-rich environments has been a major goal for MoSi2 alloys over recent decades [13,14]. One approach is to alleviate the CTE mismatch between SiO2 oxide layer and MoSi2 matrix by the addition of SiC reinforcement, which can effectively decrease the CTE value of the MoSi2 matrix [15–18]. However, with the addition of SiC, the component of oxide layer is still SiO2 phase after high temperature oxidation, which indicates that the thermal stability of the oxide layer was not enhanced [18]. Therefore, a third element, Nb or Al, is typically added to MoSi2SiC composites to enhance the thermal stability of the oxide layer and the adherence of the oxide layer to the matrix. These elements are chosen because the oxides of the Nb and Al elements have excellent
Corresponding author. E-mail address: [email protected] (J.-L. Fan). These authors contributed equally to this work.
http://dx.doi.org/10.1016/j.ceramint.2017.09.010 Received 23 June 2017; Received in revised form 25 August 2017; Accepted 3 September 2017 Available online 04 September 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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Fig. 1. Illustration of the experimental process.
Fig. 2. (a) Backscatter electronic (BSE) micrographs of the in-situ HPing MoSi2 alloy; Inset is the enlarged view of dotted box. (b) Oxidation kinetics of MoSi2 alloy at 1200 °C in air environment. (c) SEM micrograph of oxide layer which in-situ formed in the surface of the MoSi2 alloy after 120 h oxidation. (d) SEM micrographs of the cross section of MoSi2 alloy after 120 h and five thermal shock cycle tests.
oxidation resistance, and are more stable than SiO2 at high temperature [19–22]. Based on this, the MoSi2 composites prepared with additional Nb, Al alloying and SiC reinforcement elements are expected to exhibit excellent thermal shock resistance. However, the thermal shock property of such composites has not yet been reported. In this work, we synthesized different compositions of (Mo0.94Nb0.06)(Si0.97Al0.03)2-x vol% SiC, and investigated the effects of Nb-Al-SiC content on the thermal shock properties of the materials in air atmosphere at 1200 °C. The effects of oxidation time and thermal shock cycle time on the thermal shock mechanism were studied by
performing oxidation kinetics followed by characterization of the microstructure of the composites and oxidized products. 2. Experimental procedures Molybdenum (2.0–2.5 µm, 99.9% purity), silicon (− 300 mesh, 99.9% purity), niobium (− 300 mesh, 99.9% purity), aluminum (100–200 mesh, 99.0% purity), and nano-sized SiC (40 nm, 99.9% purity) powders were used to form different MoSi2 composites: MoSi2, (Mo0.94Nb0.06)(Si0.97Al0.03)2-5 vol% SiC (NAS1), (Mo0.94Nb0.06)
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Fig. 3. (a) XRD patterns of the in-situ HPing MoSi2 alloy; (b) SEM micrograph of in-situ HPing MoSi2 alloy; (c–e) Energy-dispersive spectroscopy analysis of the different phase: (c) gray area 1, (d) white area 2, (e) black area 3 in (b).
(Si0.97Al0.03)2-10 vol% SiC (NAS2), and (Mo0.94Nb0.06)(Si0.97Al0.03)215 vol% SiC (NAS3). The starting mixtures were uniformly mixed by planetary ball milling for 240 min at 450 rounds per minute using an agate jar and agate balls with alcohol as the milling medium. Next, the ball-milled powders were dried and consolidated by in situ hot pressing (HPing) at 1400 °C and 27.5 MPa for 90 min under vacuum (< 8.0 × 10−1 Pa). The sintered plates (φ 50 mm × 5 mm) were cut into test pieces by wire-electrode cutting that measured 22 mm × 4 mm × 3 mm. The specimens surface was grounded and polished to a mirror finish and cleaned in acetone before the high temperature oxidation and thermal shock tests. The overall experimental process is illustrated in Fig. 1. The specimens were placed upon a corundum support and then heated at 1200 °C in air environment in an electric furnace to investigate the isothermal oxidation and thermal shock behavior. After a specified time of oxidation at 1200 °C, the specimens were removed from the furnace and quickly cooled to room temperature. The cumulative weight change of the specimens was reported as a function of exposure time by an electronic balance with a sensitivity of ± 0.1 mg. The specimens were then returned to the furnace for the next thermal cycle. The bending strength of the in-situ HPing and the heat-treated composites was tested by a three-point loading system, using a 18 mm span and a crosshead speed of 0.5 mm min−1 (according to test methods, ASTM C1161-2013 standard method for bending strength), and at least five specimens were tested for each condition. A scanning electron microscope (SEM; JEOLJSM-6360 and JSM-890
operated at 20 kV) with energy-dispersive X-ray analysis (EDX; Phoenix EDAX 2000) was employed to characterize the surface and cross-section microstructure of the composites. The phase compositions were identified using an X-ray diffractometer (XRD, Bruker D8 Advance, Karlsruhe, Germany). For the preparation of the transmission electron microscope (TEM) sample, a FEI Helios NanoLab (FIB/SEM) was employed. The morphology, structure, and elemental composition of the oxide layer were characterized using TEM analysis (FEI Tecnai F30 operated at 300 kV). 3. Results 3.1. Thermal shock failure of pure MoSi2 alloy The backscatter electron micrograph (BSE) of the HPing alloy is shown in Fig. 2a and reveals that the microstructure of pure MoSi2 consists of three kinds of phases (gray phase, white phase, and black phase). Fig. 3a shows the XRD patterns of the HPing MoSi2 alloy. It can be seen that MoSi2 is the main phase in the alloy, and miner amount of Mo5Si3 is also found. The EDX results (Fig. 3c–e) show that the gray phase is MoSi2, the white phase is Mo5Si3, and the black phase is SiO2. The EDX results are consistent with XRD analysis. Only SiO2 phase cannot be identified by XRD analysis, which is due to that the small amount of SiO2 has lower mass absorption coefficient (36.41 cm2 g−1) than MoSi2 dose (124 cm2 g–1) [23,24]. The formation of SiO2 and Mo5Si3 is due to that the raw powders containing a small volume of
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Fig. 4. (a) XRD pattern on the surface oxide layer of MoSi2 alloy after 24 h oxidation at 1200 °C and one thermal shock cycle. (b) SEM micrograph of MoSi2 after oxidization at 1200 °C for 24 h and one thermal shock cycle, and (c) EDX analysis of (b).
oxygen, and sintering such powders results in a small fraction of amorphous silica in the final product. Consumption of Si to form SiO2 renders the materials slightly Mo rich leading to the formation of Mo5Si3. This observation has been reported previously [25]. Fig. 2b shows the oxidation kinetics of MoSi2 alloy in air at 1200 °C. Obviously, the thermal shock behavior of the MoSi2 alloy can be divided into three sections marked as I, II, and III by analysis of the oxidation curve. In section I (during the first thermal shock), the SiO2 layer determined by EDX and XRD (as shown in Fig. 4) was formed, which provides shortterm protection of the matrix (< 72 h, in section II) (Fig. 5a–c). As the thermal shock cycle was increased (in section III), the mass of the MoSi2 alloy was decreased sharply. At the same time, the surface microstructure of the oxide layer was changed greatly, a process dominated by the failure of the SiO2 oxide layer (Fig. 5d, e). After 120 h oxidation and five thermal shock cycles, the surface oxide layer was disintegrated into fragments (as show in Fig. 2c–d), due to the large mismatch of the coefficient of thermal expansion (CTE) and the thermal conductivity (TC) between the SiO2 layer and the MoSi2 matrix [14,26–28]. These cracks can provide channels that allow oxygen to permeate into the inner body and lead to the continuous volatilization of MoO3 [29], consistent with the greater mass loss during oxidation (as shown in Fig. 2b). 3.2. Microstructure and thermal shock properties of NAS1, NAS2, and NAS3 composites As shown in Fig. 6a, with the addition of Nb-Al-SiC elements, there were obvious homogeneous large black phases. Compared with the
insert picture in Figs. 2a and 6a, it can be seen that the content of small black phase (SiO2) and white phase (Mo5Si3) decreased in the MoSi2 composites. The scanning transmission electron microscopy (STEM) image of the NAS1 composite (Fig. 6b) shows that Al elements were consumed into Al2O3 [30], and Nb elements were distributed in the interior of the MoSi2 grains, as determined by EDX and shown in Fig. 7, which can increase the adherence of the oxide layer to the MoSi2 matrix [31]. The large black phases (Fig. 6a) are SiC nanocluster particles (Fig. 6c) that will decrease the thermal stress between the oxide layer and the matrix in a high temperature oxidation environment by increasing the thermal conductivity [32] and decreasing the CTE value of the matrix [18]. Compared with pure MoSi2 alloy, the NAS1, NAS2, and NAS3 composites exhibited excellent thermal shock performance. Fig. 6e shows the mass gain per unit area as a function of oxidation time for the composites oxidized at 1200 °C in air for 120 h. During the oxidation test, the weight of the composites increased to follow a near-parabolic law. The composites exhibited a much smaller weight gain than the pure MoSi2 alloy. In order to approximate the oxidation rate constants for a quantitative comparison, the relatively parabolic section of the curve was fitted to the parabolic rate equation given by [33]:
(ΔW / A)2 = kp t + C Where △W is the mass gain, A is the specimen's surface area, kp is the oxidation rate constant, and t is the oxidation time. The results are shown in Fig. 8. With the increase of Nb-Al-SiC concentration, the composites exhibited a strong trend for reduced weight gain and decreased oxidation
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Fig. 5. SEM images of surface of oxide layer of the MoSi2 alloy at different times. (a) 24 h; (b) 48 h; (c) 72 h; (d) 96 h; (e) 120 h.
Fig. 6. (a) BSE micrograph of the surface of NAS1composite; Inset is the enlarged view of dotted box; (b), (c) STEM image of NAS1composite, a selected area electron diffraction pattern recorded on dash line circle area is shown in the inset of (c); (d) STEM image of in-situ HPing NAS1 composite after bending test; (e) Oxidation kinetics of NAS1-3 composites; (f)–(h) SEM micrographs of oxide layer of NAS1(f), NAS2(g) and NAS3(h) composites.
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Fig. 7. (a)TEM analysis of in-situ HPing NAS1 composite; (b) EDX analysis of particle A in (a); (c) EDX analysis of particle B in (a).
analyses of the surface oxide layer of NAS1, NAS2, and NAS3 composites after 120 h oxidation at 1200 °C and five thermal shock cycles are shown in Fig. 9a–c. The results show that the dense glassy layers are composed of SiO2, fiber structure Nb2O5, and mullite (3Al2O3·2SiO2) phase. The bright-field TEM image of the oxide layer cross-section revealed that the fibrous Nb2O5 appeared as a single crystal with a polygon structure embedded in the surface of the oxide layer and the mullite phase was nano-sized and located in the interface of the oxide layer and matrix (Fig. 9d–f). Additionally, the three phases were immiscible, which indicated that a multi-component oxide layer was synthesized. 4. Discussion
Fig. 8. Wight gain2 plotted as a function of exposure time.
Table 1 Calculation results of volume ratio of SiO2 matrix and fiber structure Nb2O5 in surface oxide layer of NAS1-3 composites after five thermal shock cycle through Image J2x software. Composites
NAS1 NAS2 NAS3
Volume ration SiO2 matrix
Fiber structure Nb2O5
88.8 86.3 83.8
11.2 13.7 16.2
rate, not only in the early stage but also in the stable stage of oxidation. Furthermore, the fragmentized oxide layer present in the surface of the pure MoSi2 alloy (Fig. 2c, d) disappeared completely. A fiber-structured dense oxide layer was formed in-situ on the surface of the NAS1, NAS2, and NAS3 composites, and this layer may effectively prevent the matrix from further oxidation during the 120 h treatment at 1200 °C and five thermal shock cycles (Fig. 6f–h). At the same time, the grain sizes of the fiber structure phase became finer and the number of the phase increased gradually. The volume ratios of the fiber structure phase in the oxide layer of NAS1-NAS3 composites were 11.2%, 13.7%, and 16.2%, respectively (Table 1). We now present details of the phase composition and microstructure of the surface and cross section fibrous structures oxide layer formed insitu on the surface of the NAS1-3 composites. This will allow us to explain the excellent thermal shock resistance of MoSi2 composites resulting from the addition of Nb-Al-SiC elements. The XRD and EDX
In order to fully understand the anti-oxidation and anti-thermal shock mechanisms, the thermodynamic and dynamic behaviors of the oxidation of the composites were next studied. Thermodynamic calculations (delta-Gibbs free energy of reaction) show that once the composites were put into the furnace, oxidation (reaction (1)–(7)) can occur and some gases by-products are generated, such as CO2, CO, MoO3, whose emission does not increase the protection of the composites [34]. Depending on the thermodynamic calculations, MoSi2 may be oxidized into Mo5Si3 and SiO2 (reaction (1)), and then Mo5Si3 was oxidized into MoO3 and SiO2 (reaction (2)) [35]. However, Ibano [36] reported that simultaneous oxidation of molybdenum and silicon occurs at the initial stage of oxidation at ambient pressure. Zhu [37] also reported that simultaneous oxidation occurs at the initial stage of oxidation (< 1 h) in all temperature ranges from 900 °C to 1500 °C in an air environment. Thus, for the HPing composites, the simultaneous oxidation of molybdenum and silicon occurred at the initial stage of the oxidation process in the air environment, as shown in reaction (3). Under constant temperature and pressure, the enthalpy change of reaction is equal to the heat evolved by the reaction [38]. At the same time, the more heat evolved by the oxidation reaction, the higher the affinity of the component for oxygen. According to the calculated enthalpy change of the reaction (3,4), the heat values of exothermic reactions during oxidation are in the following sequence: Qreaction (4) > Qreaction (3) , indicating that the affinity of the Nb for oxygen is higher than that of Si. Prigogine et al. [39] reported that the thermodynamics of irreversible phenomena, based mainly on the concept of entropy production, leads to linear and homogeneous relationships between affinities and reaction rates. In such composites, fiber structured Nb2O5 can be preferentially formed and then could be enwrapped by SiO2 during oxidation, indicating that fiber-structured Nb2O5 was present in the top of the oxide layer (as shown in Fig. 9d). Based on these results, a schematic illustration of the phase structure evolution during the oxidation
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Fig. 9. (a) XRD analysis on the surface of oxide layer of NAS1-3 composites at 1200 °C after heat treat; (b) EDX analysis of A region in Fig. 6f; (c) EDX analysis of B region in Fig. 6f; (d) Cross-sectional FIB microstructure of the oxide layer formed in the surface of the matrix with NAS3 which were extracted from the regions of the sample indicated in Fig. 6h using focused ion beam techniques; (e) Selected area electron diffraction (SAED) pattern of the polygon in the Fig. 9d; (f) Typical diffraction pattern of the nano crystalline region in the Fig. 9d.
Fig. 10. Schematic illustration of the formation process of the multi-component oxide layer upon oxidation and thermal shock test for different stages.
of NAS1, NAS2, and NAS3 composites is shown in Fig. 10. It has been reported that the Nb2O5 and mullite are stable oxides, with high coefficient of thermal expansion (CTE) values (5.3 × 10−6/K [40], 5.3 × 10−6/K [41]) and a high degree of oxidation resistance [22]. The uniformly distributed fiber structured Nb2O5 and nano-sized mullite should increase the viscosity of the multi-component SiO2 glass layer and reduce the oxygen diffusion rate [20], thus improving the oxidation resistance of the oxide layer. Therefore, with the increased fiberstructure Nb2O5 content, the oxygen diffusion rate will decrease gradually. This is consistent with the observed thinner oxide layer (the oxide layer thickness of the heat-treated NAS1-3 composites was 21.60 µm, 7.30 µm, and 6.19 µm, respectively, as shown in Fig. 11), and consistent with the oxidation kinetics shown in Fig. 6e. Additionally, nano-sized mullite phase was present in the interface of the oxide layer and matrix and can reduce the stress concentration [42,43] and retard the propagation of micro cracks [44,45]. Consequently, the Nb-Al-SiC elements are expected to significantly improve the thermal shock resistance of the MoSi2 composites.
The possible reactions during oxidation at 1200 °C in air environment are：
5/7MoSi2 (s) + O2 (g) → 1/7Mo5 Si3 (s) + SiO2 (s)
ΔG(1400K) = −615 kJ mol−1 ΔH(1400K) = −849 kJ mol−1
(1)
2/21Mo5 Si3 +O2 (g) → 10/21MoO3 (g) + 6/21SiO2 (s)
ΔG(1400K) = −291.87 kJ mol−1 ΔH(1400K) = −474 kJ mol−1 (2)
2/7MoSi2 (s) + O2 (g) → 2/7MoO3 (g)
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+ 4/7SiO2 (s)
ΔG(1400K) = −460 kJ mol−1 ΔH(1400K) = −666 kJ mol−1 (3)
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Fig. 11. The effect of Nb-Al-SiC elements on the thickness of oxide layer. (a)–(c) Scanning electron micrographs of the fracture plane of NAS1(a), NAS2(b) and NAS3(c) composites after oxidized at 1200 °C for 120 h.
This decrease in the strength is attributed to the formation of cracks in the thick oxide layer [46]. After introducing the Nb-Al-SiC elements into the MoSi2 matrix, the bending strength of the composites increased gradually and even exceeded that of the initial HPing composites and reached a maximum of 496 MPa (NAS3 composite). This increase was due to the formation of a multi-component dense oxide layer that could heal the surface flaws resulting from sample processing and machining [47]. At the same time, the fiber structured Nb2O5 and nano-sized mullite present in the oxide layer can act as reinforcements (like whiskers [48], nanoparticles [49], and nanoclusters [50]), and retard the propagation of cracks. And heat treatment at high temperature could relieve residual stress of in-situ MoSi2 composites and result in improving mechanical properties [51]. This also needs to be investigated in future studies. Sciti et al. [52] reported that the bending strength of ZrB2-MoSi2 increased by up to 3% after oxidation at 1200 °C for 100 h, which was quite surprising since microstructural analyses revealed remarkable microstructural degradation. Krnel et al. [53] reported that after oxidized at 1200 °C for 100 h, the AlN–SiC–MoSi2 composites showed little strength improvement.
Fig. 12. Bending strength of in-situ HPing composites and heat treated composites.
4/5Nb(s) + O2 (g) → 2/5Nb2 O5 (s)
ΔG(1400K) = −520 kJ mol−1 ΔH(1400K) = −700 kJ mol−1
(4) 5. Conclusions
1/2SiC(s) + O2 (g) → 1/2SiO2 (s) + 1/2CO2 (g) ΔG(1400K) = −498 kJ mol−1
(5)
The effects of addition of Nb-Al-SiC elements on the microstructure and high temperature thermal shock resistance of MoSi2 composites were investigated. The following conclusions were derived:
2/3SiC(s) + O2 (g) → 2/3SiO2 (s) + 2/3CO (g) ΔG(1400 K) = −556 kJ mol−1
(6)
3Al2O3 (s) + 2SiO2 (s) → 3Al2O3·2SiO2 (s) ΔG(1400 K) = −20 kJ mol−1 (7) To investigate the mechanical properties of MoSi2 composites before and after high temperature oxidation and five thermal shock cycles, the three-point bending test was performed and the results are shown in Fig. 12. With increased amount of Nb-Al-SiC elements, the bending strength of the in-situ HPing composites increased gradually. This is caused by inhibition of dislocation by the Al2O3, Nb particles, and SiC nanocluster (as shown in Fig. 6c–d) and results in a strengthening effect [25]. After 120 h oxidation and five thermal shock cycles (△T = 1200 °C), the retained bending strength of MoSi2 and the NAS1, NAS2, NAS3 composites exhibited different trends. The bending strength of the pure MoSi2 after heat treatment decreased from 330 MPa (initial HPing alloy) to 258 MPa with ~ 78% retention in bending strength.
(a) The elements of Nb-Al2O3-SiC were dispersed in the MoSi2 matrix homogeneously after in-situ hot pressing. (b) After the thermal shock test, the fragmentized oxide layer present in the surface of pure MoSi2 alloy disappeared completely and a dense multi-component oxide layer was formed in-situ on the surface of the NAS1, NAS2, and NAS3 composites. This dense multi-component oxide layer was composed of SiO2, fiber-structured Nb2O5, and nano-sized mullite. (c) With increased Nb-Al-SiC content, the bending strength of the composites after hot treat increased gradually due to the formation of multi-component of oxide layer.
Acknowledgements This research was supported by the National Natural Science Fund Project of China (51534009).
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References [1] M. Yamaguchi, H. Inui, K. Ito, High-temperature structural intermetallics, Acta Mater. 48 (2000) 307–322. [2] J.H. Perepezko, The hotter the engine, the better, Science 326 (2009) 1068–1069. [3] S.P. Murarka, D.B. Fraser, A.K. Sinha, H.J. Levinstein, Refractory silicides of titanium and tantalum for low-resistivity gates and interconnects, IEEE J. Solid-State Circuit 15 (1980) 474–482. [4] A.K. Vasudevan, J.J. Petrovic, A comparative overview of molybdenum disilicide composites, Mater. Sci. Eng. A 155 (1992) 1–17. [5] P.J. Meschter, D.S. Schwartz, Silicide-matrix materials for high-temperature applications, JOM 41 (1989) 52–55. [6] R. Mitra, Mechanical behavior and oxidation resistance of structural silicides, Int. Mater. Rev. 51 (2006) 13–64. [7] T. Feng, H.J. Li, Q.G. Fu, X. Yang, H. Wu, High-temperature erosion resistance and aerodynamic oxidation mechanism of multi-layer MoSi2–CrSi2–Si/SiC coated carbon/carbon composites in a wind tunnel at 1873 K, Carbon 50 (2012) 2171–2178. [8] Y. Wang, D.Z. Wang, J.H. Yan, A.K. Sun, Preparation and characterization of molybdenum disilicide coating on molybdenum substrate by air plasma spraying, Appl. Surf. Sci. 284 (2013) 881–888. [9] L. Su, O. Lu-Steffes, H. Zhang, J.H. Perepezko, An ultra-high temperature Mo-Si-B based coating for oxidation protection of NbSS/Nb5Si3 composites, Appl. Surf. Sci. 337 (2015) 38–44. [10] J.F. Huang, H.J. Li, X.R. Zeng, X.B. Xiong, Y.W. Fu, Progress on the oxidation protective coating of carbon-carbon composites, New Carbon Mater. 20 (2005) 373–379. [11] S.P. Lee, J.K. Lee, C.H. Leem, Cyclic thermal shock properties of MoSi2 intermetallic materials, Int. J. Mod. Phys. B 24 (2010) 2922–2927. [12] S.L. Lee, Y.H. Kim, Y.S. Shin, J.L. Lee, J.Y. Park, Thermal shock damage and mechanical property of MoSi2 materials by NDE technique, Key Eng. Mater. 321–323 (2006) 930–933. [13] Y.L. Zhang, H.J. Li, X.F. Qiang, K.Z. Li, S.Y. Zhang, C/SiC/MoSi2–simultilayer coatings for carbon/carbon composites for protection against oxidation, Corros. Sci. 53 (2011) 3840–3844. [14] Q.G. Fu, H.J. Li, Y.J. Wang, K.Z. Li, X.H. Shi, B2O3 modified SiC–MoSi2 oxidation resistant coating for carbon/carbon composites by a two-step pack cementation, Corros. Sci. 51 (2009) 2450–2454. [15] M.G. Hebsur, Development and characterization of SiC(f)/MoSi2–Si3N4(p) hybrid composites, Mater. Sci. Eng. A 261 (1999) 24–37. [16] H. Chen, M. Suzuki, S. Sodeoka, T. Inoue, K. Ueno, New approach to MoSi2/SiC intermetallic-ceramic composite with B4C, J. Mater. Sci. 36 (2001) 5773–5777. [17] J.F. Huang, B. Wang, H.J. Li, M. Liu, L.Y. Cao, C.Y. Yao, A MoSi2/SiC oxidation protective coating for carbon/carbon composites, Corros. Sci. 53 (2011) 834–839. [18] J. Liu, Q.M. Gong, Y. Shao, D.M. Zhuang, J. Liang, In-situ fabrication of MoSi2/ SiC–Mo2C gradient anti-oxidation coating on Mo substrate and the crucial effect of Mo2C barrier layer at high temperature, Appl. Surf. Sci. 308 (2014) 261–268. [19] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for 2000 C+ hypersonic aerosurfaces: theoretical considerations and historical experience, J. Mater. Sci. 39 (2004) 5887–5904. [20] Z.H. Hao, W. Sun, X. Xiong, Z.K. Chen, Y.L. Wang, Effects of Ti/Al addition on the microstructures and ablation properties of Cf/C–MoSi2–SiC composites, J. Eur. Ceram. Soc. 36 (2016) 457–464. [21] J. Cheng, S. Yi, J.S. Park, Oxidation behavior of Nb-Si-B alloys with the NbSi2 coating layer formed by a pack cementation technique, Int. J. Refract. Met. Hard Mater. 41 (2013) 103–109. [22] L. Ingemarsson, M. Halvarsson, J. Engkvist, T. Jonsson, K. Hellström, L.G. Johansson, J.E. Svensson, Oxidation behavior of a Mo(Si,Al)2-based composite at 300–1000 °C, Intermetallics 18 (2010) 633–640. [23] Q. Lu, G.M. Zhu, X.H. Wang, Z. Wang, P.Z. Feng, Effects of raw materials on synthesis, microstructure and properties of MoSi2-10 vol% SiC composites, Trans. Indian Ceram. Soc. 75 (2016) 33–39. [24] B.D. Cullity, Elements of X-ray Diffraction, 1st edition, Addison-Wesley Publishing, USA, 1959, pp. 466–467. [25] Q. Lu, G.M. Zhu, X.H. Wang, J.L. Fan, Synthesis and mechanical properties of (Mo0.94Nb0.06)(Si0.97Al0.03)2-x vol% SiC composites fabricated via SHS-HP, Ceram. Int. 43 (2017) 3424–3430. [26] M.C. Chu, S.J. Cho, Y.C. Lee, H.M. Park, Crack healing in silicon carbide, J. Am.
Ceram. Soc. 87 (2004) 490–492. [27] C. Hu, M. Kiene, P.S. Ho, Thermal conductivity and interfacial thermal resistance of polymeric low k films, Appl. Phys. Lett. 79 (2001) 4121–4123. [28] W.W. Beaver, A.J. Stonehouse, R.M. Paine, Development of intermetallic compounds for aerospace applications, in: Plansee Proceedings 5th, 1965, pp. 682–700. [29] Y.L. Zhang, H.J. Li, X.Y. Yao, K.L. Li, X.F. Qiang, Oxidation protection of C/SiC coated carbon/carbon composites with Si-Mo coating at high temperature, Corros. Sci. 53 (2011) 2075–2079. [30] I.M. Allam, H.C. Akuezue, D.P. Whittle, Influence of small Pt additions on Al2O3 scale adherence, Oxid. Met. 14 (1980) 517–530. [31] R. Mitra, K. Sadananda, C.R. Feng, Effect of microstructural parameters and Al alloying on creep behavior, threshold stress and activation volumes of molybdenum disilicides, Intermetallics 12 (2004) 827–836. [32] J.Y. Qiu, Y. Hotta, K. Watari, K. Mitsuishi, Enhancement of densification and thermal conductivity in AlN ceramics by addition of nano-sized particles, J. Am. Ceram. Soc. 89 (2006) 377–380. [33] D. Monceau, B. Pieraggi, Determination of parabolic rate constants from a local analysis of mass-gain curves, Oxid. Met. 50 (1998) 477–493. [34] A.A. Sharif, High-temperature oxidation of MoSi2, J. Mater. Sci. 45 (2010) 865–870. [35] Q. Zhu, K. Shobu, Y. Zeng, T. Watanabe, Oxidation behavior of hot-pressed MoSi2TiC composite, J. Mater. Sci. 36 (2001) 313–319. [36] A. Ibano, K. Yoshimi, A. Yamauchi, R. Tu, K. Maruyama, K. Kurokawa, T. Goto, High temperature oxidation behavior of MoSi2 under low pressure atmosphere, Mater. Sci. Forum 561 (2007) 427–430. [37] G.M. Zhu, X.H. Wang, Q. Lu, G.Z. Wu, P.Z. Feng, High-temperature crack-healing behavior and strength recovery of (MoNb)Si2, Appl. Surf. Sci. 343 (2015) 41–48. [38] R.A. Alberty, Standard Gibbs free energy, enthalpy, and entropy changes as a function of pH and pMg for several reactions involving adenosine phosphates, J. Biol. Chem. 244 (1969) 3290–3302. [39] I. Prigogine, P. Outer, C.L. Herbo, Affinity and reaction rate close to equilibrium, J. Phys. Chem. 52 (1948) 321–331. [40] W.R. Manning, O. Hunter, F.W. Calderwood, D.W. Stacy, Thermal expansion of Nb2O5, J. Am. Ceram. Soc. 55 (1972) 342–347. [41] Y.S. Toulouhain, R.K. Kirby, R.E. Talor, T.Y. Lee, Thermophysical Properties of Matter 13 IFI-Plenum, New York, 1977, p. 25. [42] Q.G. Fu, Y.C. Shan, S.Q. Wen, H.J. Li, Y.H. Chu, X.S. Liu, Oxidation pre-treatment to improve the mechanical property and oxidation resistance of Si–Mo–Cr coated C/C composites, Corros. Sci. 83 (2014) 103–110. [43] S.M. Choi, H. Awaji, Nanocomposites – a new material design concept, Sci. Technol. Adv. Mater. 6 (2005) 2–10. [44] T. Aoki, H. Hatta, T. Hitomi, H. Fukuda, I. Shiota, SiC/C multi-layered coating contributing to the antioxidation of C/C composites and the suppression of throughthickness cracks in the layer, Carbon 39 (2001) 1477–1483. [45] Q.G. Fu, J.Y. Jing, H.J. Li, X.X. Jin, L. Zhuang, L. Li, Design of an inlaid interface structure to improve the oxidation protective ability of SiC-MoSi2-ZrB2 coatings for C/C composites, Ceram. Int. 42 (2016) 4212–4220. [46] G.M. Zhu, X.H. Wang, Z. Sun, P.Z. Feng, F. Akhtar, Effect of annealing environment on the crack healing and mechanical properties of (Mo0.97Nb0.03)(Si0.97Al0.03)2, J. Alloy. Compd. 634 (2015) 109–114. [47] W.M. Guo, G.J. Zhang, Oxidation resistance and strength retention of ZrB2–SiC ceramics, J. Eur. Ceram. Soc. 30 (2010) 2387–2395. [48] Y. Mao, L. Peng, S. Wang, L.X. Xi, Microstructural characterization of graphite/ CuCrZr joints brazed with CuTiH2 Ni-based fillers, J. Alloy. Compd. 716 (2017) 81–87. [49] Z. Zhan, R. Xu, X. Zheng, Q. Fu, M. Wu, Y. Lei, Effective approach to strengthen plasmon resonance localized on top surfaces of Ag nanoparticles and application in surface-enhanced Raman spectroscopy, Nanotechnology 27 (2016) 445301. [50] S.S. Xu, et al., Independence of work hardening and precipitation strengthening in a nanocluster strengthened steel, J. Alloy. Compd. 712 (2017) 573–578. [51] P. Peralta, R. Dickerson, J.R. Michael, Residual thermal stresses in MoSi2–Mo5Si3 insitu composites, Mater. Sci. Eng. A 261 (1999) 261–269. [52] D. Sciti, M. Brach, A. Bellosi, Long-term oxidation behavior and mechanical strength degradation of a pressurelessly sintered ZrB2–MoSi2 ceramic, Scr. Mater. 53 (2005) 1297–1302. [53] K. Krnel, D. Sciti, A. Bellosi, Influence of long term oxidation on the microstructure, mechanical and electrical properties of pressureless sintered AlN–SiC–MoSi2 ceramic composites, J. Eur. Ceram. Soc. 23 (2003) 3135–3146.
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Refractory metal silicides reinforced by in-situ formed Nb2O5 fibers and mullite nanoclusters Jing-Lian Fana, a b
⁎,1
MARK
, Qiong Lua,1, Pei-Zhong Fengb, Wei Lia
State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China School of Materials Science and Engineering, China University of Mining and Technology, Xuzhou 221116, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Refractory metal silicides In-situ hot pressing Thermal shock resistance Bending strength
There is keen interest in the use of refractory metal silicides as structural materials or thermal barrier coatings for a high temperature environment. However, a long-standing problem for these materials is their poor thermal shock property. To address this challenge, Nb-Al-SiC elements were introduced into the MoSi2 matrix and consolidated by in-situ hot pressing. We find that this treatment leads to improved performance of MoSi2 composites in high temperature thermal shock resistance and bending strength. After in-situ HPing, the Nb, Al2O3 particles, and SiC nanoclusters were uniformly dispersed in the MoSi2 matrix and inhibited the movement of dislocation, resulting in a strengthening effect. During the thermal shock process, the fragmentized oxide layer present in the surface of the pure MoSi2 alloy disappeared completely, and a dense multi-component oxide layer was formed in-situ on the surface of the MoSi2 composites. The dense multi-component oxide layer was composed of SiO2 glass, fiber-structured Nb2O5, and nano-sized mullite phases. The fiber structured and nano-sized oxide phases play an important role in strengthening the oxide layer.
1. Introduction Refractory metal silicides (for example, MoSi2, WSi2, NbSi2, TaSi2, and TiSi2) are potential candidate materials for applications in aerospace [1], turbine engines [2], and very large-scale integration [3], as they offer high melting point, high thermal and electrical conductivity, and high strength and high creep properties at high temperatures. Among these silicides, MoSi2 [4,5] has received especially strong attention due to its relatively low density (6.24 g/cm3) and its thermodynamic compatibility with many other ceramic reinforcements. MoSi2 also has excellent isothermal oxidation resistance at high temperature (1000–1600 °C) [6] and has been widely used as an oxidation protection coating [7] for high-temperature metal [8] and composites materials [9] due to the formation of a continuous SiO2 oxide layer with low oxygen permeability at elevated temperature (10–13 g/cm s at 1200 °C, 10–11 g/cm s at 2200 °C) [10]. However, the use of MoSi2 is limited by its poor high temperature thermal shock property. As most high-temperature materials experience thermal shock, especially between room temperature and high temperature, the cycle thermal shock properties of MoSi2 are recognized as an important baseline of thermal stability at the operating temperature [11]. In the aerospace field, for example, turbine blades can reach 1200–1300 °C during take off [2] and are
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cooled to room temperature, requiring both excellent isothermal oxidation and thermal shock performance of the materials. So, it is urgent to find an effective method to improve the thermal shock property of MoSi2 alloy. The poor thermal shock resistance of MoSi2 is due to the initiation of micro-cracks in the SiO2 glassy film and their extensive propagation [12], which will gradually decrease the bending strength of MoSi2 materials with increased thermal shock cycle. This extensive creation of micro-cracks is due to the mismatch of the thermal conductivity and the coefficient of thermal expansion (CTE) between the oxide layer and the MoSi2 matrix [13]. Improving the high temperature thermal shock properties in air or oxygen-rich environments has been a major goal for MoSi2 alloys over recent decades [13,14]. One approach is to alleviate the CTE mismatch between SiO2 oxide layer and MoSi2 matrix by the addition of SiC reinforcement, which can effectively decrease the CTE value of the MoSi2 matrix [15–18]. However, with the addition of SiC, the component of oxide layer is still SiO2 phase after high temperature oxidation, which indicates that the thermal stability of the oxide layer was not enhanced [18]. Therefore, a third element, Nb or Al, is typically added to MoSi2SiC composites to enhance the thermal stability of the oxide layer and the adherence of the oxide layer to the matrix. These elements are chosen because the oxides of the Nb and Al elements have excellent
Corresponding author. E-mail address: [email protected] (J.-L. Fan). These authors contributed equally to this work.
http://dx.doi.org/10.1016/j.ceramint.2017.09.010 Received 23 June 2017; Received in revised form 25 August 2017; Accepted 3 September 2017 Available online 04 September 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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Fig. 1. Illustration of the experimental process.
Fig. 2. (a) Backscatter electronic (BSE) micrographs of the in-situ HPing MoSi2 alloy; Inset is the enlarged view of dotted box. (b) Oxidation kinetics of MoSi2 alloy at 1200 °C in air environment. (c) SEM micrograph of oxide layer which in-situ formed in the surface of the MoSi2 alloy after 120 h oxidation. (d) SEM micrographs of the cross section of MoSi2 alloy after 120 h and five thermal shock cycle tests.
oxidation resistance, and are more stable than SiO2 at high temperature [19–22]. Based on this, the MoSi2 composites prepared with additional Nb, Al alloying and SiC reinforcement elements are expected to exhibit excellent thermal shock resistance. However, the thermal shock property of such composites has not yet been reported. In this work, we synthesized different compositions of (Mo0.94Nb0.06)(Si0.97Al0.03)2-x vol% SiC, and investigated the effects of Nb-Al-SiC content on the thermal shock properties of the materials in air atmosphere at 1200 °C. The effects of oxidation time and thermal shock cycle time on the thermal shock mechanism were studied by
performing oxidation kinetics followed by characterization of the microstructure of the composites and oxidized products. 2. Experimental procedures Molybdenum (2.0–2.5 µm, 99.9% purity), silicon (− 300 mesh, 99.9% purity), niobium (− 300 mesh, 99.9% purity), aluminum (100–200 mesh, 99.0% purity), and nano-sized SiC (40 nm, 99.9% purity) powders were used to form different MoSi2 composites: MoSi2, (Mo0.94Nb0.06)(Si0.97Al0.03)2-5 vol% SiC (NAS1), (Mo0.94Nb0.06)
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Fig. 3. (a) XRD patterns of the in-situ HPing MoSi2 alloy; (b) SEM micrograph of in-situ HPing MoSi2 alloy; (c–e) Energy-dispersive spectroscopy analysis of the different phase: (c) gray area 1, (d) white area 2, (e) black area 3 in (b).
(Si0.97Al0.03)2-10 vol% SiC (NAS2), and (Mo0.94Nb0.06)(Si0.97Al0.03)215 vol% SiC (NAS3). The starting mixtures were uniformly mixed by planetary ball milling for 240 min at 450 rounds per minute using an agate jar and agate balls with alcohol as the milling medium. Next, the ball-milled powders were dried and consolidated by in situ hot pressing (HPing) at 1400 °C and 27.5 MPa for 90 min under vacuum (< 8.0 × 10−1 Pa). The sintered plates (φ 50 mm × 5 mm) were cut into test pieces by wire-electrode cutting that measured 22 mm × 4 mm × 3 mm. The specimens surface was grounded and polished to a mirror finish and cleaned in acetone before the high temperature oxidation and thermal shock tests. The overall experimental process is illustrated in Fig. 1. The specimens were placed upon a corundum support and then heated at 1200 °C in air environment in an electric furnace to investigate the isothermal oxidation and thermal shock behavior. After a specified time of oxidation at 1200 °C, the specimens were removed from the furnace and quickly cooled to room temperature. The cumulative weight change of the specimens was reported as a function of exposure time by an electronic balance with a sensitivity of ± 0.1 mg. The specimens were then returned to the furnace for the next thermal cycle. The bending strength of the in-situ HPing and the heat-treated composites was tested by a three-point loading system, using a 18 mm span and a crosshead speed of 0.5 mm min−1 (according to test methods, ASTM C1161-2013 standard method for bending strength), and at least five specimens were tested for each condition. A scanning electron microscope (SEM; JEOLJSM-6360 and JSM-890
operated at 20 kV) with energy-dispersive X-ray analysis (EDX; Phoenix EDAX 2000) was employed to characterize the surface and cross-section microstructure of the composites. The phase compositions were identified using an X-ray diffractometer (XRD, Bruker D8 Advance, Karlsruhe, Germany). For the preparation of the transmission electron microscope (TEM) sample, a FEI Helios NanoLab (FIB/SEM) was employed. The morphology, structure, and elemental composition of the oxide layer were characterized using TEM analysis (FEI Tecnai F30 operated at 300 kV). 3. Results 3.1. Thermal shock failure of pure MoSi2 alloy The backscatter electron micrograph (BSE) of the HPing alloy is shown in Fig. 2a and reveals that the microstructure of pure MoSi2 consists of three kinds of phases (gray phase, white phase, and black phase). Fig. 3a shows the XRD patterns of the HPing MoSi2 alloy. It can be seen that MoSi2 is the main phase in the alloy, and miner amount of Mo5Si3 is also found. The EDX results (Fig. 3c–e) show that the gray phase is MoSi2, the white phase is Mo5Si3, and the black phase is SiO2. The EDX results are consistent with XRD analysis. Only SiO2 phase cannot be identified by XRD analysis, which is due to that the small amount of SiO2 has lower mass absorption coefficient (36.41 cm2 g−1) than MoSi2 dose (124 cm2 g–1) [23,24]. The formation of SiO2 and Mo5Si3 is due to that the raw powders containing a small volume of
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Fig. 4. (a) XRD pattern on the surface oxide layer of MoSi2 alloy after 24 h oxidation at 1200 °C and one thermal shock cycle. (b) SEM micrograph of MoSi2 after oxidization at 1200 °C for 24 h and one thermal shock cycle, and (c) EDX analysis of (b).
oxygen, and sintering such powders results in a small fraction of amorphous silica in the final product. Consumption of Si to form SiO2 renders the materials slightly Mo rich leading to the formation of Mo5Si3. This observation has been reported previously [25]. Fig. 2b shows the oxidation kinetics of MoSi2 alloy in air at 1200 °C. Obviously, the thermal shock behavior of the MoSi2 alloy can be divided into three sections marked as I, II, and III by analysis of the oxidation curve. In section I (during the first thermal shock), the SiO2 layer determined by EDX and XRD (as shown in Fig. 4) was formed, which provides shortterm protection of the matrix (< 72 h, in section II) (Fig. 5a–c). As the thermal shock cycle was increased (in section III), the mass of the MoSi2 alloy was decreased sharply. At the same time, the surface microstructure of the oxide layer was changed greatly, a process dominated by the failure of the SiO2 oxide layer (Fig. 5d, e). After 120 h oxidation and five thermal shock cycles, the surface oxide layer was disintegrated into fragments (as show in Fig. 2c–d), due to the large mismatch of the coefficient of thermal expansion (CTE) and the thermal conductivity (TC) between the SiO2 layer and the MoSi2 matrix [14,26–28]. These cracks can provide channels that allow oxygen to permeate into the inner body and lead to the continuous volatilization of MoO3 [29], consistent with the greater mass loss during oxidation (as shown in Fig. 2b). 3.2. Microstructure and thermal shock properties of NAS1, NAS2, and NAS3 composites As shown in Fig. 6a, with the addition of Nb-Al-SiC elements, there were obvious homogeneous large black phases. Compared with the
insert picture in Figs. 2a and 6a, it can be seen that the content of small black phase (SiO2) and white phase (Mo5Si3) decreased in the MoSi2 composites. The scanning transmission electron microscopy (STEM) image of the NAS1 composite (Fig. 6b) shows that Al elements were consumed into Al2O3 [30], and Nb elements were distributed in the interior of the MoSi2 grains, as determined by EDX and shown in Fig. 7, which can increase the adherence of the oxide layer to the MoSi2 matrix [31]. The large black phases (Fig. 6a) are SiC nanocluster particles (Fig. 6c) that will decrease the thermal stress between the oxide layer and the matrix in a high temperature oxidation environment by increasing the thermal conductivity [32] and decreasing the CTE value of the matrix [18]. Compared with pure MoSi2 alloy, the NAS1, NAS2, and NAS3 composites exhibited excellent thermal shock performance. Fig. 6e shows the mass gain per unit area as a function of oxidation time for the composites oxidized at 1200 °C in air for 120 h. During the oxidation test, the weight of the composites increased to follow a near-parabolic law. The composites exhibited a much smaller weight gain than the pure MoSi2 alloy. In order to approximate the oxidation rate constants for a quantitative comparison, the relatively parabolic section of the curve was fitted to the parabolic rate equation given by [33]:
(ΔW / A)2 = kp t + C Where △W is the mass gain, A is the specimen's surface area, kp is the oxidation rate constant, and t is the oxidation time. The results are shown in Fig. 8. With the increase of Nb-Al-SiC concentration, the composites exhibited a strong trend for reduced weight gain and decreased oxidation
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Fig. 5. SEM images of surface of oxide layer of the MoSi2 alloy at different times. (a) 24 h; (b) 48 h; (c) 72 h; (d) 96 h; (e) 120 h.
Fig. 6. (a) BSE micrograph of the surface of NAS1composite; Inset is the enlarged view of dotted box; (b), (c) STEM image of NAS1composite, a selected area electron diffraction pattern recorded on dash line circle area is shown in the inset of (c); (d) STEM image of in-situ HPing NAS1 composite after bending test; (e) Oxidation kinetics of NAS1-3 composites; (f)–(h) SEM micrographs of oxide layer of NAS1(f), NAS2(g) and NAS3(h) composites.
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Fig. 7. (a)TEM analysis of in-situ HPing NAS1 composite; (b) EDX analysis of particle A in (a); (c) EDX analysis of particle B in (a).
analyses of the surface oxide layer of NAS1, NAS2, and NAS3 composites after 120 h oxidation at 1200 °C and five thermal shock cycles are shown in Fig. 9a–c. The results show that the dense glassy layers are composed of SiO2, fiber structure Nb2O5, and mullite (3Al2O3·2SiO2) phase. The bright-field TEM image of the oxide layer cross-section revealed that the fibrous Nb2O5 appeared as a single crystal with a polygon structure embedded in the surface of the oxide layer and the mullite phase was nano-sized and located in the interface of the oxide layer and matrix (Fig. 9d–f). Additionally, the three phases were immiscible, which indicated that a multi-component oxide layer was synthesized. 4. Discussion
Fig. 8. Wight gain2 plotted as a function of exposure time.
Table 1 Calculation results of volume ratio of SiO2 matrix and fiber structure Nb2O5 in surface oxide layer of NAS1-3 composites after five thermal shock cycle through Image J2x software. Composites
NAS1 NAS2 NAS3
Volume ration SiO2 matrix
Fiber structure Nb2O5
88.8 86.3 83.8
11.2 13.7 16.2
rate, not only in the early stage but also in the stable stage of oxidation. Furthermore, the fragmentized oxide layer present in the surface of the pure MoSi2 alloy (Fig. 2c, d) disappeared completely. A fiber-structured dense oxide layer was formed in-situ on the surface of the NAS1, NAS2, and NAS3 composites, and this layer may effectively prevent the matrix from further oxidation during the 120 h treatment at 1200 °C and five thermal shock cycles (Fig. 6f–h). At the same time, the grain sizes of the fiber structure phase became finer and the number of the phase increased gradually. The volume ratios of the fiber structure phase in the oxide layer of NAS1-NAS3 composites were 11.2%, 13.7%, and 16.2%, respectively (Table 1). We now present details of the phase composition and microstructure of the surface and cross section fibrous structures oxide layer formed insitu on the surface of the NAS1-3 composites. This will allow us to explain the excellent thermal shock resistance of MoSi2 composites resulting from the addition of Nb-Al-SiC elements. The XRD and EDX
In order to fully understand the anti-oxidation and anti-thermal shock mechanisms, the thermodynamic and dynamic behaviors of the oxidation of the composites were next studied. Thermodynamic calculations (delta-Gibbs free energy of reaction) show that once the composites were put into the furnace, oxidation (reaction (1)–(7)) can occur and some gases by-products are generated, such as CO2, CO, MoO3, whose emission does not increase the protection of the composites [34]. Depending on the thermodynamic calculations, MoSi2 may be oxidized into Mo5Si3 and SiO2 (reaction (1)), and then Mo5Si3 was oxidized into MoO3 and SiO2 (reaction (2)) [35]. However, Ibano [36] reported that simultaneous oxidation of molybdenum and silicon occurs at the initial stage of oxidation at ambient pressure. Zhu [37] also reported that simultaneous oxidation occurs at the initial stage of oxidation (< 1 h) in all temperature ranges from 900 °C to 1500 °C in an air environment. Thus, for the HPing composites, the simultaneous oxidation of molybdenum and silicon occurred at the initial stage of the oxidation process in the air environment, as shown in reaction (3). Under constant temperature and pressure, the enthalpy change of reaction is equal to the heat evolved by the reaction [38]. At the same time, the more heat evolved by the oxidation reaction, the higher the affinity of the component for oxygen. According to the calculated enthalpy change of the reaction (3,4), the heat values of exothermic reactions during oxidation are in the following sequence: Qreaction (4) > Qreaction (3) , indicating that the affinity of the Nb for oxygen is higher than that of Si. Prigogine et al. [39] reported that the thermodynamics of irreversible phenomena, based mainly on the concept of entropy production, leads to linear and homogeneous relationships between affinities and reaction rates. In such composites, fiber structured Nb2O5 can be preferentially formed and then could be enwrapped by SiO2 during oxidation, indicating that fiber-structured Nb2O5 was present in the top of the oxide layer (as shown in Fig. 9d). Based on these results, a schematic illustration of the phase structure evolution during the oxidation
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Fig. 9. (a) XRD analysis on the surface of oxide layer of NAS1-3 composites at 1200 °C after heat treat; (b) EDX analysis of A region in Fig. 6f; (c) EDX analysis of B region in Fig. 6f; (d) Cross-sectional FIB microstructure of the oxide layer formed in the surface of the matrix with NAS3 which were extracted from the regions of the sample indicated in Fig. 6h using focused ion beam techniques; (e) Selected area electron diffraction (SAED) pattern of the polygon in the Fig. 9d; (f) Typical diffraction pattern of the nano crystalline region in the Fig. 9d.
Fig. 10. Schematic illustration of the formation process of the multi-component oxide layer upon oxidation and thermal shock test for different stages.
of NAS1, NAS2, and NAS3 composites is shown in Fig. 10. It has been reported that the Nb2O5 and mullite are stable oxides, with high coefficient of thermal expansion (CTE) values (5.3 × 10−6/K [40], 5.3 × 10−6/K [41]) and a high degree of oxidation resistance [22]. The uniformly distributed fiber structured Nb2O5 and nano-sized mullite should increase the viscosity of the multi-component SiO2 glass layer and reduce the oxygen diffusion rate [20], thus improving the oxidation resistance of the oxide layer. Therefore, with the increased fiberstructure Nb2O5 content, the oxygen diffusion rate will decrease gradually. This is consistent with the observed thinner oxide layer (the oxide layer thickness of the heat-treated NAS1-3 composites was 21.60 µm, 7.30 µm, and 6.19 µm, respectively, as shown in Fig. 11), and consistent with the oxidation kinetics shown in Fig. 6e. Additionally, nano-sized mullite phase was present in the interface of the oxide layer and matrix and can reduce the stress concentration [42,43] and retard the propagation of micro cracks [44,45]. Consequently, the Nb-Al-SiC elements are expected to significantly improve the thermal shock resistance of the MoSi2 composites.
The possible reactions during oxidation at 1200 °C in air environment are：
5/7MoSi2 (s) + O2 (g) → 1/7Mo5 Si3 (s) + SiO2 (s)
ΔG(1400K) = −615 kJ mol−1 ΔH(1400K) = −849 kJ mol−1
(1)
2/21Mo5 Si3 +O2 (g) → 10/21MoO3 (g) + 6/21SiO2 (s)
ΔG(1400K) = −291.87 kJ mol−1 ΔH(1400K) = −474 kJ mol−1 (2)
2/7MoSi2 (s) + O2 (g) → 2/7MoO3 (g)
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+ 4/7SiO2 (s)
ΔG(1400K) = −460 kJ mol−1 ΔH(1400K) = −666 kJ mol−1 (3)
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Fig. 11. The effect of Nb-Al-SiC elements on the thickness of oxide layer. (a)–(c) Scanning electron micrographs of the fracture plane of NAS1(a), NAS2(b) and NAS3(c) composites after oxidized at 1200 °C for 120 h.
This decrease in the strength is attributed to the formation of cracks in the thick oxide layer [46]. After introducing the Nb-Al-SiC elements into the MoSi2 matrix, the bending strength of the composites increased gradually and even exceeded that of the initial HPing composites and reached a maximum of 496 MPa (NAS3 composite). This increase was due to the formation of a multi-component dense oxide layer that could heal the surface flaws resulting from sample processing and machining [47]. At the same time, the fiber structured Nb2O5 and nano-sized mullite present in the oxide layer can act as reinforcements (like whiskers [48], nanoparticles [49], and nanoclusters [50]), and retard the propagation of cracks. And heat treatment at high temperature could relieve residual stress of in-situ MoSi2 composites and result in improving mechanical properties [51]. This also needs to be investigated in future studies. Sciti et al. [52] reported that the bending strength of ZrB2-MoSi2 increased by up to 3% after oxidation at 1200 °C for 100 h, which was quite surprising since microstructural analyses revealed remarkable microstructural degradation. Krnel et al. [53] reported that after oxidized at 1200 °C for 100 h, the AlN–SiC–MoSi2 composites showed little strength improvement.
Fig. 12. Bending strength of in-situ HPing composites and heat treated composites.
4/5Nb(s) + O2 (g) → 2/5Nb2 O5 (s)
ΔG(1400K) = −520 kJ mol−1 ΔH(1400K) = −700 kJ mol−1
(4) 5. Conclusions
1/2SiC(s) + O2 (g) → 1/2SiO2 (s) + 1/2CO2 (g) ΔG(1400K) = −498 kJ mol−1
(5)
The effects of addition of Nb-Al-SiC elements on the microstructure and high temperature thermal shock resistance of MoSi2 composites were investigated. The following conclusions were derived:
2/3SiC(s) + O2 (g) → 2/3SiO2 (s) + 2/3CO (g) ΔG(1400 K) = −556 kJ mol−1
(6)
3Al2O3 (s) + 2SiO2 (s) → 3Al2O3·2SiO2 (s) ΔG(1400 K) = −20 kJ mol−1 (7) To investigate the mechanical properties of MoSi2 composites before and after high temperature oxidation and five thermal shock cycles, the three-point bending test was performed and the results are shown in Fig. 12. With increased amount of Nb-Al-SiC elements, the bending strength of the in-situ HPing composites increased gradually. This is caused by inhibition of dislocation by the Al2O3, Nb particles, and SiC nanocluster (as shown in Fig. 6c–d) and results in a strengthening effect [25]. After 120 h oxidation and five thermal shock cycles (△T = 1200 °C), the retained bending strength of MoSi2 and the NAS1, NAS2, NAS3 composites exhibited different trends. The bending strength of the pure MoSi2 after heat treatment decreased from 330 MPa (initial HPing alloy) to 258 MPa with ~ 78% retention in bending strength.
(a) The elements of Nb-Al2O3-SiC were dispersed in the MoSi2 matrix homogeneously after in-situ hot pressing. (b) After the thermal shock test, the fragmentized oxide layer present in the surface of pure MoSi2 alloy disappeared completely and a dense multi-component oxide layer was formed in-situ on the surface of the NAS1, NAS2, and NAS3 composites. This dense multi-component oxide layer was composed of SiO2, fiber-structured Nb2O5, and nano-sized mullite. (c) With increased Nb-Al-SiC content, the bending strength of the composites after hot treat increased gradually due to the formation of multi-component of oxide layer.
Acknowledgements This research was supported by the National Natural Science Fund Project of China (51534009).
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References [1] M. Yamaguchi, H. Inui, K. Ito, High-temperature structural intermetallics, Acta Mater. 48 (2000) 307–322. [2] J.H. Perepezko, The hotter the engine, the better, Science 326 (2009) 1068–1069. [3] S.P. Murarka, D.B. Fraser, A.K. Sinha, H.J. Levinstein, Refractory silicides of titanium and tantalum for low-resistivity gates and interconnects, IEEE J. Solid-State Circuit 15 (1980) 474–482. [4] A.K. Vasudevan, J.J. Petrovic, A comparative overview of molybdenum disilicide composites, Mater. Sci. Eng. A 155 (1992) 1–17. [5] P.J. Meschter, D.S. Schwartz, Silicide-matrix materials for high-temperature applications, JOM 41 (1989) 52–55. [6] R. Mitra, Mechanical behavior and oxidation resistance of structural silicides, Int. Mater. Rev. 51 (2006) 13–64. [7] T. Feng, H.J. Li, Q.G. Fu, X. Yang, H. Wu, High-temperature erosion resistance and aerodynamic oxidation mechanism of multi-layer MoSi2–CrSi2–Si/SiC coated carbon/carbon composites in a wind tunnel at 1873 K, Carbon 50 (2012) 2171–2178. [8] Y. Wang, D.Z. Wang, J.H. Yan, A.K. Sun, Preparation and characterization of molybdenum disilicide coating on molybdenum substrate by air plasma spraying, Appl. Surf. Sci. 284 (2013) 881–888. [9] L. Su, O. Lu-Steffes, H. Zhang, J.H. Perepezko, An ultra-high temperature Mo-Si-B based coating for oxidation protection of NbSS/Nb5Si3 composites, Appl. Surf. Sci. 337 (2015) 38–44. [10] J.F. Huang, H.J. Li, X.R. Zeng, X.B. Xiong, Y.W. Fu, Progress on the oxidation protective coating of carbon-carbon composites, New Carbon Mater. 20 (2005) 373–379. [11] S.P. Lee, J.K. Lee, C.H. Leem, Cyclic thermal shock properties of MoSi2 intermetallic materials, Int. J. Mod. Phys. B 24 (2010) 2922–2927. [12] S.L. Lee, Y.H. Kim, Y.S. Shin, J.L. Lee, J.Y. Park, Thermal shock damage and mechanical property of MoSi2 materials by NDE technique, Key Eng. Mater. 321–323 (2006) 930–933. [13] Y.L. Zhang, H.J. Li, X.F. Qiang, K.Z. Li, S.Y. Zhang, C/SiC/MoSi2–simultilayer coatings for carbon/carbon composites for protection against oxidation, Corros. Sci. 53 (2011) 3840–3844. [14] Q.G. Fu, H.J. Li, Y.J. Wang, K.Z. Li, X.H. Shi, B2O3 modified SiC–MoSi2 oxidation resistant coating for carbon/carbon composites by a two-step pack cementation, Corros. Sci. 51 (2009) 2450–2454. [15] M.G. Hebsur, Development and characterization of SiC(f)/MoSi2–Si3N4(p) hybrid composites, Mater. Sci. Eng. A 261 (1999) 24–37. [16] H. Chen, M. Suzuki, S. Sodeoka, T. Inoue, K. Ueno, New approach to MoSi2/SiC intermetallic-ceramic composite with B4C, J. Mater. Sci. 36 (2001) 5773–5777. [17] J.F. Huang, B. Wang, H.J. Li, M. Liu, L.Y. Cao, C.Y. Yao, A MoSi2/SiC oxidation protective coating for carbon/carbon composites, Corros. Sci. 53 (2011) 834–839. [18] J. Liu, Q.M. Gong, Y. Shao, D.M. Zhuang, J. Liang, In-situ fabrication of MoSi2/ SiC–Mo2C gradient anti-oxidation coating on Mo substrate and the crucial effect of Mo2C barrier layer at high temperature, Appl. Surf. Sci. 308 (2014) 261–268. [19] M.M. Opeka, I.G. Talmy, J.A. Zaykoski, Oxidation-based materials selection for 2000 C+ hypersonic aerosurfaces: theoretical considerations and historical experience, J. Mater. Sci. 39 (2004) 5887–5904. [20] Z.H. Hao, W. Sun, X. Xiong, Z.K. Chen, Y.L. Wang, Effects of Ti/Al addition on the microstructures and ablation properties of Cf/C–MoSi2–SiC composites, J. Eur. Ceram. Soc. 36 (2016) 457–464. [21] J. Cheng, S. Yi, J.S. Park, Oxidation behavior of Nb-Si-B alloys with the NbSi2 coating layer formed by a pack cementation technique, Int. J. Refract. Met. Hard Mater. 41 (2013) 103–109. [22] L. Ingemarsson, M. Halvarsson, J. Engkvist, T. Jonsson, K. Hellström, L.G. Johansson, J.E. Svensson, Oxidation behavior of a Mo(Si,Al)2-based composite at 300–1000 °C, Intermetallics 18 (2010) 633–640. [23] Q. Lu, G.M. Zhu, X.H. Wang, Z. Wang, P.Z. Feng, Effects of raw materials on synthesis, microstructure and properties of MoSi2-10 vol% SiC composites, Trans. Indian Ceram. Soc. 75 (2016) 33–39. [24] B.D. Cullity, Elements of X-ray Diffraction, 1st edition, Addison-Wesley Publishing, USA, 1959, pp. 466–467. [25] Q. Lu, G.M. Zhu, X.H. Wang, J.L. Fan, Synthesis and mechanical properties of (Mo0.94Nb0.06)(Si0.97Al0.03)2-x vol% SiC composites fabricated via SHS-HP, Ceram. Int. 43 (2017) 3424–3430. [26] M.C. Chu, S.J. Cho, Y.C. Lee, H.M. Park, Crack healing in silicon carbide, J. Am.
Ceram. Soc. 87 (2004) 490–492. [27] C. Hu, M. Kiene, P.S. Ho, Thermal conductivity and interfacial thermal resistance of polymeric low k films, Appl. Phys. Lett. 79 (2001) 4121–4123. [28] W.W. Beaver, A.J. Stonehouse, R.M. Paine, Development of intermetallic compounds for aerospace applications, in: Plansee Proceedings 5th, 1965, pp. 682–700. [29] Y.L. Zhang, H.J. Li, X.Y. Yao, K.L. Li, X.F. Qiang, Oxidation protection of C/SiC coated carbon/carbon composites with Si-Mo coating at high temperature, Corros. Sci. 53 (2011) 2075–2079. [30] I.M. Allam, H.C. Akuezue, D.P. Whittle, Influence of small Pt additions on Al2O3 scale adherence, Oxid. Met. 14 (1980) 517–530. [31] R. Mitra, K. Sadananda, C.R. Feng, Effect of microstructural parameters and Al alloying on creep behavior, threshold stress and activation volumes of molybdenum disilicides, Intermetallics 12 (2004) 827–836. [32] J.Y. Qiu, Y. Hotta, K. Watari, K. Mitsuishi, Enhancement of densification and thermal conductivity in AlN ceramics by addition of nano-sized particles, J. Am. Ceram. Soc. 89 (2006) 377–380. [33] D. Monceau, B. Pieraggi, Determination of parabolic rate constants from a local analysis of mass-gain curves, Oxid. Met. 50 (1998) 477–493. [34] A.A. Sharif, High-temperature oxidation of MoSi2, J. Mater. Sci. 45 (2010) 865–870. [35] Q. Zhu, K. Shobu, Y. Zeng, T. Watanabe, Oxidation behavior of hot-pressed MoSi2TiC composite, J. Mater. Sci. 36 (2001) 313–319. [36] A. Ibano, K. Yoshimi, A. Yamauchi, R. Tu, K. Maruyama, K. Kurokawa, T. Goto, High temperature oxidation behavior of MoSi2 under low pressure atmosphere, Mater. Sci. Forum 561 (2007) 427–430. [37] G.M. Zhu, X.H. Wang, Q. Lu, G.Z. Wu, P.Z. Feng, High-temperature crack-healing behavior and strength recovery of (MoNb)Si2, Appl. Surf. Sci. 343 (2015) 41–48. [38] R.A. Alberty, Standard Gibbs free energy, enthalpy, and entropy changes as a function of pH and pMg for several reactions involving adenosine phosphates, J. Biol. Chem. 244 (1969) 3290–3302. [39] I. Prigogine, P. Outer, C.L. Herbo, Affinity and reaction rate close to equilibrium, J. Phys. Chem. 52 (1948) 321–331. [40] W.R. Manning, O. Hunter, F.W. Calderwood, D.W. Stacy, Thermal expansion of Nb2O5, J. Am. Ceram. Soc. 55 (1972) 342–347. [41] Y.S. Toulouhain, R.K. Kirby, R.E. Talor, T.Y. Lee, Thermophysical Properties of Matter 13 IFI-Plenum, New York, 1977, p. 25. [42] Q.G. Fu, Y.C. Shan, S.Q. Wen, H.J. Li, Y.H. Chu, X.S. Liu, Oxidation pre-treatment to improve the mechanical property and oxidation resistance of Si–Mo–Cr coated C/C composites, Corros. Sci. 83 (2014) 103–110. [43] S.M. Choi, H. Awaji, Nanocomposites – a new material design concept, Sci. Technol. Adv. Mater. 6 (2005) 2–10. [44] T. Aoki, H. Hatta, T. Hitomi, H. Fukuda, I. Shiota, SiC/C multi-layered coating contributing to the antioxidation of C/C composites and the suppression of throughthickness cracks in the layer, Carbon 39 (2001) 1477–1483. [45] Q.G. Fu, J.Y. Jing, H.J. Li, X.X. Jin, L. Zhuang, L. Li, Design of an inlaid interface structure to improve the oxidation protective ability of SiC-MoSi2-ZrB2 coatings for C/C composites, Ceram. Int. 42 (2016) 4212–4220. [46] G.M. Zhu, X.H. Wang, Z. Sun, P.Z. Feng, F. Akhtar, Effect of annealing environment on the crack healing and mechanical properties of (Mo0.97Nb0.03)(Si0.97Al0.03)2, J. Alloy. Compd. 634 (2015) 109–114. [47] W.M. Guo, G.J. Zhang, Oxidation resistance and strength retention of ZrB2–SiC ceramics, J. Eur. Ceram. Soc. 30 (2010) 2387–2395. [48] Y. Mao, L. Peng, S. Wang, L.X. Xi, Microstructural characterization of graphite/ CuCrZr joints brazed with CuTiH2 Ni-based fillers, J. Alloy. Compd. 716 (2017) 81–87. [49] Z. Zhan, R. Xu, X. Zheng, Q. Fu, M. Wu, Y. Lei, Effective approach to strengthen plasmon resonance localized on top surfaces of Ag nanoparticles and application in surface-enhanced Raman spectroscopy, Nanotechnology 27 (2016) 445301. [50] S.S. Xu, et al., Independence of work hardening and precipitation strengthening in a nanocluster strengthened steel, J. Alloy. Compd. 712 (2017) 573–578. [51] P. Peralta, R. Dickerson, J.R. Michael, Residual thermal stresses in MoSi2–Mo5Si3 insitu composites, Mater. Sci. Eng. A 261 (1999) 261–269. [52] D. Sciti, M. Brach, A. Bellosi, Long-term oxidation behavior and mechanical strength degradation of a pressurelessly sintered ZrB2–MoSi2 ceramic, Scr. Mater. 53 (2005) 1297–1302. [53] K. Krnel, D. Sciti, A. Bellosi, Influence of long term oxidation on the microstructure, mechanical and electrical properties of pressureless sintered AlN–SiC–MoSi2 ceramic composites, J. Eur. Ceram. Soc. 23 (2003) 3135–3146.
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