In situ synthesis, microstructure, mechanical

Int. Journal of Refractory Metals and Hard Materials 46 (2014) 101–108

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In situ synthesis, microstructure, mechanical properties and thermal shock resistance of (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites Jian Yang a,⁎, Dongdong Guo a, Lei Yu a, Limei Pan a, Tai Qiu a, Jingxian Zhang b a b

College of Materials Science and Engineering, Nanjing Tech University, Nanjing 210009, China State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China

a r t i c l e

i n f o

Article history: Received 6 March 2014 Accepted 7 May 2014 Available online 3 June 2014 Keywords: (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites In situ synthesis Mechanical properties Thermal physical properties Thermal shock resistance

a b s t r a c t Dense (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites with adjustable content of (ZrB2 + SiC) reinforcements (0–30 vol.%) were prepared by in situ hot-pressing. The microstructure, room and high temperature mechanical and thermal physical properties, as well as thermal shock resistance of the composites were investigated and compared with monolithic Zr2[Al(Si)]4C5 ceramic. ZrB2 and SiC incorporated by in situ reaction significantly improve the mechanical properties of Z2[Al(Si)]4C5 by the synergistic action of many mechanisms including particulate reinforcement, crack deflection, branching, bridging, “self-reinforced” microstructure and grain-refinement. With (ZrB2 + SiC) content increasing, the flexural strength, toughness and Vickers hardness show a nearly linear increase from 353 to 621 MPa, 3.88 to 7.85 MPa·m1/2, and 11.7 to 16.7 GPa, respectively. Especially, the 30 vol.% (ZrB2 + SiC)/Zr2 [Al(Si)]4C5 composite retains a high modulus up to 1511 °C (357 GPa, 86% of that at 25 °C) and superior strength (404 MPa) at 1300 °C in air. The composite shows higher thermal conductivity (25–1200 °C) and excellent thermal shock resistance at ΔT up to 550 °C. Superior properties render the composites a promising prospect as ultra-high-temperature ceramics. © 2014 Elsevier Ltd. All rights reserved.

Introduction Zr2[Al(Si)]4C5 is a new quaternary layered carbide formed by adding Al and Si to ZrC, which shows excellent properties, such as high melting point, retained strength at elevated temperatures, and relatively good thermal shock resistance [1–5]. Especially, its Young's modulus decreases slowly with increasing temperature and can remain 81% of the room temperature modulus to 1580 °C. Therefore, Zr2[Al(Si)]4C5 is believed to be a promising candidate for ultra-high-temperature ceramics (UHTCs) used in re-entry and hypersonic vehicles [3–5]. However, compared with other high-temperature structural materials, its hardness, toughness and oxidation resistance are still unsatisfactory [5], which limits its application. It is well known that incorporating a second phase is an effective method to overcome these weaknesses of structural ceramics. He et al [6,7] and Chen et al [8,9] prepared SiC reinforced Zr2 [Al(Si)]4C5 composites, which show superior comprehensive performance compared to Zr2[Al(Si)]4C5 ceramic. In recent years, a synergy mechanism between two or more strengthening and toughening phases/methods in ceramic-matrix composites has attracted much attention, which shows more significant strengthening and toughening effects than individual toughening phases/methods. Yang et al [10] and Song et al [11] prepared (TiB2 + TiC)/Ti3SiC2 composites and (TiB2 + SiC)/Ti3SiC2 composites by in ⁎ Corresponding author. Tel.: +86 25 83587262; fax: +86 25 83587268. E-mail address: [email protected] (J. Yang).

http://dx.doi.org/10.1016/j.ijrmhm.2014.05.014 0263-4368/© 2014 Elsevier Ltd. All rights reserved.

situ hot-pressing sintering respectively, in which columnar TiB2 grains and equiaxial TiC or SiC grains incorporated by in situ reaction improve the mechanical, tribological properties and oxidation resistance of Ti3SiC2 matrix dramatically. Comparing with traditional preparation process, in situ synthesis shows many advantages in fabrication of composite ceramics, such as uniform distribution and fine grain size of reinforcements, clean and tightly bonded grain boundary, simplified process and excellent properties of composites [12]. As the typical representative of refractory transitional metal borides, ZrB2 shows high melting point, strength, elastic modulus, resistance to chemical attack and can maintain high strength at elevated temperatures. With the addition of SiC as a second phase, the ZrB2–SiC composites show improved strength, fracture toughness and superior high temperature oxidation resistance [13–15]. As a result, ZrB2–SiC has been widely researched as the main UHTCs. In consideration of the characteristics of Zr2[Al(Si)]4C5 and ZrB2–SiC, it can be easily inferred that ZrB2–SiC will be ideal binary reinforcements for Zr2 [Al(Si)]4C5 ceramics and the composites will show excellent comprehensive properties and promising prospect as UHTCs. In our previous work [16], with an aim to design and develop novel UHTCs, (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite ceramics were successfully prepared by in situ hot-pressing based on thermodynamic analysis. In the present work, (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites with variable (ZrB2 + SiC) content were prepared. The microstructure, room and especially high temperature mechanical and thermal physical properties as well as thermal shock resistance of the composites were systematically

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J. Yang et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 101–108

investigated and compared with monolithic Zr2[Al(Si)]4C5 ceramic. The results indicate the composites a promising prospect as UHTCs.

Experimental procedure Material preparation and microstructure characterization Zirconium hydride (99.9%, 0.04 mm), silicon (99.9%, 0.04 mm), aluminum (99.9%, 0.08 mm), graphite (99.9%, 0.08 mm), and B4 C powders (99%, 0.35 mm) were used as starting materials. They were precisely weighed according to the desired composition and dry-mixed by ball milling for 24 h in a polyethene jar filled with Argon. Then the powder mixtures were put into a graphite mold pre-sprayed with a layer of BN under the pressure of 3 MPa. The compacted mixture was first heated at a rate of 20 °C/min from room temperature to 700 °C and held for 1 h, then was hot-pressed at 1850 °C for 1 h under the pressure of 25 MPa in Ar atmosphere. For comparison, monolithic Zr2[Al(Si)]4C5 ceramic was also prepared following the same process using zirconium hydride, silicon, aluminum and graphite powders as raw materials. The bulk density and apparent porosity of the samples were determined by the Archimedes method. Phase identification was conducted via an X-ray diffractometer (XRD) with Cu Kα radiation (ARL X'TRA, Switzerland). The microstructure of polished and fracture surfaces of the samples were observed by scanning electron microscope (SEM) (JEOL, JSM-5900), coupled with an energy-dispersive spectroscope (EDS) for chemical analysis. A transmission electron microscope (TEM; JEOL, JEM-2010 UHR) equipped with an EDS system was used to characterize the microstructure of the samples. Thin-foil specimens for TEM observations were prepared by slicing, mechanical grinding to b60 μm and ion milling at 5.0 kV.

Mechanical property test The sintered materials were cut, ground and polished into strips with the sizes of 3 mm × 4 mm × 40 mm and 2 mm × 4 mm × 20 mm for strength and toughness tests, respectively. Three point bending method was applied to measure room temperature bending strength using a span of 30 mm with a crossing speed of 0.5 mm/min. The fracture toughness (K IC ) was measured using single-edge notched beam (SENB) method with a crosshead speed of 0.05 mm/min. The notch machined by inner circle cutting machine was 4 mm in depth and ≤ 0.28 mm in width with a notch radius of 0.06 mm. Five samples were used in both strength and toughness measurements. Indentation tests were employed to determine Vickers hardness at room temperature using a Vickers' diamond indenter. The indention parameters were made using a load of 9.8 N with a dwell time of 10 s. Five measurement runs were carried out to determine the average value. The hightemperature three-point flexural strength of the samples was tested in a high-temperature universal testing machine in the air. To reveal the temperature dependence of Young's modulus of the samples, a rectangular beam-like sample with dimension of 3 mm × 15 mm × 40 mm was suspended in the nodes of their first bending vibration mode, and measured in a graphite furnace (HTVP 1750 °C, IMCE, Diepenbeek, Belgium) at a heating rate of 3 °C/min in vacuum on the order of 10−1 Pa. The vibration signal, captured by a laser vibrometer, was analyzed with the resonance frequency and damping analyzer [17]. The Young's modulus was calculated from the flexural resonant frequen 2 mf L3 T 1 , with m, L, cy, ff, according to ASTM E 1876–97: E ¼ 0:9465 b f t3 b and t, the sample weight, length, width, and thickness, respectively. T1 is a correction factor, depending on the Poisson's ratio and the thickness/ length ratio.

Thermal physical property measurement The coefficient of thermal expansion (CTE) was measured using a Setsys-24 thermal mechanical analyzer (TMA) (Linseis L75 Platinum Series) from 100 to 1200 °C with a heating rate of 5 °C·min−1 under flowing Ar. The dimension of the sample for CTE measurement was 20 mm × 3.5 mm × 3.5 mm. The specific heat capacity cp and thermal diffusivity α (sample size Φ12.7 mm × 2 mm) were measured using laser-flash technique from 25 to 1200 °C with a heating rate of 20 °C/min under flowing Ar. Prior to the test, the sample was sprayed with a thin layer of colloidal graphite approximately 10-μm-thick to ensure complete and uniform absorption of the laser pulse. Three measurements were taken at each temperature and the data was calculated using software (Anter FL5000). The thermal conductivity, k, was calculated according to the following equation: k = ρ ⋅ Cp ⋅ α, in which ρ, Cp and α were the sample density, the specific heat capacity and the thermal diffusivity, respectively. Thermal shock experiments Water quenching tests under air atmosphere were conducted to evaluate the thermal shock resistance of the samples. The test specimens were heated up to the testing temperatures in air atmosphere and held for 10 min, which was followed by quenching into water bath quickly. The time taken for transferring the specimens from the furnace to the water bath was less than 1 s. The temperature of the water bath was maintained at about 20 °C by adjusting the cooling water flow. The water quenching temperature differences (exposure temperature minus the water bath temperature, ΔT) were 100, 200, 300, 400, 600, and 800 °C, respectively. Before water quenching, all specimens were ground and polished with diamond slurries down to a 1 μm finish and the edges of all specimens were chamfered to minimize the effect of stress concentration due to machining flaws. The residual flexural strength of the shocked specimens was measured under the same conditions as those of the unshocked ones. Residual strength ratio was also calculated, which is defined as the ratio of flexural strength after and before thermal shocking. Results and discussion During the in situ hot pressing process, Zr produced by decomposition of ZrH2 reacts with other raw materials as follows: 2Zr þ 3:6Al þ 0:4Si þ 5C→Zr2 ½AlðSiÞ4 C5

ð1Þ

2Zr þ Si þ B4 C→2ZrB2 þ SiC

ð2Þ

which has been confirmed by thermodynamic analysis and experiment results in our previous work [16]. The general reaction for the (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites is described as follows: ZrH2 þ Al þ Si þ C þ B4 C→Zr2 ½AlðSiÞ4 C5 þ ZrB2 þ SiC þ H2 :

ð3Þ

In this work, the composites with different volume fractions of (ZrB 2 + SiC) (10, 20, and 30 vol.%) were employed according to this reaction. Phase and microstructure analysis Fig. 1 shows the XRD patterns of (ZrB 2 + SiC)/Zr 2 [Al(Si)] 4 C 5 composites. For comparison, the XRD pattern of monolithic Zr 2 [Al(Si)] 4 C 5 ceramic was also given. It can be seen that three phases, Zr 2 [Al(Si)] 4 C 5 , ZrB 2 and SiC(4H) are identified in the composites and Zr 2 [Al(Si)] 4 C 5 is the major crystal phase. With the increase of (ZrB2 + SiC) content, the intensity of diffraction

J. Yang et al. / Int. Journal of Refractory Metals and Hard Materials 46 (2014) 101–108 Zr2 [Al(Si)4]C5 ZrB2 SiC

Intensity (arb. unites)

30vol.%(ZrB2+SiC)

20vol.%(ZrB2+SiC)

10vol.%(ZrB2+SiC]

Zr2 [Al(Si)]4C5

10

20

30

2θ θ/°

40

50

60

Fig. 1. XRD patterns of (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites with different (ZrB2 + SiC) content as well as Zr2[Al(Si)]4C5 for comparison.

peaks for ZrB2 and SiC increases gradually. It suggests that in situ (ZrB 2 + SiC)/Zr 2 [Al(Si)] 4 C 5 composites with different (ZrB 2 + SiC) content have been successfully prepared. Fig. 2 is the sintering properties of (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites. It can be seen that all the samples are fully dense with an apparent porosity lower that 0.5%. Both monolithic Zr2[Al(Si)]4C5 and the composites with (ZrB2 + SiC) content up to 30 vol.% can be densified by sintering at 1850 °C. The incorporation of (ZrB2 + SiC) and increasing their content show little influence on the sintering of the materials. Owing to higher density of ZrB2 and SiC compared with Zr2[Al(Si)]4C5 matrix, the bulk density of the materials increases gradually with the increasing (ZrB2 + SiC) content. Fig. 3 shows TEM and HRTEM images of the 30 vol.% (ZrB 2 + SiC)/Zr 2 [Al(Si)] 4 C 5 composite. The successful synthesis of the composites is furthermore confirmed by EDS analysis (the spectra are omitted). It can be seen that owing to the superiority of in situ synthesis, the grain boundaries are clean and clear, with no glass phase, which indicates a tight bonding. In addition, the SiC reinforcement formed by in situ reaction shows a nano-level grain size of about 60 nm, implying that Zr 2[Al(Si)]4C5-matrix nano composites have been obtained. Backscattered electron image of the polished surface of the composite is shown in Fig. 4. Based on the brightness and EDS results (omitted), three phases can be easily identified. The gray phase is Zr3[Al(Si)]4C6, the offwhite phase is ZrB2 and the white phase is SiC. This is in good agreement with XRD and TEM-EDS results. It also can be seen that ZrB2 and SiC show fine grain size and homogenous distribution in the composites. The ZrB2 is usually columnar grain with an aspect ratio of 3–4, displaying a preferential growth vertical to the c-axis. The 1.0

0.8

Bulk density Apparent porosity

4.6

0.6 4.4

0.4

Apparent porosity /%

Bulk density /g/cm3

4.8

0.2

4.2

0

5

10

15

20

25

30

(ZrB2+SiC) content /vol% Fig. 2. Sintering performance of the samples as a function of (ZrB2 + SiC) content.

103

elongated ZrB2 grains constitute a “self-reinforced” microstructure which might generate distinctive strengthening and toughening effects. Fig. 5 shows SEM images of the fracture surfaces of the composites with different (ZrB2 + SiC) contents. All the samples exhibit a dense microstructure, which is in agreement with the sintering performance results. The fracture mode of monolithic Zr2[Al(Si)]4C5 is mainly transgranular fracture and the fracture surface (Fig. 4(a)) is relatively flat. For the composites, however, the fracture surfaces (Fig. 4(b–d)) become rough and the fracture pattern changes to a mixture of intergranular and transgranular fracture modes. In addition, it also can be clearly seen that with the increase of (ZrB2 + SiC) content, the materials show smaller grain size, suggesting that the incorporation of (ZrB2 + SiC) significantly inhibits the grain growth of Zr2[Al(Si)]4C5 matrix. Room-and high-temperature mechanical properties Fig. 6 shows the room temperature mechanical properties of the materials as a function of (ZrB2 + SiC) content. It can be seen that with the content of (ZrB2 + SiC) increasing up to 30 vol.%, bending strength, fracture toughness and Vickers hardness of the materials show a nearly linear increase from 353 to 621 MPa, 3.88 to 7.85 MPa·m1/2, and 11.7 to 16.7 GPa, respectively. The flexural strength and fracture toughness of 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite are about two times of those of Zr2[Al(Si)]4C5. It can also be identified that the increase of Vickers hardness of the composites follows the linear relation Hv ðGPaÞ ¼ 11:93 þ 0:168V f :

ð4Þ

Table 1 gives the comparison of some mechanical properties between 30 vol.% (ZrB 2 + SiC)/Zr 2 [Al(Si)] 4 C 5 composite and Zr 2 [Al(Si)]4 C 5 ceramic and 30 vol.% SiC/Zr 2 [Al(Si)]4 C 5 composite [6]. The 30 vol.% (ZrB 2 + SiC)/Zr 2 [Al(Si)] 4 C 5 composite also shows a higher Young's modulus (425 GPa) compared with monolithic Zr2 [Al(Si)] 4 C5 ceramic(361 GPa). On the whole, the 30 vol.% (ZrB2 + SiC)/Zr2 [Al(Si)] 4 C5 composite shows much higher room temperature mechanical properties than monolithic Zr2[Al(Si)]4C5 ceramic as well as SiC individually reinforced Zr 2 [Al(Si)] 4 C 5 composites. The above results suggest that the in situ incorporation of ZrB 2 and SiC two reinforcements produces multivariate synergic strengthening and toughening effect, which is much more significant than that derived by SiC individual reinforcement. It is well known that both ZrB2 and SiC have a higher hardness (20–30 GPa, 23 GPa) and stiffness (Young's modulus of 403 GPa and 489 GPa) compared with Zr2[Al(Si)]4C5. Therefore, the incorporation of them could effectively improve the Vickers hardness and stiffness of the composites. Furthermore, the specific stiffness of the composites is greatly enhanced by adding the two reinforcements with lower density. At the same time, the higher hardness and Young's modulus as well as mismatched CTE of ZrB 2 and SiC compared with Zr2[Al(Si)]4C5 matrix will lead to higher deformation resistance and residual stress toughening. When a composite is cooled down from the fabrication temperature, thermal residual stress arises due to the thermo-elastic mismatch between constituents [18]. The CTEs of both ZrB2 and SiC are lower than that of Zr2 [Al(Si)]4C5 (8.1 × 10− 6/K for Zr2[Al(Si)]4C5 [3,6], 6.9 × 10− 6/K for ZrB2 [19] and 5.1 × 10− 6/K for SiC [20]), therefore the residual stresses at the Zr2[Al(Si)]4C5/ZrB2 and Zr2[Al(Si)]4C5/SiC interfaces are in a compressive stress [21], which can make the cracks be deflected and propagate along the grain boundary as they meet ZrB 2 or SiC particles. Therefore, the crack path becomes zigzag and prolonged, consuming a lot of fracture energy, which benefits for the improvement in strength and toughness of the composites. Fig. 7 shows the magnified image for the Vickers' indentation-induced crack propagation path of the 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite.

104


Zr2[Al(Si)]4C5

Fig. 3. TEM and HRTEM images of the 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite.

It can be seen that the cracks show dramatic deflection, branching and bridging during propagation. In addition, the incorporation of ZrB2 and SiC reinforcing agents significantly inhibits the grain-boundary migration and grain growth of Zr2 [Al(Si)]4C5, resulting in a fine-grained microstructure. Hall–Petch equation [22] describes the relation between the yield stress point and the grain size.

δy ¼ δ0 þ kd

‐1=2

ð5Þ

where δy is the yield stress for a polycrystal, δ0 is the yield stress for a single crystal or a polycrystal with an infinitely large grain size, d is the average grain size, and k is the Hall–Petch parameter. Therefore, the incorporation of (ZrB2 + SiC) and the increase of their content also induce a grain-refinement strengthening effect. Especially, both the nano SiC reinforcements and “self-reinforced” microstructure of ZrB2 will make an important contribution to the strength and toughness of the composites. In short, it is the synergistic action of many strengthening and toughening mechanisms such as particulate reinforcement, crack deflection, branching and bridging, “self-reinforced” microstructure as well as grain-refinement strengthening, which results in the excellent room temperature mechanical properties of (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites.

Undoubtedly, high-temperature mechanical properties make more significance for UHTCs. Fig. 8 shows the temperature dependence of Young's modulus of the 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite and Zr2[Al(Si)]4C5 [3,6]. A slight linear decrease of modulus is observed for both the composite and monolithic Zr2[Al(Si)]4C5 with increasing temperature, and no obvious sharp decrease appeared for the composite even at 1511 °C. The composite shows higher modulus compared to monolithic Zr2[Al(Si)]4C5 and SiC individually reinforced Zr2[Al(Si)]4C5 composite [6] from room temperature to 1511 °C. In consideration of the clean grain boundary with no glass phase, the grain boundary sliding in the composite at high temperature will be significantly suppressed. Therefore, it is reasonable to expect a low viscoelastic response and high macroscopic deformation resistance at elevated temperatures for the composite. Because of the limitation of testing instrument, the test result above 1511 °C was not obtained. However, we suppose that the Young's modulus of the 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite can be retained up to a higher temperature than 1511 °C. The remaining Young's modulus of the composite at 1511 °C is 361 GPa (86% of that at ambient temperature), which is even higher than that of Zr2 [Al(Si)]4C5 (306 GPa) [3,6] and most refractory compounds [23] at room temperature. The flexural strength of 30 vol.% (ZrB2 + SiC)/Zr2 [Al(Si)]4C5 composite and monolithic Zr2[Al(Si)]4C5 at 1300 °C in air was also tested, which can better reflect their high temperature mechanical properties under practical serving conditions. It can be seen that at 1300 °C in air the composite still shows a strength as high as (404 ± 13 MPa), which is about 65% of that at room temperature, and is about 50% higher than that of monolithic Zr2[Al(Si)]4C5 ceramic (261 ± 26 MPa) at 1300 °C in air. It is thus clear that, just as being expected, the in situ incorporation of (ZrB2 + SiC) improves significantly not only the room temperature mechanical properties, more importantly, but also the high temperature mechanical properties of the materials, which renders the (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites as a potential UHTC. Thermal physical properties

Fig. 4. Backscattered electron image of the polished surface of the 30 vol.%(ZrB 2 + SiC)/Zr 2 [Al(Si)] 4C 5 composite.

A least-square fit of thermal expansion curve of the 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite in the temperature range of 100–1200 °C yields the CTE of 7.99 × 10−6 K−1, which is lower than that of Zr2 [Al(Si)]4C5 ceramic(8.1 × 10− 6 K− 1) [3,6]. Since both ZrB2 and SiC have lower CTE compared with Zr2[Al(Si)]4C5, the incorporation of them will decrease the CTE of the composites, which is favorable to the improvement in thermal shock resistance of the materials. Fig. 9


105

Fig. 5. SEM micrographs of the fracture surfaces of the samples with different (ZrB2 + SiC) contents.

shows the temperature dependence of specific heat capacity of 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite. For comparison, that of Zr2 [Al(Si)]4C5 is also given. Curve fitting of the data for the composite yields a Cp∝T expression −3

5

T−73:94 10 T

−2

ð6Þ

with a R2 of 0.99. It can be seen that the composite shows the similar variation tendency and close value to Zr2[Al(Si)]4C5 in the whole temperature range of 25–1200 °C. The specific heat capacities of the composite at 25 °C and 1200 °C are 544 and 1045 J kg−1 K−1, respectively. The similar specific heat capacity value can be attributed to the higher one for SiC (620–770 J kg−1 K−1 [20]) and lower one for ZrB2 (427 J kg−1 K−1 [19]) compared with Zr2[Al(Si)]4C5. Thermal conductivity, especially at high temperatures, is another important parameter determining the thermal shock performance of

600

Bending strength Fracture toughness

8

500 6

400

300

0%

10%

(ZrB2+SiC)

20%

content /vol%

30%

1/2

4

Fracture toughness ( MPa /m )

Bending strength (MPa)

ð7Þ

with a R2 of 0.974. The thermal conductivities of monolithic Zr2[Al(Si)]4C5

10

700

200

k ¼ 14:38 þ 2802:75=T

Vickers Hardness (GPa)

Cp ¼ 619:44 þ 161:08 10

ceramics under high temperature service conditions. High thermal conductivity at high temperatures is favorable to reduce the temperature gradient born by the materials and consequently improve the thermal shock resistance. Therefore, investigation on high-temperature thermal conductivity for UHTCs has recently attracted much attention of researchers. Fig. 9 also shows the temperature dependence of thermal conductivity of 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite and Zr2 [Al(Si)]4C5 [6]from room temperature up to 1200 °C. It can be seen that both the two materials show a slow decrease in thermal conductivity with temperature, which nearly keeps constant at temperatures above 1000 °C. A least-square fit of the data of the composite yields the following relationship:

16.5

15.0

13.5

12.0

0

10 (ZrB2+SiC)

20

content /vol%

Fig. 6. Room temperature mechanical properties of the samples as a function of (ZrB2 + SiC) content.

30

106


Table 1 Comparison of some properties of 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5, Zr2[Al(Si)]4C5, and 30 vol.% SiC/Zr2[Al(Si)]4C5 [6]. Properties

30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5

Zr2[Al(Si)]4C5

30 vol.%SiC/Zr2[Al(Si)]4C5

Measured density (g·cm−3) Young's modulus (GPa) Shear modulus (GPa) Bulk modulus (GPa) Poisson's ratio Vickers hardness (GPa) Bending strength (MPa) Bending strength at 1300 °C in air (MPa) Fracture toughness (MPa·m1/2) Coefficient of thermal expansion (×10−6 K−1) Specific heat capacity (J·kg−1·K−1) Thermal conductivity at room temperature (W·m−1·K−1) Thermal shock resistance parameter R (K) RIV (μm) Rst (μm1/2·K)

4.76 422 177 228 0.173 16.7 621 ± 33 404 ± 13 7.85 ± 0.29 7.99 544.3 22.9 151 192.4 1638.1

4.45 361 153 188 0.18 11.7 353 ± 10 261 ± 26 3.88 ± 0.16 8.1 567 12 99 147.3 938.3

4.01 386 163 204 0.18 16.4 353 ± 19 – 6.62 ± 0.19 7.2 584 25.3 – – –

at 25 °C and 1200 °C are 12.0 and 10 Wm−1 K−1, respectively. As both ZrB2 and SiC have much higher thermal conductivity than Zr 2 [Al(Si)]4C5, the composite shows higher thermal conductivity than Zr 2[Al(Si)]4C 5 over the entire temperature range. The thermal conductivities of the composite at 25 °C and 1200 °C are 22 and 17.4 Wm− 1 K− 1, respectively. Although the test was only performed under 1200 °C because of the restrainment of instrument, it is conceivable to expect a higher thermal conductivity for the composite than Zr2 [Al(Si)]4 C5 at higher temperatures. Therefore, compared with Zr2[Al(Si)]4C5, the thermal conductivity of the composite was significantly improved at both room and high temperature.

Thermal shock resistance Severe thermal shock is a critical test confronting UHTCs under practical serving conditions, which is usually one of the main causes of their failure. Therefore, thermal shock resistance is vitally important for UHTCs. The curves of residual strength as well as residual strength ratio versus temperature difference for the monolithic Zr2[Al(Si)]4C5 ceramic and 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite are shown in Fig. 10. It can be seen that both the residual strength and residual strength ratio of the two materials decrease gradually with increasing temperature differences. The composite shows a higher residual strength over the entire test temperatures and a higher residual strength ratio at temperature difference up to 550 °C compared with monolithic Zr2[Al(Si)]4C5. The critical temperature difference ΔTc is an important index to evaluate the thermal shock resistance of materials. According to ASTM standard C1525-04 [24], the ΔTc is determined as

70% of the room temperature strength. The residual strength data indicates that the ΔTc of about 336 °C is obtained for the composite, which is 103 °C higher than that for Zr3[Al(Si)]4C6(233 °C). These reveal that the incorporation of ZrB2 and SiC significantly improves the thermal shock resistance of the materials. Three thermal shock resistance parameters, R, RIV and Rst, are proposed and used to quantitatively predict and evaluate the thermal shock performance of ceramics from different angles [25–28]. They also disclose the effects of intrinsic mechanical and thermal physical properties on thermal shock resistance of ceramics. They are defined as follows respectively.

R¼

IV

R

σ ð1−νÞ Εα

ð8Þ

2

¼

K IC σ ð1−νÞ 2

Rst ¼

γ α2 E

ð9Þ

1 2

ð10Þ

where σ is the tensile strength, E is Young's modulus, v is Poisson's ratio, α is thermal expansion coefficient, KIC is fracture toughness and γ is the fracture surface energy(which can be converted according to the Irwin equation: K2IC = 2γE). R is the thermal stress fracture resistance parameter, which predicts the maximum change in temperature (the critical

Fig. 7. Typical micrograph of crack propagation on the polished surface of the 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite.


Young's Modulus (GPa)

450

107

significantly increases the strength, fracture toughness and thermal conductivity and simultaneously decreases the thermal expansion coefficient of the materials.

400

Conclusions

350

300

□

30 vol.% (ZrB2+SiC)/Zr2[Al(Si)]4C5 Zr2[Al(Si)]4C5

250 0

200

400

600

800

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thermal shock temperature difference) that can occur without the initiation of cracks. The calculation assumes an instantaneous change in surface temperature (infinite heat transfer). RIV is the thermal stress damage resistance parameter, which reflects the resistance to catastrophic crack propagation of ceramics under a critical temperature difference ΔTc. Rst is the thermal stress crack stability parameter, which indicates the resistance to crack repropagation after a critical temperature difference, ΔTc. The calculated thermal shock resistance parameters of the two materials are also compared in Table 1. It can be seen that the three parameters of 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite are 151 K, 192.4 μm, and 1638.1 μm1/2·K, respectively, which are much higher than those of monolithic Zr2[Al(Si)]4C5. The calculation results are in good agreement with the thermal shock experiment results, which indicate an improved resistance against crack initiation, propagation and repropagation for the composites, and also further confirm the beneficial effect of the in situ incorporation of ZrB2 and SiC on the thermal shock resistance of the materials. In the present work, for both the composite and monolithic Zr2[Al(Si)]4C5, actual ΔTc is significantly larger than R value because of the finite heat transfer coefficient, h, between the specimen and the quenching medium, which reduces the thermal stresses during the thermal shock tests [24]. According to the above results and analysis, the composite has the similar Poisson's ratio to and little higher Young's modulus than Zr2[Al(Si)]4C5. Therefore, the significantly improved thermal shock resistance of the composite should be attributed to the fact that the in situ incorporation of (ZrB2 + SiC)

Dense (ZrB 2 + SiC)/Zr 2 [Al(Si)] 4 C 5 composites, with different volume fraction of (ZrB 2 + SiC) reinforcements were successfully prepared by in situ hot pressing method. ZrB2 and SiC incorporated by in situ reaction significantly improve the mechanical properties of the Z2[Al(Si)]4C5 matrix at both room and high temperatures by the synergistic action of many strengthening and toughening mechanisms such as particulate reinforcement, crack deflection, branching and bridging, “self-reinforced” microstructure as well as grain-refinement strengthening. With the (ZrB 2 + SiC) content increasing from 0 to 30 vol.%, the flexural strength, fracture toughness and Vickers hardness of the materials show a nearly linear increase from 353 to 621 MPa, 3.88 to 7.85 MPa·m 1/2 , and 11.7 to 16.7 GPa, respectively. Especially, owing to clean grain boundaries with no glass phase, the 30 vol.% (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composite shows much higher modulus than Zr2[Al(Si)]4C5 from room temperature up to 1511 °C, and could retain high modulus of 357 GPa at 1511 °C (86% of that at ambient temperature). The composite also could retain a superior flexural strength up to 404 ± 13 MPa at 1300 °C in air. In addition, the composite exhibits a higher thermal conductivity from 25 °C (24.6 Wm − 1 K − 1 ) to 1200 °C (16.6 Wm − 1 K − 1 ) compared with Z 2 [Al(Si)] 4 C 5 . As a result, excellent thermal shock resistance was obtained for the composite, which shows a higher residual strength over the entire test temperatures and a higher residual strength ratio at temperature difference up to 550 °C compared with monolithic Zr 2 [Al(Si)] 4 C 5 . Superior comprehensive properties suggest a promising prospect of (ZrB2 + SiC)/Zr2[Al(Si)]4C5 composites as UHTCs. Acknowledgments This work was supported by the National Natural Science Foundation of China (No. 50872052), the Natural Science Foundation of Jiangsu Province of China (No. BK2011806), the Opening Project of State Key Laboratory of High Performance Ceramics and Superfine Microstructure (No. SKL201108SIC), the Priority Academic Program Development of Jiangsu Higher Education Institutions (PAPD), and the Program for Changjiang Scholars and Innovative Research Team in University (PCSIRT), IRT1146. References [1] Fukuda K, Hisamura M, Iwata T, Tera N, Sato K. Synthesis, crystal structure, and thermoelectric properties of a new carbide Zr2[Al3.56Si0.44]C5. J Solid State Chem 2007;180:1809–15. [2] Lin ZJ, He LF, Wang JY, Li MS, Bao YW, Zhou YC. Atomic-scale microstructures and elastic properties of quaternary Zr–Al–Si–C ceramics. Acta Mater 2008;56:2022–31.

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