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Mechanical properties and thermal shock behavior of bionic laminated ZrB2–SiC–G ceramics
Int. Journal of Refractory Metals and Hard Materials 51 (2015) 233–238
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Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
Mechanical properties and thermal shock behavior of bionic laminated ZrB2–SiC–G ceramics Chuncheng Wei ⁎, Changshou Ye School of Material Science and Engineering, Shandong University of Technology, Zibo 255049, PR China
a r t i c l e
i n f o
Article history: Received 20 December 2014 Received in revised form 9 April 2015 Accepted 20 April 2015 Available online 22 April 2015 Keywords: Ceramics Mechanical properties Microstructure
a b s t r a c t Laminated ZrB2–SiC–G ceramics were prepared by the combination of tape casting and hot pressing. The various effects of the addition of ZrB2 and SiC particles to graphite layers on the microstructure, mechanical properties and fracture behavior, as well as thermal shock resistance were investigated in detail. When 30 vol.% ZrB2 and 10 vol.% SiC were added to the graphite layers, the laminated ZrB2–SiC–G ceramics were obtained with bending strength of 610 MPa, fracture toughness of 13.7 MPa·m1/2, fracture work of 617 J/m2, critical thermal shock temperature (ΔTC) of 508 °C. These results were attributed to the appropriate residual thermal stress between interface layers, which were important to the improvement of the mechanical properties and thermal shock resistance of the laminated ZrB2–SiC–G ceramics. © 2015 Published by Elsevier Ltd.
1. Introduction Zirconium diboride (ZrB2) and hafnium diboride (HfB2) are members of a family of materials known as ultrahigh temperature ceramics [1,2]. Several carbides and nitrides of the group IVB and VB transition metals are also considered ultrahigh temperature ceramics based on melting temperatures in excess of 3000 °C and other properties [3,4]. Very few elements or compounds from any class of ceramic materials have melting temperatures approaching 3000 °C [5–8]. Within the family of ultrahigh temperature ceramics, ZrB2 has the lowest theoretical density (6.09 g/cm3) [8,9], which make it an important member of ultrahigh temperature ceramics family, with a unique combination of properties for applications at extreme temperatures in very hostile environments [10,11], as for example those encountered in the field of structural applications, such as cutting tools, wear-resistant and vacuum metallization, high-temperature structural materials, neutron absorbers and lightweight impact resistant armor material, due to its high temperature properties like high melting point, low coefficient of thermal expansion, high thermal conductivity and excellent mechanical properties, such as high hardness, high Young's modulus, high strength to weight ratio, and chemical stability [12–14]. Although the ZrB2-based ceramics have many advantages, intrinsic characteristics such as low fracture toughness (premature failure due to brittle fracture) and poor thermal shock resistance are still obstacles for them to be used widely, especially for applications in extreme environment [15–17]. Fortunately, ceramics with bionic laminated and fibrous monolithic structures have been found to be excellent fracture toughness and thermal shock resistance [18]. Laminated ZrB2-based ceramics were ⁎ Corresponding author. E-mail address: [email protected] (C. Wei).
http://dx.doi.org/10.1016/j.ijrmhm.2015.04.023 0263-4368/© 2015 Published by Elsevier Ltd.
prepared in our previous work, and fracture toughness of 14 MPa·m1/2 was at least two times stronger than monolithic ZrB2-based ceramics [15,19]. However, intrinsic relationships between mechanical properties and composition/microstructure of laminated ZrB2-based ceramics have not been well understood, and very little attention has been paid to the thermal shock resistance of laminated ZrB2-based UHTCs. In this work, ZrB2–SiC and graphite green sheets were formed by tape casting, respectively; the ZrB2–SiC green sheets and graphite green sheets were stacked layer by layer and then hot-pressed. Finally, the laminated ZrB2–SiC–G ceramics were obtained successfully. The mechanical properties, fracture behavior and thermal shock behavior of the laminated ZrB2–SiC–G ceramics were investigated in detail. 2. Experimental procedures Commercially available ZrB2 (2 μm, N99.5% purity, Northwest Institute for Non-ferrous Metal Research, China), SiC (0.5 μm, N99.5% purity, Weifang Kaihua Micro-powder Co., Ltd., China) and graphite flake (mean diameter and thickness are 15 and 1.5 μm, respectively, N 99% purity, Qingdao Tiansheng Graphite Co., Ltd., China) powders were used as raw materials. The starting mixtures of 80 vol.% ZrB2 plus 20 vol.% SiC were first ball milled for 10 h in a polyethylene bottle using ZrO2 balls and ethanol as the grinding media. Then adhesive and plasticizer were added to the slurry and a further ball-milling for 4 h to obtain a uniform slurry. Polyvinyl butyral resin and polyethylene glycol-6000 were chosen as the adhesive and plasticizer, respectively. The mass ratio of polyvinyl butyral resin, polyethylene glycol-6000 and ethanol to ZrB2–SiC mixtures in the slurry was 1:1:15:10. The ZrB2–SiC green sheets were formed by tape casting on home-made instruments. By the same route the graphite green sheets were formed. After drying, the ZrB2– SiC green sheets and graphite green sheets were stacked alternately
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Table 1 Composition of the as-prepared bionic laminated ZrB2–SiC–G ceramics. Specimens
ZrB2–SiC layers
Graphite layers
LZS-1 LZS-2 LZS-3
ZrB2–20 vol.% SiC ZrB2–20 vol.% SiC ZrB2–20 vol.% SiC
Graphite–20 vol.% ZrB2–10 vol.% SiC Graphite–30 vol.% ZrB2–10 vol.% SiC Graphite–40 vol.% ZrB2–10 vol.% SiC
until the desired compositions were achieved. The laminated material was placed into BN covered graphite molds, vacuum degreased at 650 °C for 30 min, and the laminated ZrB2–SiC–G UHTCs were heated at heat speed of 10 °C/min (Hot-pressing Furnace ZYD-60, Jinzhou Hangxing Vacuum Equipment Co., LTD, China), then the laminated ZrB2–SiC–G UHTCs were hot-pressed at 1900 °C under a uniaxial load of 30 MPa for 1 h, finally the laminated ZrB2–SiC–G UHTCs were cooled from 1900 °C to room temperature naturally. Table 1 illustrates the composition of the as-prepared ZrB2–SiC–G UHTCs samples in this work. Flexural strength (σ) was measured using three-point bending tests (Model 5569, Instron, USA) on 36 mm × 3 mm × 4 mm test bars, with a loading span of 30 mm and a crosshead speed of 0.5 mm/min at room temperature. Young's modulus is calculated by stress divided by strain. Fracture toughness (KIC) was evaluated using single-edge notched bend (SENB, Model 5569, Instron, USA) beams 22 mm × 4 mm × 2 mm, with a notched depth and width of 2 and 0.2 mm, respectively, a span of 16 mm and a crosshead speed of 0.05 mm/min. The loading direction and the notch plane lie perpendicular to the general plane of orientation of the graphite interface layers. Mechanical performance was evaluated by the average test result of 7–10 specimens. Indentation method was used to measure Vickers hardness with a 49 N load for 15 s (Wilson® VH1150, Buehler an ITW Co., Ltd, Germany). The thermal shock behavior was assessed by water quenching method. The polished test bars (36 mm × 3 mm × 4 mm) were heated in air up to the desired temperature and held for 10 min. The temperature of the water bath was controlled at about 20 °C. The thermal quenching of single thermal shock was conducted at the initial temperatures set as 220, 420, 620, 820, and 1020 °C, which correspond to the thermal shock temperature differences (ΔT) of 200, 400, 600, 800, 1000 °C, respectively. The retained flexural strength (σf) after water quenching was measured. Scanning electron microscope (SEM, FEI Sirion, Holland) and Energy dispersive spectrometer
a
(EDS, EDAX Inc., USA) were employed to analyze the microstructures and element composition of the tested specimens. 3. Results and discussion SEM images of the cross-section of the laminated ZrB2–SiC–G ceramics are shown in Fig. 1. It was found that the thicknesses for ZrB2– SiC layers and graphite layers were about 300 and 30 μm, respectively. It was observed that all of layers were relatively flat and arranged alternatively without faults, as shown in Fig. 1a, b and c. The Vickers hardness of ZrB2–SiC layers and graphite layers was about 16.5 and 4.8 GPa, respectively. Because the thickness of graphite layer was 30 μm and the diagonals of Vickers indentations on graphite layer using a load of 49 N are extremely more than 30 μm, Vickers hardness of apparent graphite layer (indentation tip has been partially extended to the ceramic layer) was about 4.8 GPa. In the present work, Vickers hardness tester was used to measure the hardness of ZrB2–SiC layers and graphite layers for comparison under the same test conditions. The hardness of the graphite layers was relatively lower than ZrB2–SiC layers, which could be related to the falling-off phenomenon of low hardness graphite particles during polishing [15], as shown in Fig. 1d. As shown in Fig. 1e and f, the graphite layers were distinctly dense and the height difference between ZrB2–SiC and graphite layers decreased with the addition of ZrB2 particles into graphite layers. In addition, graphite flakes were found to be perpendicular to the hot-pressing direction, which was beneficial to the improvement of fracture toughness owing to deflection cracks [20,21]. Fig. 2 shows the fracture behavior of the laminated ZrB2–SiC–G ceramics. It was observed that the fracture surface in ZrB2–SiC layers was relatively flat (Fig. 2a and b), whereas the fracture surface in graphite layers had an obvious deflection behavior. Graphite flakes were found to be parallel to the general plane of orientation of the graphite layers, i.e. perpendicular to the hot-pressing direction during fabrication, and the results demonstrated that hot-pressed graphite flakes presented preferred orientation, which was in good accordance with the results reported by Wang [22,23]. In the present study, preferred orientation was presumed to generate crack deflection in graphite layers under stress, which resulted in the improvement of the fracture toughness of as-prepared bionic laminated ZrB2–SiC–G ceramics. In Fig. 2c, crack deflection behavior was weakened significantly and the
b
500 µm
d
c
500 µm
e
20 µm
500 µm
f
20 µm
20 µm
Fig. 1. SEM images of the cross-section of the as-prepared bionic laminated ZrB2–SiC–G ceramics: (a and d) LZS-1, (b and e) LZS-2, (c and f) LZS-3.
C. Wei, C. Ye / Int. Journal of Refractory Metals and Hard Materials 51 (2015) 233–238
a
b
c
500 µm
d
235
500 µm
e
500 µm
f
2 µm
2 µm
2 µm
Fig. 2. SEM images of fracture surface of laminated ZrB2–SiC–G ceramics in the three-point bending test: (a and d) LZS-1, (b and e) LZS-2, (c and f) LZS-3.
step-like fracture surface disappeared with more heat-resistant particles added into graphite layers. The reason was attributed to the residual thermal stresses caused by the thermal expansion coefficient discrepancy between ZrB2–SiC layers and graphite layers. The thermal expansion coefficients of ZrB2–SiC and pure graphite were 7.18 × 10−6/K and 30 × 10−6/K, respectively [24]. Great residual thermal stress among layers was believed to result in the poor interfacial adherence and even delamination [20,21], while little residual thermal stress could weaken crack deflection resulting in the reduction of fracture toughness. The key factor for the improvement of mechanical properties and thermal shock behavior of laminated ZrB2–SiC–G ceramics is the optimal residual thermal stresses between the ZrB2–SiC layers and graphite layers, which could be calculated by Chartier formula [25], In Eqs. (1) and (2), σres, E, ν, α, and ΔT refer to residual thermal stresses, modulus of elasticity, Poisson's ratio, coefficients of thermal expansion, and difference between sintering and initial temperatures, respectively. In the present study, ZrB2–SiC and graphite layers were indexed by subscripts 1 and 2, respectively. σ res1 ¼ −
σ res2 ¼
E1 E2 h2 ðα 2 −α 1 ÞΔT ð1−ν 1 ÞE2 h2 þ ð1−ν2 ÞE1 h1
ð1Þ
E1 E2 h1 ðα 2 −α 1 ÞΔT ð1−ν 1 ÞE2 h2 þ ð1−ν 2 ÞE1 h1
ð2Þ
a
Residual thermal stresses in laminated ZrB2–SiC–G ceramics were presented by adding σres1 and σres2, where E1 and E2 were 450 and 15 GPa, respectively. ν1 and ν2 were 0.13 and 0.25, respectively. σres2 was calculated as tensile stress of 852 MPa, while σres1 was calculated as compressive stress of − 85 MPa. Compressive stress was supposed to inhibit the propagation of the cracks [20], which lead to a higher bending strength of laminated ZrB2–SiC–G ceramics than monolithic ZrB2–SiC ceramics [25,26]. With the addition of heat-resistant particles (ZrB2 and SiC) into graphite layers, the crack deflection was weakened because of the decrease in residual thermal stresses which was caused by the decrease in thermal expansion of graphite layers and the accordingly decrease in expansion differences between ZrB2–SiC layers and graphite layers. The optimal residual thermal stresses between the ZrB2–SiC layers and graphite layers can be obtained by the addition of a certain amount of heat-resistant particles into graphite layers. Besides, it should be noted that the addition of heat-resistant particles into graphite layers could improve high temperature oxidation resistance of laminated ZrB2–SiC–G ceramics. Fig. 2d shows the fracture surface of the ZrB2–SiC layers for LZS-1 specimens. The fracture mode of the ZrB2–SiC layers for LZS-1 specimens exhibited transgranular fracture, which meant consuming more fracture energy. The fracture mode of the ZrB2–SiC layers for LZS-2 specimens exhibited transgranular fracture and intergranular fracture (Fig. 2e), while that for LZS-3 specimens exhibited intergranular fracture and many grain pulled out (Fig. 2f). It
b
500 µm
c
500 µm
Fig. 3. SEM images of laminated ZrB2–SiC–G ceramics in the SENB test: (a) LZS-1, (b) LZS-2, and (c) LZS-3.
500 µm
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was revealed that the residual thermal stress between interface layers can influence fracture behavior of laminated ZrB2–SiC–G ceramics and fracture morphology of the ZrB2–SiC layers. To compare fracture behavior of the laminated ZrB2–SiC–G ceramics, a series of mechanical tests were carried out. Fig. 3 shows the SEM images of the crack propagation of laminated ZrB2–SiC–G ceramics after the SENB test. It can be seen from Fig. 3a and b, major crack deflection occurred repeatedly along graphite layers, the crack propagation path was significantly prolonged and the step-like fracture surface was clearly observed. Furthermore, the branching of major cracks was also found in Fig. 3a. The deflection of the cracks was obviously weakened and the crack propagation path was shorted with the addition of heat-resistance particles into graphite layers (Fig. 3c). The measured load–displacement curves of laminated ZrB2–SiC–G ceramics were shown in Fig. 4. The laminated ZrB2–SiC–G ceramics with different fracture curves were analyzed and compared to monolithic ZrB2–SiC ceramics. Fracture curves for LZS-1 and LZS-2 exhibited a two-time fracture, i.e. the load further increased after first-time partly fracture, and while that for LZS-3 exhibited catastrophic fracture at the maximum load. The two-time fracture mode can effectively consume fracture energy, and thus enhance reliability of the laminated ceramics as the structural components. The rupture work could be calculated according to Eq. (3) [27].
γ wo f ¼
1 2A
Z pdδ
ð3Þ
where γwof refers to the rupture work; A refers to the projected area of fracture surface; ∫ pdδ refers to the total work of applied load, which can be calculated by the area of load–displacement curves. With standard samples employed in the test, larger projected area indicated more rupture work of the specimens. As shown in Fig. 4, rupture work was found to decrease with an increase in the addition of heatresistant particles into graphite layers. The measured density and mechanical properties of the laminated ZrB2–SiC–G ceramics were listed in Table 2. It was found that the density (ρ), flexural strength (σ), fracture toughness (KIC) and Young's modulus (E) increased with the increase in the heat-resistant particles adding into graphite layers. Compared to monolithic ZrB2–SiC ceramics, the fracture toughness of LZS-3 specimens raised to 13.7 MPa·m1/2 by 185% as a result of an effective extension of crack propagation caused by crack deflection and branching. In addition, the bending strength of LZS-3 specimens was tested to be 670 MPa and the increase of 75% compared to monolithic ZrB2–SiC ceramics [7,9], which was attributed to the inhibition of crack propagation by compressive stress of ZrB2–SiC layers. While, the rupture work of laminated ZrB2–SiC–G ceramics decreased
Fig. 4. Load-deflection curves of laminated ZrB2–SiC–G specimens.
Table 2 Mechanical properties of the laminated ZrB2–SiC–G ceramics. Materials
LZS-1 LZS-2 LZS-3
ρ
σf
KIC
E
rwof
(g/cm3)
(MPa)
(MPa·m1/2)
(GPa)
(J/m2)
5.18 5.24 5.31
587 ± 23 610 ± 18 670 ± 27
12.1 ± 0.1 13.2 ± 0.3 13.7 ± 0.4
259 ± 19 277 ± 24 293 ± 27
737 ± 28 617 ± 31 480 ± 35
with the addition of heat-resistant particles into graphite layers. Compared to laminated LZS-1 specimens, the rupture work of LZS-3 specimens was reduced to 480 J/m2 by 35%. The curves of retained flexural strength of laminated ZrB2–SiC–G ceramics versus temperature difference (ΔT, exposure temperature minus the water quench temperature) were plotted in Fig. 5. By contrast with room temperature strength of laminated ZrB2–SiC–G ceramics, i.e. the bending strength before thermal shock, the curves of flexural strength for LZS-1 and LZS-3 showed a similar sharp drop as ΔT increased to 200 °C, followed by a modest change as ΔT increased to 400 °C, while that for LZS-2 decreased gradually and slowly. The minimum bending strength was obtained in laminated ZrB2–SiC–G specimens after thermal shock at ΔT of 800 °C and a slight increase was found as ΔT increased to 1000 °C. The critical value of the temperature difference (ΔTC), an important index to evaluate the thermal shock resistance, was determined by a 30% reduction in flexural strength compared to the average bending strength of the tested specimens [27–29]. ΔTC for LZS-1, LZS-2 and LZS-3 was calculated as 470 °C, 508 °C and 400 °C, respectively. The highest ΔTC for LZS-2 specimens, raised by 8% for LZS-1, 27% for LZS-3 and 32% for monolithic ZrB2–SiC [28–30], illustrated that it was the optimized residual thermal stresses between ZrB2–SiC layers and graphite layers that play a key role in the improvement of laminated ZrB2–SiC–G ceramics thermal shock behavior, which revealed that too large or too small residual thermal stress among layers was not conducive to the improvement of thermal shock resistance. To further analyze the reason for low residual strength after thermal shock, the microstructures of LZS-2 specimens under different temperatures of thermal shock were observed. SEM images of LZS-2 specimens after thermal shock at different temperatures are shown in Fig. 6. As indicated, smooth surface with no cracks was shown for specimens quenched at 420 °C (Fig. 6a) and 620 °C (Fig. 6b), indicating superior thermal shock resistance of materials, while obvious cracks with a certain depth were found for specimens quenched at 820 °C (Fig. 6c), indicating that the destruction of materials massive flakes observed at 1020 °C in Fig. 6d was determined to be ZrO2 by EDS analysis. This was also supposed
Fig. 5. Residual strength versus temperature difference for LZS-1, LZS-2, and LZS-3 specimens.
C. Wei, C. Ye / Int. Journal of Refractory Metals and Hard Materials 51 (2015) 233–238
237
b
a
5 µm
c
5 µm
d
5 µm
5 µm
Fig. 6. SEM images of LZS-2 specimens after thermal shock at different temperatures: (a) 420 °C, (b) 620 °C, (c) 820 °C, and (d) 1020 °C.
to be the oxidation product of ZrB2 particles and it's the oxidation product that benefited cracks healing and accordingly accounted for the slight increase of residual strength as ΔT increased from 820 °C to 1020 °C. In order to further evaluate the thermal shock properties of laminated ZrB2–SiC–G ceramics, two thermal shock resistance parameters were used to predict the crack initiation and propagation behavior induced by the thermal shock stress of ceramic matrix composites. One was the thermal stress fracture resistance parameter R, expressed in Eq. (4) in terms of the bending strength of the material (σf), Young's modulus (E), Poisson's ratio (ν) and the coefficient of thermal expansion (α). The thermal shock parameter R predicts the maximum change in temperature (the critical thermal shock temperature difference) that can occur without the initiation of cracks. The other parameter was the thermal stress damage resistance parameter RIV, as expressed in Eq. (5), where σf is the bending strength of the material, KIC is the fracture toughness and ν is the Poisson's ratio, respectively. RIV decides the resistance to catastrophic crack propagation of ceramics under a critical temperature difference. R¼
R
IV
σ f ð1−ν Þ Eα
¼
ð4Þ
2 1 KIC 1−ν σ f
ð5Þ
For the simplicity of the calculation, the magnitudes of the Poisson's ratio (0.13) and the thermal expansion coefficient (7.18 × 10−6/K) for laminated ZrB2–SiC-G ceramics were supposed to be identical. Thermal shock resistance parameters of tested specimens were shown in Table 3. As listed, LZS-2 presented the highest RIV value, 110% of LZS-1, 121% of
LZS-3 and 560% of monolithic ZrB2–SiC ceramics (96 μm) [14], and similar R values were found for LZS-1, LZS-2 and LZS-3, ~182% of R value (147 °C) for monolithic ZrB2–SiC ceramics [14,30], indicating superior thermal shock behavior of laminated ZrB2–SiC–G as demonstrated in Fig. 5. 4. Conclusions The laminated ZrB2–SiC–G ceramics with improved mechanical properties and thermal shock resistance were obtained and investigated in detail. Flexural strength and fracture toughness of the laminated ZrB2–SiC–G ceramics were increased with the addition of heatresistant particles into graphite layers, while the rupture work was decreased gradually. It was found that further improvement of the mechanical properties and thermal shock resistance can be made through optimizing residual thermal stresses by changing the amount of heatresistant particles adding into graphite layers. High performance laminated ZrB2–SiC–G UHTCs were prepared with 30 vol.% ZrB2 and 10 vol.% SiC in graphite layers, with bending strength of 610 MPa, fracture toughness of 13.2 MPa·m1/2, rupture work of 617 J/m2 and critical thermal shock temperature (ΔTC) of 508 °C. Acknowledgment This work was supported by the National Science Foundation of China (50972029) and project (HIT.KLOF.2009026) supported by the Key laboratory Opening Funding of National Key Laboratory of Advanced Composites in Special Environments. References
Table 3 Thermal shock resistance parameters of the laminated ZrB2–SiC–G ceramics. Thermal shock resistance parameters
LZS-1
LZS-2
LZS-3
ΔTC (°C) R (°C) RIV (μm)
470 275 488
508 267 538
400 277 480
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Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
Mechanical properties and thermal shock behavior of bionic laminated ZrB2–SiC–G ceramics Chuncheng Wei ⁎, Changshou Ye School of Material Science and Engineering, Shandong University of Technology, Zibo 255049, PR China
a r t i c l e
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Article history: Received 20 December 2014 Received in revised form 9 April 2015 Accepted 20 April 2015 Available online 22 April 2015 Keywords: Ceramics Mechanical properties Microstructure
a b s t r a c t Laminated ZrB2–SiC–G ceramics were prepared by the combination of tape casting and hot pressing. The various effects of the addition of ZrB2 and SiC particles to graphite layers on the microstructure, mechanical properties and fracture behavior, as well as thermal shock resistance were investigated in detail. When 30 vol.% ZrB2 and 10 vol.% SiC were added to the graphite layers, the laminated ZrB2–SiC–G ceramics were obtained with bending strength of 610 MPa, fracture toughness of 13.7 MPa·m1/2, fracture work of 617 J/m2, critical thermal shock temperature (ΔTC) of 508 °C. These results were attributed to the appropriate residual thermal stress between interface layers, which were important to the improvement of the mechanical properties and thermal shock resistance of the laminated ZrB2–SiC–G ceramics. © 2015 Published by Elsevier Ltd.
1. Introduction Zirconium diboride (ZrB2) and hafnium diboride (HfB2) are members of a family of materials known as ultrahigh temperature ceramics [1,2]. Several carbides and nitrides of the group IVB and VB transition metals are also considered ultrahigh temperature ceramics based on melting temperatures in excess of 3000 °C and other properties [3,4]. Very few elements or compounds from any class of ceramic materials have melting temperatures approaching 3000 °C [5–8]. Within the family of ultrahigh temperature ceramics, ZrB2 has the lowest theoretical density (6.09 g/cm3) [8,9], which make it an important member of ultrahigh temperature ceramics family, with a unique combination of properties for applications at extreme temperatures in very hostile environments [10,11], as for example those encountered in the field of structural applications, such as cutting tools, wear-resistant and vacuum metallization, high-temperature structural materials, neutron absorbers and lightweight impact resistant armor material, due to its high temperature properties like high melting point, low coefficient of thermal expansion, high thermal conductivity and excellent mechanical properties, such as high hardness, high Young's modulus, high strength to weight ratio, and chemical stability [12–14]. Although the ZrB2-based ceramics have many advantages, intrinsic characteristics such as low fracture toughness (premature failure due to brittle fracture) and poor thermal shock resistance are still obstacles for them to be used widely, especially for applications in extreme environment [15–17]. Fortunately, ceramics with bionic laminated and fibrous monolithic structures have been found to be excellent fracture toughness and thermal shock resistance [18]. Laminated ZrB2-based ceramics were ⁎ Corresponding author. E-mail address: [email protected] (C. Wei).
http://dx.doi.org/10.1016/j.ijrmhm.2015.04.023 0263-4368/© 2015 Published by Elsevier Ltd.
prepared in our previous work, and fracture toughness of 14 MPa·m1/2 was at least two times stronger than monolithic ZrB2-based ceramics [15,19]. However, intrinsic relationships between mechanical properties and composition/microstructure of laminated ZrB2-based ceramics have not been well understood, and very little attention has been paid to the thermal shock resistance of laminated ZrB2-based UHTCs. In this work, ZrB2–SiC and graphite green sheets were formed by tape casting, respectively; the ZrB2–SiC green sheets and graphite green sheets were stacked layer by layer and then hot-pressed. Finally, the laminated ZrB2–SiC–G ceramics were obtained successfully. The mechanical properties, fracture behavior and thermal shock behavior of the laminated ZrB2–SiC–G ceramics were investigated in detail. 2. Experimental procedures Commercially available ZrB2 (2 μm, N99.5% purity, Northwest Institute for Non-ferrous Metal Research, China), SiC (0.5 μm, N99.5% purity, Weifang Kaihua Micro-powder Co., Ltd., China) and graphite flake (mean diameter and thickness are 15 and 1.5 μm, respectively, N 99% purity, Qingdao Tiansheng Graphite Co., Ltd., China) powders were used as raw materials. The starting mixtures of 80 vol.% ZrB2 plus 20 vol.% SiC were first ball milled for 10 h in a polyethylene bottle using ZrO2 balls and ethanol as the grinding media. Then adhesive and plasticizer were added to the slurry and a further ball-milling for 4 h to obtain a uniform slurry. Polyvinyl butyral resin and polyethylene glycol-6000 were chosen as the adhesive and plasticizer, respectively. The mass ratio of polyvinyl butyral resin, polyethylene glycol-6000 and ethanol to ZrB2–SiC mixtures in the slurry was 1:1:15:10. The ZrB2–SiC green sheets were formed by tape casting on home-made instruments. By the same route the graphite green sheets were formed. After drying, the ZrB2– SiC green sheets and graphite green sheets were stacked alternately
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Table 1 Composition of the as-prepared bionic laminated ZrB2–SiC–G ceramics. Specimens
ZrB2–SiC layers
Graphite layers
LZS-1 LZS-2 LZS-3
ZrB2–20 vol.% SiC ZrB2–20 vol.% SiC ZrB2–20 vol.% SiC
Graphite–20 vol.% ZrB2–10 vol.% SiC Graphite–30 vol.% ZrB2–10 vol.% SiC Graphite–40 vol.% ZrB2–10 vol.% SiC
until the desired compositions were achieved. The laminated material was placed into BN covered graphite molds, vacuum degreased at 650 °C for 30 min, and the laminated ZrB2–SiC–G UHTCs were heated at heat speed of 10 °C/min (Hot-pressing Furnace ZYD-60, Jinzhou Hangxing Vacuum Equipment Co., LTD, China), then the laminated ZrB2–SiC–G UHTCs were hot-pressed at 1900 °C under a uniaxial load of 30 MPa for 1 h, finally the laminated ZrB2–SiC–G UHTCs were cooled from 1900 °C to room temperature naturally. Table 1 illustrates the composition of the as-prepared ZrB2–SiC–G UHTCs samples in this work. Flexural strength (σ) was measured using three-point bending tests (Model 5569, Instron, USA) on 36 mm × 3 mm × 4 mm test bars, with a loading span of 30 mm and a crosshead speed of 0.5 mm/min at room temperature. Young's modulus is calculated by stress divided by strain. Fracture toughness (KIC) was evaluated using single-edge notched bend (SENB, Model 5569, Instron, USA) beams 22 mm × 4 mm × 2 mm, with a notched depth and width of 2 and 0.2 mm, respectively, a span of 16 mm and a crosshead speed of 0.05 mm/min. The loading direction and the notch plane lie perpendicular to the general plane of orientation of the graphite interface layers. Mechanical performance was evaluated by the average test result of 7–10 specimens. Indentation method was used to measure Vickers hardness with a 49 N load for 15 s (Wilson® VH1150, Buehler an ITW Co., Ltd, Germany). The thermal shock behavior was assessed by water quenching method. The polished test bars (36 mm × 3 mm × 4 mm) were heated in air up to the desired temperature and held for 10 min. The temperature of the water bath was controlled at about 20 °C. The thermal quenching of single thermal shock was conducted at the initial temperatures set as 220, 420, 620, 820, and 1020 °C, which correspond to the thermal shock temperature differences (ΔT) of 200, 400, 600, 800, 1000 °C, respectively. The retained flexural strength (σf) after water quenching was measured. Scanning electron microscope (SEM, FEI Sirion, Holland) and Energy dispersive spectrometer
a
(EDS, EDAX Inc., USA) were employed to analyze the microstructures and element composition of the tested specimens. 3. Results and discussion SEM images of the cross-section of the laminated ZrB2–SiC–G ceramics are shown in Fig. 1. It was found that the thicknesses for ZrB2– SiC layers and graphite layers were about 300 and 30 μm, respectively. It was observed that all of layers were relatively flat and arranged alternatively without faults, as shown in Fig. 1a, b and c. The Vickers hardness of ZrB2–SiC layers and graphite layers was about 16.5 and 4.8 GPa, respectively. Because the thickness of graphite layer was 30 μm and the diagonals of Vickers indentations on graphite layer using a load of 49 N are extremely more than 30 μm, Vickers hardness of apparent graphite layer (indentation tip has been partially extended to the ceramic layer) was about 4.8 GPa. In the present work, Vickers hardness tester was used to measure the hardness of ZrB2–SiC layers and graphite layers for comparison under the same test conditions. The hardness of the graphite layers was relatively lower than ZrB2–SiC layers, which could be related to the falling-off phenomenon of low hardness graphite particles during polishing [15], as shown in Fig. 1d. As shown in Fig. 1e and f, the graphite layers were distinctly dense and the height difference between ZrB2–SiC and graphite layers decreased with the addition of ZrB2 particles into graphite layers. In addition, graphite flakes were found to be perpendicular to the hot-pressing direction, which was beneficial to the improvement of fracture toughness owing to deflection cracks [20,21]. Fig. 2 shows the fracture behavior of the laminated ZrB2–SiC–G ceramics. It was observed that the fracture surface in ZrB2–SiC layers was relatively flat (Fig. 2a and b), whereas the fracture surface in graphite layers had an obvious deflection behavior. Graphite flakes were found to be parallel to the general plane of orientation of the graphite layers, i.e. perpendicular to the hot-pressing direction during fabrication, and the results demonstrated that hot-pressed graphite flakes presented preferred orientation, which was in good accordance with the results reported by Wang [22,23]. In the present study, preferred orientation was presumed to generate crack deflection in graphite layers under stress, which resulted in the improvement of the fracture toughness of as-prepared bionic laminated ZrB2–SiC–G ceramics. In Fig. 2c, crack deflection behavior was weakened significantly and the
b
500 µm
d
c
500 µm
e
20 µm
500 µm
f
20 µm
20 µm
Fig. 1. SEM images of the cross-section of the as-prepared bionic laminated ZrB2–SiC–G ceramics: (a and d) LZS-1, (b and e) LZS-2, (c and f) LZS-3.
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a
b
c
500 µm
d
235
500 µm
e
500 µm
f
2 µm
2 µm
2 µm
Fig. 2. SEM images of fracture surface of laminated ZrB2–SiC–G ceramics in the three-point bending test: (a and d) LZS-1, (b and e) LZS-2, (c and f) LZS-3.
step-like fracture surface disappeared with more heat-resistant particles added into graphite layers. The reason was attributed to the residual thermal stresses caused by the thermal expansion coefficient discrepancy between ZrB2–SiC layers and graphite layers. The thermal expansion coefficients of ZrB2–SiC and pure graphite were 7.18 × 10−6/K and 30 × 10−6/K, respectively [24]. Great residual thermal stress among layers was believed to result in the poor interfacial adherence and even delamination [20,21], while little residual thermal stress could weaken crack deflection resulting in the reduction of fracture toughness. The key factor for the improvement of mechanical properties and thermal shock behavior of laminated ZrB2–SiC–G ceramics is the optimal residual thermal stresses between the ZrB2–SiC layers and graphite layers, which could be calculated by Chartier formula [25], In Eqs. (1) and (2), σres, E, ν, α, and ΔT refer to residual thermal stresses, modulus of elasticity, Poisson's ratio, coefficients of thermal expansion, and difference between sintering and initial temperatures, respectively. In the present study, ZrB2–SiC and graphite layers were indexed by subscripts 1 and 2, respectively. σ res1 ¼ −
σ res2 ¼
E1 E2 h2 ðα 2 −α 1 ÞΔT ð1−ν 1 ÞE2 h2 þ ð1−ν2 ÞE1 h1
ð1Þ
E1 E2 h1 ðα 2 −α 1 ÞΔT ð1−ν 1 ÞE2 h2 þ ð1−ν 2 ÞE1 h1
ð2Þ
a
Residual thermal stresses in laminated ZrB2–SiC–G ceramics were presented by adding σres1 and σres2, where E1 and E2 were 450 and 15 GPa, respectively. ν1 and ν2 were 0.13 and 0.25, respectively. σres2 was calculated as tensile stress of 852 MPa, while σres1 was calculated as compressive stress of − 85 MPa. Compressive stress was supposed to inhibit the propagation of the cracks [20], which lead to a higher bending strength of laminated ZrB2–SiC–G ceramics than monolithic ZrB2–SiC ceramics [25,26]. With the addition of heat-resistant particles (ZrB2 and SiC) into graphite layers, the crack deflection was weakened because of the decrease in residual thermal stresses which was caused by the decrease in thermal expansion of graphite layers and the accordingly decrease in expansion differences between ZrB2–SiC layers and graphite layers. The optimal residual thermal stresses between the ZrB2–SiC layers and graphite layers can be obtained by the addition of a certain amount of heat-resistant particles into graphite layers. Besides, it should be noted that the addition of heat-resistant particles into graphite layers could improve high temperature oxidation resistance of laminated ZrB2–SiC–G ceramics. Fig. 2d shows the fracture surface of the ZrB2–SiC layers for LZS-1 specimens. The fracture mode of the ZrB2–SiC layers for LZS-1 specimens exhibited transgranular fracture, which meant consuming more fracture energy. The fracture mode of the ZrB2–SiC layers for LZS-2 specimens exhibited transgranular fracture and intergranular fracture (Fig. 2e), while that for LZS-3 specimens exhibited intergranular fracture and many grain pulled out (Fig. 2f). It
b
500 µm
c
500 µm
Fig. 3. SEM images of laminated ZrB2–SiC–G ceramics in the SENB test: (a) LZS-1, (b) LZS-2, and (c) LZS-3.
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was revealed that the residual thermal stress between interface layers can influence fracture behavior of laminated ZrB2–SiC–G ceramics and fracture morphology of the ZrB2–SiC layers. To compare fracture behavior of the laminated ZrB2–SiC–G ceramics, a series of mechanical tests were carried out. Fig. 3 shows the SEM images of the crack propagation of laminated ZrB2–SiC–G ceramics after the SENB test. It can be seen from Fig. 3a and b, major crack deflection occurred repeatedly along graphite layers, the crack propagation path was significantly prolonged and the step-like fracture surface was clearly observed. Furthermore, the branching of major cracks was also found in Fig. 3a. The deflection of the cracks was obviously weakened and the crack propagation path was shorted with the addition of heat-resistance particles into graphite layers (Fig. 3c). The measured load–displacement curves of laminated ZrB2–SiC–G ceramics were shown in Fig. 4. The laminated ZrB2–SiC–G ceramics with different fracture curves were analyzed and compared to monolithic ZrB2–SiC ceramics. Fracture curves for LZS-1 and LZS-2 exhibited a two-time fracture, i.e. the load further increased after first-time partly fracture, and while that for LZS-3 exhibited catastrophic fracture at the maximum load. The two-time fracture mode can effectively consume fracture energy, and thus enhance reliability of the laminated ceramics as the structural components. The rupture work could be calculated according to Eq. (3) [27].
γ wo f ¼
1 2A
Z pdδ
ð3Þ
where γwof refers to the rupture work; A refers to the projected area of fracture surface; ∫ pdδ refers to the total work of applied load, which can be calculated by the area of load–displacement curves. With standard samples employed in the test, larger projected area indicated more rupture work of the specimens. As shown in Fig. 4, rupture work was found to decrease with an increase in the addition of heatresistant particles into graphite layers. The measured density and mechanical properties of the laminated ZrB2–SiC–G ceramics were listed in Table 2. It was found that the density (ρ), flexural strength (σ), fracture toughness (KIC) and Young's modulus (E) increased with the increase in the heat-resistant particles adding into graphite layers. Compared to monolithic ZrB2–SiC ceramics, the fracture toughness of LZS-3 specimens raised to 13.7 MPa·m1/2 by 185% as a result of an effective extension of crack propagation caused by crack deflection and branching. In addition, the bending strength of LZS-3 specimens was tested to be 670 MPa and the increase of 75% compared to monolithic ZrB2–SiC ceramics [7,9], which was attributed to the inhibition of crack propagation by compressive stress of ZrB2–SiC layers. While, the rupture work of laminated ZrB2–SiC–G ceramics decreased
Fig. 4. Load-deflection curves of laminated ZrB2–SiC–G specimens.
Table 2 Mechanical properties of the laminated ZrB2–SiC–G ceramics. Materials
LZS-1 LZS-2 LZS-3
ρ
σf
KIC
E
rwof
(g/cm3)
(MPa)
(MPa·m1/2)
(GPa)
(J/m2)
5.18 5.24 5.31
587 ± 23 610 ± 18 670 ± 27
12.1 ± 0.1 13.2 ± 0.3 13.7 ± 0.4
259 ± 19 277 ± 24 293 ± 27
737 ± 28 617 ± 31 480 ± 35
with the addition of heat-resistant particles into graphite layers. Compared to laminated LZS-1 specimens, the rupture work of LZS-3 specimens was reduced to 480 J/m2 by 35%. The curves of retained flexural strength of laminated ZrB2–SiC–G ceramics versus temperature difference (ΔT, exposure temperature minus the water quench temperature) were plotted in Fig. 5. By contrast with room temperature strength of laminated ZrB2–SiC–G ceramics, i.e. the bending strength before thermal shock, the curves of flexural strength for LZS-1 and LZS-3 showed a similar sharp drop as ΔT increased to 200 °C, followed by a modest change as ΔT increased to 400 °C, while that for LZS-2 decreased gradually and slowly. The minimum bending strength was obtained in laminated ZrB2–SiC–G specimens after thermal shock at ΔT of 800 °C and a slight increase was found as ΔT increased to 1000 °C. The critical value of the temperature difference (ΔTC), an important index to evaluate the thermal shock resistance, was determined by a 30% reduction in flexural strength compared to the average bending strength of the tested specimens [27–29]. ΔTC for LZS-1, LZS-2 and LZS-3 was calculated as 470 °C, 508 °C and 400 °C, respectively. The highest ΔTC for LZS-2 specimens, raised by 8% for LZS-1, 27% for LZS-3 and 32% for monolithic ZrB2–SiC [28–30], illustrated that it was the optimized residual thermal stresses between ZrB2–SiC layers and graphite layers that play a key role in the improvement of laminated ZrB2–SiC–G ceramics thermal shock behavior, which revealed that too large or too small residual thermal stress among layers was not conducive to the improvement of thermal shock resistance. To further analyze the reason for low residual strength after thermal shock, the microstructures of LZS-2 specimens under different temperatures of thermal shock were observed. SEM images of LZS-2 specimens after thermal shock at different temperatures are shown in Fig. 6. As indicated, smooth surface with no cracks was shown for specimens quenched at 420 °C (Fig. 6a) and 620 °C (Fig. 6b), indicating superior thermal shock resistance of materials, while obvious cracks with a certain depth were found for specimens quenched at 820 °C (Fig. 6c), indicating that the destruction of materials massive flakes observed at 1020 °C in Fig. 6d was determined to be ZrO2 by EDS analysis. This was also supposed
Fig. 5. Residual strength versus temperature difference for LZS-1, LZS-2, and LZS-3 specimens.
C. Wei, C. Ye / Int. Journal of Refractory Metals and Hard Materials 51 (2015) 233–238
237
b
a
5 µm
c
5 µm
d
5 µm
5 µm
Fig. 6. SEM images of LZS-2 specimens after thermal shock at different temperatures: (a) 420 °C, (b) 620 °C, (c) 820 °C, and (d) 1020 °C.
to be the oxidation product of ZrB2 particles and it's the oxidation product that benefited cracks healing and accordingly accounted for the slight increase of residual strength as ΔT increased from 820 °C to 1020 °C. In order to further evaluate the thermal shock properties of laminated ZrB2–SiC–G ceramics, two thermal shock resistance parameters were used to predict the crack initiation and propagation behavior induced by the thermal shock stress of ceramic matrix composites. One was the thermal stress fracture resistance parameter R, expressed in Eq. (4) in terms of the bending strength of the material (σf), Young's modulus (E), Poisson's ratio (ν) and the coefficient of thermal expansion (α). The thermal shock parameter R predicts the maximum change in temperature (the critical thermal shock temperature difference) that can occur without the initiation of cracks. The other parameter was the thermal stress damage resistance parameter RIV, as expressed in Eq. (5), where σf is the bending strength of the material, KIC is the fracture toughness and ν is the Poisson's ratio, respectively. RIV decides the resistance to catastrophic crack propagation of ceramics under a critical temperature difference. R¼
R
IV
σ f ð1−ν Þ Eα
¼
ð4Þ
2 1 KIC 1−ν σ f
ð5Þ
For the simplicity of the calculation, the magnitudes of the Poisson's ratio (0.13) and the thermal expansion coefficient (7.18 × 10−6/K) for laminated ZrB2–SiC-G ceramics were supposed to be identical. Thermal shock resistance parameters of tested specimens were shown in Table 3. As listed, LZS-2 presented the highest RIV value, 110% of LZS-1, 121% of
LZS-3 and 560% of monolithic ZrB2–SiC ceramics (96 μm) [14], and similar R values were found for LZS-1, LZS-2 and LZS-3, ~182% of R value (147 °C) for monolithic ZrB2–SiC ceramics [14,30], indicating superior thermal shock behavior of laminated ZrB2–SiC–G as demonstrated in Fig. 5. 4. Conclusions The laminated ZrB2–SiC–G ceramics with improved mechanical properties and thermal shock resistance were obtained and investigated in detail. Flexural strength and fracture toughness of the laminated ZrB2–SiC–G ceramics were increased with the addition of heatresistant particles into graphite layers, while the rupture work was decreased gradually. It was found that further improvement of the mechanical properties and thermal shock resistance can be made through optimizing residual thermal stresses by changing the amount of heatresistant particles adding into graphite layers. High performance laminated ZrB2–SiC–G UHTCs were prepared with 30 vol.% ZrB2 and 10 vol.% SiC in graphite layers, with bending strength of 610 MPa, fracture toughness of 13.2 MPa·m1/2, rupture work of 617 J/m2 and critical thermal shock temperature (ΔTC) of 508 °C. Acknowledgment This work was supported by the National Science Foundation of China (50972029) and project (HIT.KLOF.2009026) supported by the Key laboratory Opening Funding of National Key Laboratory of Advanced Composites in Special Environments. References
Table 3 Thermal shock resistance parameters of the laminated ZrB2–SiC–G ceramics. Thermal shock resistance parameters
LZS-1
LZS-2
LZS-3
ΔTC (°C) R (°C) RIV (μm)
470 275 488
508 267 538
400 277 480
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