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Self-toughening behavior of nano yttria partially stabilized hafnia ceramics
Ceramics International 45 (2019) 21467–21474
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Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Self-toughening behavior of nano yttria partially stabilized hafnia ceramics a
Chun Li , Yue Ma
a,b
a,b
, Jian He
, Hongbo Guo
T
a,b,*
a
School of Materials Science and Engineering, Beihang University (BUAA), No. 37 Xueyuan Road, Beijing, 100191, China Key Laboratory of High-temperature Structural Materials & Coatings Technology (Ministry of Industry and Information Technology), Beihang University (BUAA), No. 37 Xueyuan Road, Beijing, 100191, China
b
A R T I C LE I N FO
A B S T R A C T
Keywords: Yttria stabilized hafnia (YSH) Fracture toughness Geometry Finite element simulation (FEM)
Sandwich structure of cubic (C) phase/monoclinic (M) phase/C phase was formed in 8% mol Y2O3 partially stabilized HfO2 (Hf0.92Y0.08O1.96, YSH8) samples which were produced by sintering of YSH8 green body cold pressed from nano-powders of C phase at 1500 °C. The geometry of the sandwich structure was approximately tailored by controlling the sintering time. The YSH8 samples with sandwich structure from nano-powders yielded fracture toughness of 2.1–2.4 MPa m0.5, which was about 33% higher than that of conventional YSH8 sample sintered from micron powders. Both the experimental results and numerical simulation indicated that the fracture toughness is closely related to the geometry of the sandwich structure.
1. Introduction Y2O3 partially stabilized ZrO2 (YSZ) ceramic has long been considered as the most successful thermal insulation material, such as thermal barrier coatings (TBCs) [1–3], mostly because of its excellent mechanical properties and relatively low thermal conductivity (2.3 W/ (m·K) at 1000 °C for dense bulk) [4–7]. Owing to the increasing demand for raising the thrust weight ratio of modern aero engine, it is urgent to raise the service temperature up to 1200 °C, which will cause severe phase transformation of YSZ coating and lead to the failure of TBCs [7–9]. HfO2 based materials are believed as one of the candidate ceramics for TBCs since the chemical properties is quite similar to YSZ, but the thermal phase stability and thermal sintering properties are much better than YSZ [10–15]. Since the fracture toughness of the hafnia is very low regardless of its excellent temperature tolerance, many efforts have been made to improve the fracture toughness of HfO2 based ceramics from different perspectives, including oxides doping [16–20]. The fracture toughness was effectively enhanced by doping of Er2O3 in the case of presence of tetragonal (T) phase [17,18]. In our previous work, the phases of YSH (Y2O3 stabilized HfO2) were accurately controlled by varying the Y3+ doping content [21]. A composite structure consisted of about 53% monoclinic (M) phase and 47% cubic (C) phase was obtained by 8% mol Y2O3 stabilized HfO2 (YSH8) sintered at 1600 °C for 5 h. However, T phase were undetectable by means of XRD, Raman, and TEM. Although the traditional way to use T→M martensitic transformation toughening seems to be ineffective for YSH ceramics, it's still possible to
*
enhance the fracture toughness since the thermal expansion mismatch of M phase and C phase will cause residual stress at the boundary of the M/C grain boundaries [22]. Besides, phase stability of both M phase and C phase are excellent below 1700 °C [10], the potential risk of 3–5% volume expansion caused by catastrophic T→M martensitic phase transformation can be avoided in TBCs. There are still several factors to be considered: 1, how to get tightly bonded M/C interfaces, 2, effects of the microstructure on the fracture toughness as well as the toughening mechanisms, 3, preparation process to control the microstructure. In the present work, we fabricated the partially stabilized YSH ceramics with C/M/C sandwich structure using sol-gel method. Nanopowders composed of mainly C phase goes through spinodal decomposition to form compact M/C boundary. Utilizing the anisotropy of M phase, the morphology of the sandwich structure is optimized during the preparing process. The toughening mechanism and the microstructure design are proposed based on ANSYS. This work provides new insights into microstructure design of dual-phase ceramics to obtain high fracture toughness. 2. Experimental procedure YSH8 nano-powders were synthesized by sol-gel method. HfOCl2·8H2O (98%, Alfa Aesar Chemical Co., Ltd.) and Y(NO3)3·6H2O (99.9%, Aladdin Biochemical Technology Co., Ltd.) were dissolved in deionized water and stirred for 10 min. Citric acid was added into the above solution in a mole ratio of 1:4 (metal cations: citrate). To ensure a
Corresponding author. School of Materials Science and Engineering, Beihang University (BUAA), No. 37 Xueyuan Road, Beijing, 100191, China. E-mail address: [email protected] (H. Guo).
https://doi.org/10.1016/j.ceramint.2019.07.137 Received 18 June 2019; Received in revised form 8 July 2019; Accepted 12 July 2019 Available online 12 July 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Ceramics International 45 (2019) 21467–21474
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better dispersion of the organic precursor, glycol was added afterward. During the reaction process of the solution, water bath heating at 80 °C and vigorous stirring were applied. The acquired wet gel with quite high viscosity was heated at 120 °C for one day to get dried gel, and finally calcined at 1000 °C for 2 h to form nano-crystalline powder. The powder was sieved through a 200-mesh screen and dry-pressed into tablets (Φ 8 mm × 3 mm). The YSH8 ceramic green bodies were sintered at 1500 °C for 2 h, 4 h, 8 h and 16 h (marked as YSH8-2, YSH8-4, YSH8-8 and YSH8-16) respectively to obtain highly dense ceramic bulks with different morphology. For comparison, conventional YSH8 bulk with starting micron powder of HfO2 and Y2O3 was sintered at 1500 °C for 5 h. Phase composition of the specimens was determined by X-ray diffractometry (XRD; Rigaku D/max 2500PC, Japan) in the 2θ range of 20–90° using Cu Kα radiation. The volume fraction of M phase and C phase were calculated by the formula proposed by Yujun Zhang [23]:
Vm =
1.6031Im (1 ‾ 11) 1.6031Im (1 ‾ 11) + Ic (111)
(1)
and Im and Ic are diffraction intensity of monoclinic phase and cubic phase, respectively. Surface morphologies of YSH8 bulks were observed by scanning electron microscopy (SEM, FEI Quanta 600, 20 kV) equipped with energy diffraction spectrum (EDS). Photoshop software was utilized to mark out the typical sandwich structure, and then quantitative statistics of geometric parameters was carried out by Image-Proplus software. Experimental density of YSH8 specimens (ρ) was measured by Archimedes method. The hardness (Hv), fracture toughness (KIC) and fracture energy (Γ) were calculated by indentation method [24] with the following formulas:
KIC = 0.16Hv⋅a0.5⋅(c /a)−3/2 Γ = 2ξ 2⋅P⋅
d2 c3
Hv = 0.464
P d2
(2) (3) (4)
where a, c, P and ξ = 0.016 stand for half the length of the indentation diagonal, the crack length from crack tip to the indentation center, loading force, and the geometry factor of the Vickers penetrator. And ultrasonic method was utilized to measure the elastic modulus (E) of YSH20 with C phase and HfO2 with M phase [21,25]:
4υt2 − 3υl2 υt2 − υl2
(5)
E (1 − ϕ2/3)1.21
(6)
E = ρvt2 E0 =
where υt and υl are the transverse and longitudinal wave speed, respectively. ϕ, E and E0 represent the porosity, experimental elastic modulus, and the theoretical elastic modulus. 3. Results and discussion 3.1. Formation and evolution of sandwich structure As seen in the XRD patterns (Fig. 1), the YSH8 nano-powder is fully stabilized to C phase owning to the lower surface energy since the particle size is less than 50 nm [10]. At temperatures below the tetragonal to monoclinic transformation temperature, a tetragonal HfO2 particle is metastable when its size is smaller than a critical size. Our previous work has demonstrated that the doping of Y element stabilizes the hafnia to cubic phase rather than tetragonal phase, and it is also indicated by first-principles calculation that the doping of rare earth oxides as well as some of the cations with rather large ion radius stabilizes hafnia to cubic phase [21,26]. For the YSH ceramics, to fully stabilize hafnia to cubic phase, the concentration of Y3+ needs to be
Fig. 1. XRD patterns of YSH8 nano-powders before and after annealing at 1500 °C for 2 h.
higher than 17.1% (sintered at 1600 °C). But when YSH8 is in nano scale, the surface effect helps to fully stabilize it to C phase despite its rather low Y3+ concentration (see Fig. 2). When the YSH8 nano-powder is exposed to high temperature, the grain size of the nano-crystals grows to a critical grain size, above which the C phase is no longer stable due to the relatively low Y3+ doping concentration. As the surface of the two relatively large metastable C phase grains contact, sintering and ions diffusion occurs at the grain boundaries at 1500 °C. The Y3+ with higher chemistry activity is likely to move to the surface from neck as seen in the schematic diagram in Fig. 3, leaving a low Y3+ zone at the neck area, which promotes the nucleation of M phase with very low concentration of Y3+.At the same time, the primary cation Hf4+ moves from inner parts of the grains towards the neck area to achieve densification. Since stable M phase hafnia belongs to the monoclinic-prismatic class, the growth rate of basal planes is much higher than that of prism planes, forming lath-shaped M phase morphology. The difference in growth rate is related to the monoclinic crystal structure, which can be attributed to an energetically more favorable attachment of a surface nucleus on a basal plane [27]. Considering the anisotropy of M phase during grain growth, C/M/C sandwich structure was formed, which is schematically presented in Fig. 3. Fig. 4 shows the XRD patterns of the YSH8 bulks sintered for 2–16 h. The diffraction peaks are very similar despite of the different sintering time, indicating very similar phase composition of the specimens. All the samples were composed of M phase and C phase, and the phase volume fraction were listed in Table 1. The volume fraction of M phase was about 53%, which was close to that of the conventional YSH8 from the micron powders. Note that the volume fraction of M phase in the dense bulk (YSH8-2) is higher than that of the powder after sintered for 2 h (Figs. 1 and 4). The phase transformation of the metastable C phase to the C + M phases is controlled mainly by the diffusion of the cations (Hf4+ and Y3+). When the YSH8 powders were pressed into tablets, the binding of the particles within the tablet is enhanced, and the contact areas are also much bigger compared with that of the freestanding powders. The tablet provides better diffusion condition for cations, as a result, the phase transformation is more sufficient and thus more M phase grains are present. The identical phase composition in the YSH8 ceramics ensures that the fracture toughness discussed below is influenced by the effect of the grain geometry instead of the contents of M phase and C phase. The morphologies of YSH8 bulks sintered for different time are shown in Fig. 5, and the experimental density and the porosity of the specimens were listed in Table 2. The YSH8 bulks sintered from nanopowders (Fig. 5a–d) are quite dense compared with that from micron powders (Fig. 5f), and the microstructure is very homogeneous, mainly
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Fig. 2. Electron back-scattering micrographs of nano-powders annealed at 1500 °C for 2 h: (a) unpolished surface and (b) cross section.
due to the much smaller grain size which enables the more efficient diffusion of ions during sintering. The grain sizes increase as the sintering time increases from 2 h to 8 h, but changes very slowly when prolonging the sintering time to 16 h. In order to distinguish M phase from C phase, EDS was utilized to analyze the chemical composition of the grains in the SEM image, and those with low Y element (< 1 at.%) belongs to M phase, otherwise they are C phase. The grains in the ceramics are generally divided into three groups (Fig. 5c): C/M/C phase sandwich structure (I), M phase cluster grain (II), and C phase grain with flat surface (III). The anisotropic growth of M phase crystal produced rather flat interface in the sandwich structure. We made detailed statistics of the volume fraction of the sandwich structure grains for the YSH8 ceramics on the SEM images, and found that they account for 8%–10% volume fraction of the YSH8 ceramics no matter how many hours during the sintering process. When the YSH8 nano grains are exposed to high temperature, the formation of the sandwich structure is one of the grain growth modes. In addition, the Y3+ in a particle moves to the adjacent particles, leaving very low Y3+ concentration in the original particle but relatively high Y3+ concentration in the around particles, this is why separate M phase and C phase grains are formed. Moreover, the formation of the separate M phase and C phase grains seems to be the primary mode during the grain growth of the YSH8 ceramics since their volume fractions are higher than that of the grains with sandwich structure. Thermal expansion mismatch between M phase and C phase will cause significant stress at the C/M/C interfaces. Considering the relatively low grain boundary binding force, the toughening effect of YSH8 is mainly contributed from the residual stress formed in these I-typed sandwich structure since the two phases are firmly bonded. Although the phase composition of YSH8 ceramics sintered from nano-powders and micron powders are almost the same, particles with sandwich structure were not observed in the latter. With starting powders of HfO2 (M phase) and Y2O3 (C phase), both of which are in
Fig. 4. XRD patterns of YSH8 bulks sintered at 1500 °C for 2–16 h. Table 1 Volume fractions of the C phase and M phase in the YSH8 ceramics. Specimen
YSH8-2
YSH8-4
YSH8-8
YSH8-16
M phase C phase
44.9% 55.1%
42.4% 57.6%
45.0% 55.0%
43.6% 56.4%
the range of 0.5–3 μm, Y3+ diffused into the crystal cell of the nearest M phase HfO2, causing the formation of C phase YSH and the disappearance of C phase Y2O3 during sintering. And those HfO2 particles around which the Y2O3 was unattainable remained in the bulk and only grain growth was accomplished. Thus, several conditions are necessary
Fig. 3. Schematic diagram showing the formation of sandwich structure. 21469
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Fig. 5. (a–d) Hot etched surface morphology of the YSH8 sintered for 2–16 h, (e) SEM of YSH8-4 treated by Photoshop showing the typical sandwich structures, and (f) the surface morphology of conventional YSH8 bulk from micron powders.
3.2. Fracture toughness
Table 2 Density, porosity, Young's modulus and Vicker's hardness in YSH8 ceramics. Specimen
YSH8-2
YSH8-4
YSH8-8
YSH8-16
Density/(g/cm3) Porosity/% Young's modulus/GPa Vicker's Hardness/GPa
9.25 6.2 181 8.02
9.18 6.9 187 8.04
9.42 4.5 209 9.01
9.21 6.6 199 9.27
for the formation of C/M/C sandwich structure: very small grain size (< 100 μm), chemical composition in the range of M phase and C phase (partially doped if stable), homogenous metastable solid solution. The sandwich structure is composed of a lath-shaped M phase inside and two semisphere outside (Fig. 5e). b/a is the value of length to width ratio of M phase, and SC/SM stands for the volume ratio (also area ratio, for simplification) of the corresponding C phase and M phase. In order to control the morphology of the sandwich structure by adjusting the sintering time, as shown in Fig. 6, b/a and SC/SM were counted quantitively according to the Photoshop treated pictures (Fig. 5e, for example). At least 100 monomer sandwich structures were counted for each sample to ensure the statistic is credible to reflect the intrinsic morphology. SC/SM primarily increases as b/a of M phase increases (the black dash line in Fig. 6a). Fig. 6b shows the cumulative probability of M phase as a function of b/a. The b/a distribution is identical for the YSH8 ceramics when b/a < 5 and b/a > 10, but is remarkably different when 5 < b/a < 10. The b/a value first increases and then decreases as the sintering time increases from 2 h to 4 h and 4 h–16 h, respectively. Generally, the average b/a increases when sintering time is less than 4 h and afterward decreases with sintering time as seen in Fig. 6c. As the schematic diagram demonstrates (Fig. 7), the grain size is quite small for the YSH8-2, and the M phase grains tend to grow along the radial direction as sintering time prolongs, resulting in a higher b/a value for YSH8-4. A network of rectangle shaped M phase is formed after sintered for 8 h, and the growth of M phase grains along the radial direction is inhibited by the adjacent M phase grains. Consequently, the b/a value is reduced due to the growth along the width direction. This is also the reason for the relatively lower grain growth rate for YSH8-8 than YSH82.
Although the experimental Young's modulus and Vicker's hardness increase slightly with the sintering time (Table 2), fracture toughness shows quite different features. Fig. 8 shows the fracture toughness and the toughening mechanisms of the ceramic bulks. Fracture toughness of YSH8 specimens exceeds 2.1 MPa m0.5, which is 33% higher than the YSH8 ceramic from micron powders mentioned before ascribed to the sandwich structure within the bulks. Moreover, it exceeds many of the potential thermal barrier coating materials, such as Gd2Zr2O7 and La2Ce2O7 (~0.8 MPa m1/2 and 1.3–1.5 MPa m1/2, measured by indentation method as in this research) [28,29], because of the unique sandwich structure which introduces high compressive stress in the M phase near the M/C interface. The crack prolongation modes, as shown in Fig. 8b–f, illustrate that many toughening mechanisms contribute to the relatively high fracture toughness. Crack bridging occurs when the crack prolongs in the ceramic, generating crack closure force between the two interfaces, and finally the stress intensity factor is enhanced with the growth of the cracks. Besides, the length of the crack propagation path lengthens if second crack forms in the previous grain as the crack enters the adjacent grain. Crack deflection is another important factor for the toughening of YSH8. As shown in Fig. 8d–f, when the crack propagates near the C/M/C sandwich structure, there's a tendency for the crack to deflect and extend along the interface of C phase and M phase. During crack deflection, the crack plane will reorient perpendicular to the direction of applied tension stress, which means that the crack propagation path will be increased. At the same time, the stress at the crack tip decreases because the crack plane is no longer perpendicular to the direction of tensile stress. As the sintering time increases, the fracture toughness increases first but decreases when the sintering time exceeds 4 h. The fracture toughness of YSH8-4 is 15% higher than that of YSH8-16 (2.4 MPa m0.5 and 2.1 MPa m0.5, respectively). Like the changing tendency of fracture toughness of YSH8 ceramics with the sintering time, the fracture energy also increases first with the sintering time, and reaches ~23 J m−2 for YSH8-4, but then decreases as the sintering time increases from 4 h to 16 h. The fracture energy of YSH8-4 is 44% higher than that of YSH816, this is also why the fracture toughness of YSH8-4 is higher than that of YSH8-16. Generally speaking, fracture energy increases with the
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Fig. 6. (a) Scatter plot of SC/SM as a function of b/a for each sandwich structure grain and (b) the cumulative distribution curve of b/a of M phase. (c) Average b/a with the sintering time.
grain growth for the brittle oxide ceramics due to the decrease of the grain boundaries which weakens the binding stress of the grains. Thus, the grain size effect fails to explain the variation of fracture toughness with the sintering time.
3.3. Finite element simulation Finite element simulation (FEM) of the stress distribution in C phase and M phase at the interfaces of C/M/C sandwich structure base on ANSYS was thus carried out to understand the influence of geometry on the fracture toughness. To simplify the FEM process, the morphology geometry of the sandwich structure in YSH8 ceramics was extracted
from the 2D SEM images. As seen in the SEM images (Fig. 5e), the C/M/ C sandwich structure is composed of a lath-shaped M phase inside and two semispheres outside, in order to avoid the singularity at the corner of the C phase semisphere, the C phase is modelled as the geometry in Fig. 9a. The SC/SM = 2.4L/a is calculated by the geometry in Fig. 9a. By varying the value of a, b and L, different geometry with various b/a and SC/SM values are established for the FEM analysis. And the variables of b/a changing from 2 to 8 is based on the statistics that the b/a in the YSH8 specimens under SEM images is mainly in this range. This is also why the SC/SM is studied in the range of 0.5–5. The elements attached to the symmetrical axis which is parallel to the Y axis are immobile in the X direction as shown in Fig. 9a. And the bottom point of the above
Fig. 7. Schematic diagram of microstructure evolution for YSH8 during annealing at 1500 °C. 21471
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Fig. 8. (a) The Fracture toughness and fracture energy of YSH8 ceramics and (b)–(f) the toughening mechanisms in YSH8.
Fig. 9. Finite element simulation of sandwich structures with various geometries. (a) Geometric model of the C/M/C structure for finite element simulation, and (b) the residual stress at the boundary of M and C phase (indicated by the red box) as a function of SC/SM when b/a = 4. (c) The comprehensive stress in M phase as a function of SC/SM for various b/a models. (d) The critical SC/SM determined by the method in Fig. 3b and the corresponding steady stress value in M phase as a function of b/a. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.) Table 3 Material properties of M phase and C phase [21]. Material
Density (g/cm3)
Young's modulus (GPa)
Poisson's ratio
Average thermal expansion coefficient (10−6/K)
M phase C phase
10.09 9.59
294 267
0.294 0.304
6.84 10.44
points is immobile in all directions. Besides, there's no relative motion between the adjacent elements at the boundary of M phase and C phase since the overlapping lines of M phase and C phase are “glued”. Material properties of M phase and C phase are listed in Table 3. Thermal
residual stress after cooling from 1500 °C (stress-free) to 25 °C during sintering was calculated by FEM analysis. By changing the length of C phase (L) while keeping the b and a of M phase constant, the stress states as a function of SC/SM are plotted in
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Fig. 9b. When b/a = 4, the compressive stress in M phase increases dramatically as SC/SM increases, but keeps constant when SC/SM > 2. Residual tensile stress is generated within the C phase, and is prone to promote the crack extension because the tensile stress enlarges the level of stress concentration at the crack tip. However, due to the existence of the residual tensile stress and the adjacent compressive region, the crack is easily deflected along the C phase where thermal tensile stress is present (Fig. 8d–f). Moreover, the FEM analysis indicates that the residual tensile stress in C phase is relatively low compared with the compressive stress in M phase when SC/SM > 1. The tensile stress in C phase generally decreases but changes very slowly when SC/SM > 2. The increasing compressive stress in M phase and the decreasing stress in C phase both help to resist the crack propagation and enhance the fracture toughness of the YSH8. Obviously when the area ratio is higher than the critical SC/SM = 1.2 value (point A) for the case b/a = 4, a steady stress value in M phase is obtained, and further increasing the SC/SM value has little effect on the fracture toughness. Fig. 9c reveals the compressive stress in M phase as a function of the SC/SM value for different b/a ranging from 2 to 8. For each sandwich structure with different b/a values, there's a critical SC/SM, above which a steady stress in M phase is obtained. Extracted from Fig. 9c and d shows the critical SC/SM value and the steady stress in M phase as a function of b/a. As the b/a value increases, the SC/SM must be enhanced to achieve the steady stress, which also increases with the b/a value. In that way, as the M phase in the sandwich structure becomes thinner, the compressive stress will be higher and contributes more to the toughening of the ceramics, but the SC/SM value must be higher than the corresponding critical value. Besides, the steady stress for b/a = 8 (~1100 MPa in compress) is 57% higher than that of b/a = 2 (~700 MPa), indicating the toughening effects of the residual stress is quite significant when the M phase is of high length to width ratio. The critical SC/SM as a function of b/a determined by the FEM results is plotted in Fig. 6a, which is lower than most of the SC/SM values in the samples, indicating that the stress state of M phase in the C/M/C sandwich structure can be simplified only as a function of the b/a value since a steady stress is reached. When the crack is going to enter M phase, the compressive stress will reduce the crack propagation driving force, thus enhances the fracture toughness. The total width (b) of the M phase within the sandwich structure declines with the sintering time due to the grain growth and the simultaneous reduction of grain boundary, resulting a reducing tendency of the toughening effect. However, the residual compressive stress at the boundary of the M phase shows first increase but then decrease as the sintering time increases as illustrated in Fig. 6. Compared with YSH8-2, YSH8-4 specimen maintains a relatively higher b/a value but lower total width of the M phase in the sandwich structure. The higher fracture toughness of YSH8-4 compared with that of YSH8-2 demonstrates that the increased compressive stress of M phase seems to be the dominant factor for the change of fracture toughness rather than the decreased total width of the M phase at the initial sintering process. But when the sintering time exceeds 4 h, the fracture toughness decreases with sintering time, resulted from the decreased compressive stress alone with the decreased total width of M phase. Furthermore, the fracture toughness of dual-phase ceramics could be tailored by controlling the preparing process not just for the YSH8 ceramics, and FEM is an effective method to roughly estimate the residual stress within the ceramics and contribute to the design of ceramics with high fracture toughness. 4. Conclusion In this work, 8% mol Y2O3 partially stabilized HfO2 (YSH8) dense sample featured with a novel C phase/M phase (lath shaped)/C phase self-toughening sandwich structure was successfully prepared by solid state sintering of nano-powders of C phase. The geometry of the sandwich structure was approximately tailored by controlling sintering time
at 1500 °C. The YSH8 samples with sandwich structure yielded fracture toughness of 2.1–2.4 MPa m0.5, which was 33% higher than that of conventional YSH8 from micron powders, owing to the toughening mechanism of crack deflection and bridging within C/M/C sandwich structure, releasing the stress concentration at the crack tip. With increasing the length to width ratio (b/a) of the lath M phase, the experimental result and the FEM simulation confirm that the toughening effect was significantly enhanced. Acknowledgements This research is sponsored by National Natural Science Foundation of China (NSFC) under grant U1537212, No. 51471019, No. 51590894 and No. 51425102, and the National Key Research and Development Program of China under grant No. 2016YFB0300901. References [1] N.P. Padture, Thermal barrier coatings for gas-turbine engine applications, Science 296 (80) (2002) 280–284, https://doi.org/10.1126/science.1068609. [2] A.K. Gain, H.-Y. Song, B.-T. 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Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Self-toughening behavior of nano yttria partially stabilized hafnia ceramics a
Chun Li , Yue Ma
a,b
a,b
, Jian He
, Hongbo Guo
T
a,b,*
a
School of Materials Science and Engineering, Beihang University (BUAA), No. 37 Xueyuan Road, Beijing, 100191, China Key Laboratory of High-temperature Structural Materials & Coatings Technology (Ministry of Industry and Information Technology), Beihang University (BUAA), No. 37 Xueyuan Road, Beijing, 100191, China
b
A R T I C LE I N FO
A B S T R A C T
Keywords: Yttria stabilized hafnia (YSH) Fracture toughness Geometry Finite element simulation (FEM)
Sandwich structure of cubic (C) phase/monoclinic (M) phase/C phase was formed in 8% mol Y2O3 partially stabilized HfO2 (Hf0.92Y0.08O1.96, YSH8) samples which were produced by sintering of YSH8 green body cold pressed from nano-powders of C phase at 1500 °C. The geometry of the sandwich structure was approximately tailored by controlling the sintering time. The YSH8 samples with sandwich structure from nano-powders yielded fracture toughness of 2.1–2.4 MPa m0.5, which was about 33% higher than that of conventional YSH8 sample sintered from micron powders. Both the experimental results and numerical simulation indicated that the fracture toughness is closely related to the geometry of the sandwich structure.
1. Introduction Y2O3 partially stabilized ZrO2 (YSZ) ceramic has long been considered as the most successful thermal insulation material, such as thermal barrier coatings (TBCs) [1–3], mostly because of its excellent mechanical properties and relatively low thermal conductivity (2.3 W/ (m·K) at 1000 °C for dense bulk) [4–7]. Owing to the increasing demand for raising the thrust weight ratio of modern aero engine, it is urgent to raise the service temperature up to 1200 °C, which will cause severe phase transformation of YSZ coating and lead to the failure of TBCs [7–9]. HfO2 based materials are believed as one of the candidate ceramics for TBCs since the chemical properties is quite similar to YSZ, but the thermal phase stability and thermal sintering properties are much better than YSZ [10–15]. Since the fracture toughness of the hafnia is very low regardless of its excellent temperature tolerance, many efforts have been made to improve the fracture toughness of HfO2 based ceramics from different perspectives, including oxides doping [16–20]. The fracture toughness was effectively enhanced by doping of Er2O3 in the case of presence of tetragonal (T) phase [17,18]. In our previous work, the phases of YSH (Y2O3 stabilized HfO2) were accurately controlled by varying the Y3+ doping content [21]. A composite structure consisted of about 53% monoclinic (M) phase and 47% cubic (C) phase was obtained by 8% mol Y2O3 stabilized HfO2 (YSH8) sintered at 1600 °C for 5 h. However, T phase were undetectable by means of XRD, Raman, and TEM. Although the traditional way to use T→M martensitic transformation toughening seems to be ineffective for YSH ceramics, it's still possible to
*
enhance the fracture toughness since the thermal expansion mismatch of M phase and C phase will cause residual stress at the boundary of the M/C grain boundaries [22]. Besides, phase stability of both M phase and C phase are excellent below 1700 °C [10], the potential risk of 3–5% volume expansion caused by catastrophic T→M martensitic phase transformation can be avoided in TBCs. There are still several factors to be considered: 1, how to get tightly bonded M/C interfaces, 2, effects of the microstructure on the fracture toughness as well as the toughening mechanisms, 3, preparation process to control the microstructure. In the present work, we fabricated the partially stabilized YSH ceramics with C/M/C sandwich structure using sol-gel method. Nanopowders composed of mainly C phase goes through spinodal decomposition to form compact M/C boundary. Utilizing the anisotropy of M phase, the morphology of the sandwich structure is optimized during the preparing process. The toughening mechanism and the microstructure design are proposed based on ANSYS. This work provides new insights into microstructure design of dual-phase ceramics to obtain high fracture toughness. 2. Experimental procedure YSH8 nano-powders were synthesized by sol-gel method. HfOCl2·8H2O (98%, Alfa Aesar Chemical Co., Ltd.) and Y(NO3)3·6H2O (99.9%, Aladdin Biochemical Technology Co., Ltd.) were dissolved in deionized water and stirred for 10 min. Citric acid was added into the above solution in a mole ratio of 1:4 (metal cations: citrate). To ensure a
Corresponding author. School of Materials Science and Engineering, Beihang University (BUAA), No. 37 Xueyuan Road, Beijing, 100191, China. E-mail address: [email protected] (H. Guo).
https://doi.org/10.1016/j.ceramint.2019.07.137 Received 18 June 2019; Received in revised form 8 July 2019; Accepted 12 July 2019 Available online 12 July 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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better dispersion of the organic precursor, glycol was added afterward. During the reaction process of the solution, water bath heating at 80 °C and vigorous stirring were applied. The acquired wet gel with quite high viscosity was heated at 120 °C for one day to get dried gel, and finally calcined at 1000 °C for 2 h to form nano-crystalline powder. The powder was sieved through a 200-mesh screen and dry-pressed into tablets (Φ 8 mm × 3 mm). The YSH8 ceramic green bodies were sintered at 1500 °C for 2 h, 4 h, 8 h and 16 h (marked as YSH8-2, YSH8-4, YSH8-8 and YSH8-16) respectively to obtain highly dense ceramic bulks with different morphology. For comparison, conventional YSH8 bulk with starting micron powder of HfO2 and Y2O3 was sintered at 1500 °C for 5 h. Phase composition of the specimens was determined by X-ray diffractometry (XRD; Rigaku D/max 2500PC, Japan) in the 2θ range of 20–90° using Cu Kα radiation. The volume fraction of M phase and C phase were calculated by the formula proposed by Yujun Zhang [23]:
Vm =
1.6031Im (1 ‾ 11) 1.6031Im (1 ‾ 11) + Ic (111)
(1)
and Im and Ic are diffraction intensity of monoclinic phase and cubic phase, respectively. Surface morphologies of YSH8 bulks were observed by scanning electron microscopy (SEM, FEI Quanta 600, 20 kV) equipped with energy diffraction spectrum (EDS). Photoshop software was utilized to mark out the typical sandwich structure, and then quantitative statistics of geometric parameters was carried out by Image-Proplus software. Experimental density of YSH8 specimens (ρ) was measured by Archimedes method. The hardness (Hv), fracture toughness (KIC) and fracture energy (Γ) were calculated by indentation method [24] with the following formulas:
KIC = 0.16Hv⋅a0.5⋅(c /a)−3/2 Γ = 2ξ 2⋅P⋅
d2 c3
Hv = 0.464
P d2
(2) (3) (4)
where a, c, P and ξ = 0.016 stand for half the length of the indentation diagonal, the crack length from crack tip to the indentation center, loading force, and the geometry factor of the Vickers penetrator. And ultrasonic method was utilized to measure the elastic modulus (E) of YSH20 with C phase and HfO2 with M phase [21,25]:
4υt2 − 3υl2 υt2 − υl2
(5)
E (1 − ϕ2/3)1.21
(6)
E = ρvt2 E0 =
where υt and υl are the transverse and longitudinal wave speed, respectively. ϕ, E and E0 represent the porosity, experimental elastic modulus, and the theoretical elastic modulus. 3. Results and discussion 3.1. Formation and evolution of sandwich structure As seen in the XRD patterns (Fig. 1), the YSH8 nano-powder is fully stabilized to C phase owning to the lower surface energy since the particle size is less than 50 nm [10]. At temperatures below the tetragonal to monoclinic transformation temperature, a tetragonal HfO2 particle is metastable when its size is smaller than a critical size. Our previous work has demonstrated that the doping of Y element stabilizes the hafnia to cubic phase rather than tetragonal phase, and it is also indicated by first-principles calculation that the doping of rare earth oxides as well as some of the cations with rather large ion radius stabilizes hafnia to cubic phase [21,26]. For the YSH ceramics, to fully stabilize hafnia to cubic phase, the concentration of Y3+ needs to be
Fig. 1. XRD patterns of YSH8 nano-powders before and after annealing at 1500 °C for 2 h.
higher than 17.1% (sintered at 1600 °C). But when YSH8 is in nano scale, the surface effect helps to fully stabilize it to C phase despite its rather low Y3+ concentration (see Fig. 2). When the YSH8 nano-powder is exposed to high temperature, the grain size of the nano-crystals grows to a critical grain size, above which the C phase is no longer stable due to the relatively low Y3+ doping concentration. As the surface of the two relatively large metastable C phase grains contact, sintering and ions diffusion occurs at the grain boundaries at 1500 °C. The Y3+ with higher chemistry activity is likely to move to the surface from neck as seen in the schematic diagram in Fig. 3, leaving a low Y3+ zone at the neck area, which promotes the nucleation of M phase with very low concentration of Y3+.At the same time, the primary cation Hf4+ moves from inner parts of the grains towards the neck area to achieve densification. Since stable M phase hafnia belongs to the monoclinic-prismatic class, the growth rate of basal planes is much higher than that of prism planes, forming lath-shaped M phase morphology. The difference in growth rate is related to the monoclinic crystal structure, which can be attributed to an energetically more favorable attachment of a surface nucleus on a basal plane [27]. Considering the anisotropy of M phase during grain growth, C/M/C sandwich structure was formed, which is schematically presented in Fig. 3. Fig. 4 shows the XRD patterns of the YSH8 bulks sintered for 2–16 h. The diffraction peaks are very similar despite of the different sintering time, indicating very similar phase composition of the specimens. All the samples were composed of M phase and C phase, and the phase volume fraction were listed in Table 1. The volume fraction of M phase was about 53%, which was close to that of the conventional YSH8 from the micron powders. Note that the volume fraction of M phase in the dense bulk (YSH8-2) is higher than that of the powder after sintered for 2 h (Figs. 1 and 4). The phase transformation of the metastable C phase to the C + M phases is controlled mainly by the diffusion of the cations (Hf4+ and Y3+). When the YSH8 powders were pressed into tablets, the binding of the particles within the tablet is enhanced, and the contact areas are also much bigger compared with that of the freestanding powders. The tablet provides better diffusion condition for cations, as a result, the phase transformation is more sufficient and thus more M phase grains are present. The identical phase composition in the YSH8 ceramics ensures that the fracture toughness discussed below is influenced by the effect of the grain geometry instead of the contents of M phase and C phase. The morphologies of YSH8 bulks sintered for different time are shown in Fig. 5, and the experimental density and the porosity of the specimens were listed in Table 2. The YSH8 bulks sintered from nanopowders (Fig. 5a–d) are quite dense compared with that from micron powders (Fig. 5f), and the microstructure is very homogeneous, mainly
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Fig. 2. Electron back-scattering micrographs of nano-powders annealed at 1500 °C for 2 h: (a) unpolished surface and (b) cross section.
due to the much smaller grain size which enables the more efficient diffusion of ions during sintering. The grain sizes increase as the sintering time increases from 2 h to 8 h, but changes very slowly when prolonging the sintering time to 16 h. In order to distinguish M phase from C phase, EDS was utilized to analyze the chemical composition of the grains in the SEM image, and those with low Y element (< 1 at.%) belongs to M phase, otherwise they are C phase. The grains in the ceramics are generally divided into three groups (Fig. 5c): C/M/C phase sandwich structure (I), M phase cluster grain (II), and C phase grain with flat surface (III). The anisotropic growth of M phase crystal produced rather flat interface in the sandwich structure. We made detailed statistics of the volume fraction of the sandwich structure grains for the YSH8 ceramics on the SEM images, and found that they account for 8%–10% volume fraction of the YSH8 ceramics no matter how many hours during the sintering process. When the YSH8 nano grains are exposed to high temperature, the formation of the sandwich structure is one of the grain growth modes. In addition, the Y3+ in a particle moves to the adjacent particles, leaving very low Y3+ concentration in the original particle but relatively high Y3+ concentration in the around particles, this is why separate M phase and C phase grains are formed. Moreover, the formation of the separate M phase and C phase grains seems to be the primary mode during the grain growth of the YSH8 ceramics since their volume fractions are higher than that of the grains with sandwich structure. Thermal expansion mismatch between M phase and C phase will cause significant stress at the C/M/C interfaces. Considering the relatively low grain boundary binding force, the toughening effect of YSH8 is mainly contributed from the residual stress formed in these I-typed sandwich structure since the two phases are firmly bonded. Although the phase composition of YSH8 ceramics sintered from nano-powders and micron powders are almost the same, particles with sandwich structure were not observed in the latter. With starting powders of HfO2 (M phase) and Y2O3 (C phase), both of which are in
Fig. 4. XRD patterns of YSH8 bulks sintered at 1500 °C for 2–16 h. Table 1 Volume fractions of the C phase and M phase in the YSH8 ceramics. Specimen
YSH8-2
YSH8-4
YSH8-8
YSH8-16
M phase C phase
44.9% 55.1%
42.4% 57.6%
45.0% 55.0%
43.6% 56.4%
the range of 0.5–3 μm, Y3+ diffused into the crystal cell of the nearest M phase HfO2, causing the formation of C phase YSH and the disappearance of C phase Y2O3 during sintering. And those HfO2 particles around which the Y2O3 was unattainable remained in the bulk and only grain growth was accomplished. Thus, several conditions are necessary
Fig. 3. Schematic diagram showing the formation of sandwich structure. 21469
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Fig. 5. (a–d) Hot etched surface morphology of the YSH8 sintered for 2–16 h, (e) SEM of YSH8-4 treated by Photoshop showing the typical sandwich structures, and (f) the surface morphology of conventional YSH8 bulk from micron powders.
3.2. Fracture toughness
Table 2 Density, porosity, Young's modulus and Vicker's hardness in YSH8 ceramics. Specimen
YSH8-2
YSH8-4
YSH8-8
YSH8-16
Density/(g/cm3) Porosity/% Young's modulus/GPa Vicker's Hardness/GPa
9.25 6.2 181 8.02
9.18 6.9 187 8.04
9.42 4.5 209 9.01
9.21 6.6 199 9.27
for the formation of C/M/C sandwich structure: very small grain size (< 100 μm), chemical composition in the range of M phase and C phase (partially doped if stable), homogenous metastable solid solution. The sandwich structure is composed of a lath-shaped M phase inside and two semisphere outside (Fig. 5e). b/a is the value of length to width ratio of M phase, and SC/SM stands for the volume ratio (also area ratio, for simplification) of the corresponding C phase and M phase. In order to control the morphology of the sandwich structure by adjusting the sintering time, as shown in Fig. 6, b/a and SC/SM were counted quantitively according to the Photoshop treated pictures (Fig. 5e, for example). At least 100 monomer sandwich structures were counted for each sample to ensure the statistic is credible to reflect the intrinsic morphology. SC/SM primarily increases as b/a of M phase increases (the black dash line in Fig. 6a). Fig. 6b shows the cumulative probability of M phase as a function of b/a. The b/a distribution is identical for the YSH8 ceramics when b/a < 5 and b/a > 10, but is remarkably different when 5 < b/a < 10. The b/a value first increases and then decreases as the sintering time increases from 2 h to 4 h and 4 h–16 h, respectively. Generally, the average b/a increases when sintering time is less than 4 h and afterward decreases with sintering time as seen in Fig. 6c. As the schematic diagram demonstrates (Fig. 7), the grain size is quite small for the YSH8-2, and the M phase grains tend to grow along the radial direction as sintering time prolongs, resulting in a higher b/a value for YSH8-4. A network of rectangle shaped M phase is formed after sintered for 8 h, and the growth of M phase grains along the radial direction is inhibited by the adjacent M phase grains. Consequently, the b/a value is reduced due to the growth along the width direction. This is also the reason for the relatively lower grain growth rate for YSH8-8 than YSH82.
Although the experimental Young's modulus and Vicker's hardness increase slightly with the sintering time (Table 2), fracture toughness shows quite different features. Fig. 8 shows the fracture toughness and the toughening mechanisms of the ceramic bulks. Fracture toughness of YSH8 specimens exceeds 2.1 MPa m0.5, which is 33% higher than the YSH8 ceramic from micron powders mentioned before ascribed to the sandwich structure within the bulks. Moreover, it exceeds many of the potential thermal barrier coating materials, such as Gd2Zr2O7 and La2Ce2O7 (~0.8 MPa m1/2 and 1.3–1.5 MPa m1/2, measured by indentation method as in this research) [28,29], because of the unique sandwich structure which introduces high compressive stress in the M phase near the M/C interface. The crack prolongation modes, as shown in Fig. 8b–f, illustrate that many toughening mechanisms contribute to the relatively high fracture toughness. Crack bridging occurs when the crack prolongs in the ceramic, generating crack closure force between the two interfaces, and finally the stress intensity factor is enhanced with the growth of the cracks. Besides, the length of the crack propagation path lengthens if second crack forms in the previous grain as the crack enters the adjacent grain. Crack deflection is another important factor for the toughening of YSH8. As shown in Fig. 8d–f, when the crack propagates near the C/M/C sandwich structure, there's a tendency for the crack to deflect and extend along the interface of C phase and M phase. During crack deflection, the crack plane will reorient perpendicular to the direction of applied tension stress, which means that the crack propagation path will be increased. At the same time, the stress at the crack tip decreases because the crack plane is no longer perpendicular to the direction of tensile stress. As the sintering time increases, the fracture toughness increases first but decreases when the sintering time exceeds 4 h. The fracture toughness of YSH8-4 is 15% higher than that of YSH8-16 (2.4 MPa m0.5 and 2.1 MPa m0.5, respectively). Like the changing tendency of fracture toughness of YSH8 ceramics with the sintering time, the fracture energy also increases first with the sintering time, and reaches ~23 J m−2 for YSH8-4, but then decreases as the sintering time increases from 4 h to 16 h. The fracture energy of YSH8-4 is 44% higher than that of YSH816, this is also why the fracture toughness of YSH8-4 is higher than that of YSH8-16. Generally speaking, fracture energy increases with the
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Fig. 6. (a) Scatter plot of SC/SM as a function of b/a for each sandwich structure grain and (b) the cumulative distribution curve of b/a of M phase. (c) Average b/a with the sintering time.
grain growth for the brittle oxide ceramics due to the decrease of the grain boundaries which weakens the binding stress of the grains. Thus, the grain size effect fails to explain the variation of fracture toughness with the sintering time.
3.3. Finite element simulation Finite element simulation (FEM) of the stress distribution in C phase and M phase at the interfaces of C/M/C sandwich structure base on ANSYS was thus carried out to understand the influence of geometry on the fracture toughness. To simplify the FEM process, the morphology geometry of the sandwich structure in YSH8 ceramics was extracted
from the 2D SEM images. As seen in the SEM images (Fig. 5e), the C/M/ C sandwich structure is composed of a lath-shaped M phase inside and two semispheres outside, in order to avoid the singularity at the corner of the C phase semisphere, the C phase is modelled as the geometry in Fig. 9a. The SC/SM = 2.4L/a is calculated by the geometry in Fig. 9a. By varying the value of a, b and L, different geometry with various b/a and SC/SM values are established for the FEM analysis. And the variables of b/a changing from 2 to 8 is based on the statistics that the b/a in the YSH8 specimens under SEM images is mainly in this range. This is also why the SC/SM is studied in the range of 0.5–5. The elements attached to the symmetrical axis which is parallel to the Y axis are immobile in the X direction as shown in Fig. 9a. And the bottom point of the above
Fig. 7. Schematic diagram of microstructure evolution for YSH8 during annealing at 1500 °C. 21471
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Fig. 8. (a) The Fracture toughness and fracture energy of YSH8 ceramics and (b)–(f) the toughening mechanisms in YSH8.
Fig. 9. Finite element simulation of sandwich structures with various geometries. (a) Geometric model of the C/M/C structure for finite element simulation, and (b) the residual stress at the boundary of M and C phase (indicated by the red box) as a function of SC/SM when b/a = 4. (c) The comprehensive stress in M phase as a function of SC/SM for various b/a models. (d) The critical SC/SM determined by the method in Fig. 3b and the corresponding steady stress value in M phase as a function of b/a. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.) Table 3 Material properties of M phase and C phase [21]. Material
Density (g/cm3)
Young's modulus (GPa)
Poisson's ratio
Average thermal expansion coefficient (10−6/K)
M phase C phase
10.09 9.59
294 267
0.294 0.304
6.84 10.44
points is immobile in all directions. Besides, there's no relative motion between the adjacent elements at the boundary of M phase and C phase since the overlapping lines of M phase and C phase are “glued”. Material properties of M phase and C phase are listed in Table 3. Thermal
residual stress after cooling from 1500 °C (stress-free) to 25 °C during sintering was calculated by FEM analysis. By changing the length of C phase (L) while keeping the b and a of M phase constant, the stress states as a function of SC/SM are plotted in
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Fig. 9b. When b/a = 4, the compressive stress in M phase increases dramatically as SC/SM increases, but keeps constant when SC/SM > 2. Residual tensile stress is generated within the C phase, and is prone to promote the crack extension because the tensile stress enlarges the level of stress concentration at the crack tip. However, due to the existence of the residual tensile stress and the adjacent compressive region, the crack is easily deflected along the C phase where thermal tensile stress is present (Fig. 8d–f). Moreover, the FEM analysis indicates that the residual tensile stress in C phase is relatively low compared with the compressive stress in M phase when SC/SM > 1. The tensile stress in C phase generally decreases but changes very slowly when SC/SM > 2. The increasing compressive stress in M phase and the decreasing stress in C phase both help to resist the crack propagation and enhance the fracture toughness of the YSH8. Obviously when the area ratio is higher than the critical SC/SM = 1.2 value (point A) for the case b/a = 4, a steady stress value in M phase is obtained, and further increasing the SC/SM value has little effect on the fracture toughness. Fig. 9c reveals the compressive stress in M phase as a function of the SC/SM value for different b/a ranging from 2 to 8. For each sandwich structure with different b/a values, there's a critical SC/SM, above which a steady stress in M phase is obtained. Extracted from Fig. 9c and d shows the critical SC/SM value and the steady stress in M phase as a function of b/a. As the b/a value increases, the SC/SM must be enhanced to achieve the steady stress, which also increases with the b/a value. In that way, as the M phase in the sandwich structure becomes thinner, the compressive stress will be higher and contributes more to the toughening of the ceramics, but the SC/SM value must be higher than the corresponding critical value. Besides, the steady stress for b/a = 8 (~1100 MPa in compress) is 57% higher than that of b/a = 2 (~700 MPa), indicating the toughening effects of the residual stress is quite significant when the M phase is of high length to width ratio. The critical SC/SM as a function of b/a determined by the FEM results is plotted in Fig. 6a, which is lower than most of the SC/SM values in the samples, indicating that the stress state of M phase in the C/M/C sandwich structure can be simplified only as a function of the b/a value since a steady stress is reached. When the crack is going to enter M phase, the compressive stress will reduce the crack propagation driving force, thus enhances the fracture toughness. The total width (b) of the M phase within the sandwich structure declines with the sintering time due to the grain growth and the simultaneous reduction of grain boundary, resulting a reducing tendency of the toughening effect. However, the residual compressive stress at the boundary of the M phase shows first increase but then decrease as the sintering time increases as illustrated in Fig. 6. Compared with YSH8-2, YSH8-4 specimen maintains a relatively higher b/a value but lower total width of the M phase in the sandwich structure. The higher fracture toughness of YSH8-4 compared with that of YSH8-2 demonstrates that the increased compressive stress of M phase seems to be the dominant factor for the change of fracture toughness rather than the decreased total width of the M phase at the initial sintering process. But when the sintering time exceeds 4 h, the fracture toughness decreases with sintering time, resulted from the decreased compressive stress alone with the decreased total width of M phase. Furthermore, the fracture toughness of dual-phase ceramics could be tailored by controlling the preparing process not just for the YSH8 ceramics, and FEM is an effective method to roughly estimate the residual stress within the ceramics and contribute to the design of ceramics with high fracture toughness. 4. Conclusion In this work, 8% mol Y2O3 partially stabilized HfO2 (YSH8) dense sample featured with a novel C phase/M phase (lath shaped)/C phase self-toughening sandwich structure was successfully prepared by solid state sintering of nano-powders of C phase. The geometry of the sandwich structure was approximately tailored by controlling sintering time
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