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Fabrication and characterization of Sialon–Si3N4 graded nano-composite ceramic tool materials
Composites: Part B 42 (2011) 1813–1820
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Fabrication and characterization of Sialon–Si3N4 graded nano-composite ceramic tool materials G.M. Zheng ⇑, J. Zhao, Y.H. Zhou, Z.J. Gao, X.B. Cui, A.H. Li Key Laboratory of High Efficiency and Clean Mechanical Manufacture of MOE, School of Mechanical Engineering, Shandong University, 17923 Jingshi Road, Jinan 250061, PR China
a r t i c l e
i n f o
Article history: Received 13 February 2011 Received in revised form 20 June 2011 Accepted 3 July 2011 Available online 14 July 2011 Keywords: A. Ceramic–matrix composites (CMCs) A. Layered structures B. Mechanical properties B. Microstructures
a b s t r a c t Sialon–Si3N4 graded nano-composite ceramic tool materials with five-layered symmetrical distribution were fabricated by using hot pressing technique. Mechanical property tests have been conducted to determine the optimal structural parameters and sintering parameters. The residual stresses in the surface layer of the graded ceramic tool materials were calculated by the indentation method. The experiment results showed that Sialon–Si3N4 graded ceramic tool materials with a thickness ratio of 0.3, which were sintered under a pressure of 35 MPa at a sintering temperature of 1700–1750 °C for 60 min, had optimum mechanical properties. And the graded structure can induce residual compressive stresses in the material surface layer. The characterization revealed a typical duplex distribution with small b-Si3N4 grains embedded in the matrix of large b-Si3N4 grains. This duplex microstructure can contribute to the improvement of flexural strength and fracture toughness. Additionally, a mix of intergranular and transgranular fracture, crack deflection and crack bridging in the material surface layer contributed to the strengthening and toughening mechanisms for Sialon–Si3N4 graded ceramic tool materials. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction The concept of functionally graded materials (FGMs) was raised in 1984 to develop heat-shielding structure materials for future space-plane program [1,2], with the composition of a FGM varying continuously from a pure metal to a pure ceramic. FGMs have the potential for use in a wide range of engineering applications, direct metal tools for industrial use, biomaterials used in artificial human implants, drug delivery devices with release rate control, armor and armament components for defense and many more [3,4]. For these applications, control of material composition is necessary both for improving a variety of properties such as toughness and strength, and for reducing interfacial stresses between dissimilar materials [5]. The introduction of the concept of FGM into the fabrication of ceramic cutting tool materials provided a new approach to improve their thermal and mechanical properties [6,7]. Functionally graded cutting ceramics with symmetrical structure have been developed to perform high speed intermittent machining of hard-to-cut materials, with higher tool lives than those of homogeneous cutting ceramics [8].
⇑ Corresponding author. Tel./fax: +86 531 88393904. E-mail address: [email protected] (G.M. Zheng). 1359-8368/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.compositesb.2011.07.007
The intrinsic drawbacks of ceramic cutting tools, such as lower strength, lower fracture toughness and lower thermal shock resistance usually make them more susceptible to excessive chipping or fracture when machining hardened materials, leading to a short tool life. According to the research [9,10], ceramic tool materials can be strengthened and toughened by adding nano-sized particles. Conventionally, Si3N4-based and Sialonbased ceramic tool materials were strengthened and toughened by the addition of nano-particles like SiC, TiC, TiN, etc. to improve the mechanical properties [11–14]. In addition, the matrix grains of the composite were refined by adding nano-Si3N4 particles. And the addition of nano-Si3N4 particles could promote the formation of the duplex distribution characteristic. The optimum flexural strength and fracture toughness were obtained when the volume fraction ratio of nano-sized Si3N4 to microsized Si3N4 is fixed at 1:3 [15]. In the present paper, Sialon–Si3N4 graded nano-composite ceramic tool materials were synthesized by hot pressing, which were designed by adding nano-sized particles and constructing a graded structure model. Scanning electron microscopy (SEM) was used to investigate microstructures. Phase identification was carried out by X-ray diffraction (XRD). The mechanical properties of the graded ceramic tool materials were tested. The residual stresses in the surface layer of the graded ceramic tool materials were calculated by the indentation method.
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G.M. Zheng et al. / Composites: Part B 42 (2011) 1813–1820 Table 2 Composite of each layer of five-layered FGM.
2. Experimental details 2.1. Model of graded nano-composite ceramic tool materials Considering the ease of both modeling and fabricating process, the compositional distribution of five-layered graded material with symmetrical structure was modeled in the present work (Fig. 1). The compositional distribution changes along the Z-axis. The thickness of the surface layer, the second layer, the third layer and the total thickness of the material, were a, b, c and H, respectively. Thickness ratio e = a/b = b/c, a structural parameter, was defined to determine the thickness of layers. For a given ceramic tool material, the value of its compressive strength is far higher than that of its tensile strength. From this point of view, formation of residual compressive stresses in the surface layer should be taken as the aim for material design to relax stresses resulting from external loadings in the cutting process. That is to say, the thermal expansion coefficient of the surface layer should be the lowest among all the layers of the material. In virtue of the thermal mismatch effect between the matrix Si3N4 (thermal expansion coefficient amatrix = 3.2 106 K1) and the TiC0.7N0.3 (aparticle = 8.6 106 K1), the layer with the highest volume fraction of TiC0.7N0.3 was put in the middle with the compositional distribution changing from the middle layer to the two surface layers. A symmetrical composition distribution was proposed to design the graded tool material, as shown in Fig. 2 (uTiC0.7N0.3 is the volume fraction of TiC0.7N0.3). Both the two opposite surfaces of an insert made by this means could be used as rake face.
Fig. 1. The model of five-layered graded material with symmetrical structure.
Fig. 2. Symmetrical composition distribution.
Code
1st (5th) layer
2nd (4th) layer
3rd layer
GSS1 GSS2
ST10 SAAT10
ST15 ST15
ST20 ST20
Table 1 shows the composition of different composites. For the composite SAAT10, b-Sialon could be produced by chemical reaction at high temperature between the major phases Si3N4, Al2O3 and AlN [16]. Table 2 shows the composite of each layer of five-layered graded material with symmetrical structure. 2.2. Materials preparation The starting materials were a-Si3N4 powders with average grain size of approximately 0.02 lm and 0.5 lm, purity 99% (Hefei Kai Nanometer Energy and Technology Co., Ltd., China), TiC0.7N0.3 particles with average grain size of approximately 0.5 lm, purity 99% (Beijing Xingrongyuan Technology Co., Ltd., China), a-Al2O3 powders with average grain size of approximately 0.1 lm, purity 99.6% (Shanghai TeamShare Nanotechnology Co., Ltd., China) and AlN powders with average grain size of approximately 0.5 lm, purity 99% (Hefei MoK Advanced Material Technology Co., Ltd., China). a-Al2O3 powders with average grain size of approximately 0.5 lm, purity 99.9% (Zibo Xinmeiyu Alumina Co., Ltd., China) and Y2O3 (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) were used as sintering additives to promote the densification of Si3N4 ceramics during the fabricating process. The surfactant polyethylene glycol (PEG, Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) and ethanol were used as dispersant and dispersing medium, respectively. The dispersion of Al2O3 nano-particles and Si3N4 nano-particles was conducted by measuring and adjusting zeta potentials and pH values (with PHS-3C pH meter, Shanghai Kangyi Instrument Co., Ltd., China) through sedimentation experiments, with the concentration of the suspensions being 2.0 vol.%. The dispersion process was assisted by ultrasonic stirring using ultra-sonic instrument (HS-10260D, China) for 15–20 min after the PEG was added. Then the pH value of the solution was adjusted with HCl (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) or NH3 (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) solutions. The relative sedimentation volume (RSV) was measured after the suspensions were sealed in 10 ml graduated cylinder for 7 days. The results indicated that when the amount of PEG (molecular weight is 4000) added is 1.5 wt.% (relative weight ratio to the particles to be dispersed) for Al2O3 nano-particles, the RSV was the smallest at pH 9.5–10. And the RSV was the smallest at pH 9.5–10 when the amount of PEG (molecular weight is 4000) added is 0.5 wt.% for Si3N4 nano-particles. Well reagglomerated and uniform suspension of Al2O3 nanoparticles and Si3N4 nano-particles were obtained. They were then mixed with micropowders of the same composite (Table 1). The mixed slurries were ball-milled for 48 h and then dried at 100 °C in a vacuum desiccator (Model ZK-40, China). The powder mixtures were sieved through a 120 mesh sieve. The composite powders
Table 1 Composition of different composites (vol.%). Composites
Si3N4 (0.5 lm)
Si3N4 (0.02 lm)
Al2O3 (0.5 lm)
Al2O3 (0.1 lm)
AlN (0.5 lm)
TiC0.7N0.3 (0.5 lm)
Y2O3
SAAT10 ST10 ST15 ST20
53.25 61.50 57.75 54.00
17.75 20.50 19.25 18.00
0 3.2 3.2 3.2
10 0 0 0
5 0 0 0
10 10 15 20
4.0 4.8 4.8 4.8
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Fig. 5. Cross-sectional surfaces of GSS2 with a thickness ratio of 0.3 sintered under 35 MPa at 1750 °C for 60 min.
with different mixture ratios were layered into the graphite mold one layer after another with a predetermined thickness ratio, according to the material design results listed in Table 2. The specimens were then sintered by a multifunctional hot pressing sintering furnace (Model ZRC85-25T, China) in a vacuum environment (the working vacuity is 6.75 102 Pa), according to predetermined sintering process. The concrete sintering procedure was described as follows: the heating rate was 75–80 °C/min as the temperature increases from room temperature to 1200 °C, a pressure of 8 MPa was added, and
Fig. 3. Effect of thickness ratio on (a) flexural strength, (b) fracture toughness (surface layer) and (c) Vicker’s hardness (surface layer) of GSS1 and GSS2 sintered under 35 MPa at 1750 °C for 60 min.
Fig. 4. XRD patterns of the surface layer of (a) GSS1 and (b) GSS2 sintered under 35 MPa at 1750 °C for 60 min.
Fig. 6. Effect of sintering temperature on (a) flexural strength, (b) fracture toughness (surface layer) and (c) Vicker’s hardness (surface layer) of GSS1 and GSS2 sintered under 35 MPa for 60 min.
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then the heating rate was 45–50 °C/min as the temperature increases from 1200 °C to sintering temperature (1650, 1700, 1750 or 1800 °C), the pressure was added from 8 MPa to 35 MPa. The holding time at the sintering temperature was 60 (20, 40 or 80) min. The cooling rate was 25–30 °C/min as the temperature decreases from sintering temperature to 1400 °C, and then cooling with furnace to room temperature. 2.3. Characterization Each sintered compact (with 42 mm diameter) was cut and machined into five bend bars with dimensions of 3 mm 4 mm 35 mm and then ground flat to a 10 lm finish. Flexural strength was measured using a three-point bending tester (Model WDW-50E, China) with a 20 mm span width at a loading crosshead rate of 0.5 mm/min. Fracture toughness measurement was performed using indentation method (IM) [17] at the top surface of the outer layer of the graded materials. Hardness measurements were performed by placing Vicker’s indentations on the top surface of the outer layer of the graded materials. The indentation load was 196 N and the holding time was 15 s. A minimum number of 10 specimens were tested for each experimental condition. The composition of the composites was characterized by XRD (D/max-rB, Japan Rigaku). The microstructures of specimens were analyzed with SEM (JSM-6380LA, Japan). The cross-sectional surfaces of specimens were observed by a Keyence VHX-600E optical microscope (Japan). The fractured specimen surfaces for SEM observations were eroded in melting NaOH at 400 °C for 60 s to reveal the grain boundaries.
3. Results and discussion 3.1. Mechanical properties and microstructure 3.1.1. Effect of thickness ratio Graded ceramic specimens were sintered under 35 MPa pressure at 1750 °C for 60 min. Fig. 3 shows the effect of thickness ratio on flexural strength (rf), fracture toughness of the surface layer (KIC) and Vicker’s hardness of the surface layer (HV), from which it can be seen that the thickness ratio has a great effect on the mechanical properties, especially the flexural strength and the fracture toughness. The flexural strength improves upon increasing the thickness ratio from 0.2 to 0.3, but the further increase of thickness ratio results in a decrease in flexural strength (Fig. 3a). The fracture toughness (surface layer) increases as the thickness ratio increases from 0.2 to 0.4 (Fig. 3b). A drop is caused by the increases of thickness ratio from 0.4 to 0.5, but it increases again with the further increase of thickness ratio. Similarly, the Vicker’s hardness (surface layer) increases as the thickness ratio increases from 0.2 to 0.3 (Fig. 3c). A slight drop is caused by the increases of thickness ratio from 0.3 to 0.5, but it increases again with the further increase of thickness ratio. Consequently, the optimum mechanical properties of the graded ceramic materials were achieved by the thickness ratio e = 0.3. The mechanical properties of GSS1 with a thickness ratio of 0.3 are rf = 865 MPa, KIC = 9.51 MPa m1/2 and HV = 16.41 GPa, respectively. The mechanical properties of GSS2 with a thickness ratio of 0.3 are rf = 810 MPa, KIC = 9.33 MPa m1/2 and HV = 16.98 GPa, respectively.
Fig. 7. SEM micrographs of etched fracture surface of the first layer of GSS2 with a thickness ratio of 0.3, sintered under 35 MPa for at (a) 1650 °C, (b) 1700 °C, (c) 1750 °C and (d) 1800 °C for 60 min.
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The XRD patterns of the surface layer of GSS1 and GSS2 are shown in Fig. 4a and b. The major phases in the surface layer of GSS1 are b-Si3N4 and TiC0.7N0.3 with no a-Si3N4 identified (Fig. 4a), revealing a complete phase transformation from a-Si3N4 to b-Si3N4. It can be seen from Fig. 4b, chemical reaction occurred between the major phases Si3N4, Al2O3 and AlN in this hot pressing sintering process, which resulted in the formation of b-Sialon (Si4Al2O2N6) phase. Sialon ceramic material is an excellent overall performance ceramic material with very high high-temperature hardness. Wear resistance and chemical stability of Sialon ceramic material are better than those of Si3N4 ceramic material. By comparing the flexural strength and the Vicker’s hardness (surface layer) (Fig. 3a and c), it also can be found that the flexural strength of GSS2 is lower than that of GSS1, but the Vicker’s hardness (surface layer) of GSS2 is higher than that of GSS1 because of the formation of Sialon phase in the surface layer of GSS2. Fig. 5 shows the cross-sectional surfaces of GSS2 observed by a VHX-600E optical microscope. The layered architectures can be clearly seen and are reasonably uniform. The layers are all highly densified without voids and the interfaces are straight. Moreover, the thickness of each layer is accordance with the design. The formation of residual compressive stresses in the material surface layer was taken as the aim for material design (Section 2.1). Indentation tests were performed on the top surface of the outer layer to measure residual stresses. The magnitude of surface compressive stress can be estimated from the difference in indentation crack length between a stressed surface and stress-free surface [18,19]. For a stress-free material the fracture toughness can be related to the indentation load and length of the relative cracks emanating from the corners of the impressions through the following equation
K IC ¼ XPc03=2
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The flexural strength of GSS1 increases as the sintering temperature increases from 1650 °C to 1700 °C. A drop is caused by the increases of sintering temperature from 1700 °C to 1750 °C, but it increases again with the further increase of sintering temperature (Fig. 6a). The fracture toughness (surface layer) of GSS1 is a little change with the increase of sintering temperature (Fig. 6b). The Vicker’s hardness (surface layer) of GSS1 improve upon increasing the sintering temperature from 1650 °C to 1700 °C, but the further increase of sintering temperature results in a decrease in Vicker’s hardness (surface layer) (Fig. 6c). It can be seen that the highest mechanical properties of GSS1 were obtained at the sintering temperature of 1700 °C. For a sintering temperature of less than 1750 °C, the flexural strength and fracture toughness (surface layer) of GSS2 increase with temperature (Fig. 6a and b). However, a further increase results in a decrease of the flexural strength and fracture toughness (surface layer). On the other hand, the highest Vicker’s hardness (surface layer) of GSS2 was obtained at 1650 °C (Fig. 6c). SEM micrographs of etched fracture surface of the first layer of GSS2 with a thickness ratio of 0.3 sintered under 35 MPa at (a) 1650 °C, (b) 1700 °C, (c) 1750 °C and (d) 1800 °C for 60 min are shown in Fig. 7a–d. It can be seen that there are very few equiaxed b-Sialon grains in Fig. 7a and d because of sintering temperature. Compared with the composites sintered at 1650 °C and 1700 °C
ð1Þ
where KIC is toughness of the material, P is indentation load, c0 is crack length in the stress free material, and X is a dimensionless constant. X is related to the magnitude of the residual stress field associated with the indentation. For many ceramics, the constant is approximated as 0.016(E/HV)1/2 where E is the Young’s modulus of the material and HV is the hardness of the material [19,20]. Young’s modulus E and hardness HV can be measured using the experimental method. In the presence of a uniform residual stress field, Eq. (1) becomes
K IC ¼ XPc13=2 þ Y rres c1=2 1
ð2Þ
where c1 is crack length in the stressed material, Y = 1.26 is a geometrical factor of crack (for a semicircular surface crack) [19], rres is residual stress, while the other symbols have the same meaning as above. Using the value for KIC (measured using the indentation method), P, X and c1 through Eq. (2) allows us to calculate rres. The calculation results of residual stress in the material surface layer of GSS1 and GSS2 with a thickness ratio of 0.3 are 494 MPa and 442 MPa, respectively. It can be seen that residual compressive stresses were introduced to the graded material surface layer, which is accordance with the design. The formation of residual compressive stresses in the material surface layer is contribution to an improvement of cutting performance [21]. 3.1.2. Effect of sintering temperature Graded ceramic specimens with a thickness ratio of 0.3 were sintered under 35 MPa pressure at different sintering temperatures (1650 °C, 1700 °C, 1750 °C and 1800 °C) for 60 min. Fig. 6 shows the effect of sintering temperature on the mechanical properties.
Fig. 8. Effect of holding time on (a) flexural strength, (b) fracture toughness (surface layer) and (c) Vicker’s hardness (surface layer) of GSS1 and GSS2 sintered under 35 MPa at 1700 °C and 1750 °C, respectively.
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(Fig. 7a and b), more equiaxed b-Sialon grains were formed on the fracture surface of the first layer sintered at 1750 °C (Fig. 7c). This indicates that b-Sialon grains coarsened quickly at the sintering temperature of 1800 °C (Fig. 7d). Furthermore, the TiC0.7N0.3 grains also grew to agglomerations. This was detrimental to the mechanical properties of the graded nano-composite ceramic materials. Thus, the optimum sintering temperature of GSS2 was 1750 °C.
3.1.3. Effect of holding time Then GSS1 with a thickness ratio of 0.3 was sintered at sintering temperature 1700 °C. GSS2 with a thickness ratio of 0.3 was sintered at sintering temperature 1750 °C. The effects of holding time on the mechanical properties of GSS1 and GSS2 are shown in Fig. 8a–c. For a holding time of less than 60 min, the flexural strength of GSS1 and GSS2 increases with time (Fig. 8a). However, a further increase results in a decrease of the flexural strength. For a holding time is less than 40 min, the fracture toughness (surface layer) of GSS1 increases with time (Fig. 8b), and then a further increase results in a decrease of the fracture toughness (surface layer). However, the fracture toughness (surface layer) of GSS2 increases with an increase in holding time. It also can be seen that the Vicker’s hardness (surface layer) of GSS1 and GSS2 decreases as the holding time increases from 20 min to 40 min (Fig. 8c). A slight rise is caused by the increases of holding time from 40 min to 60 min, but it decreases again with the further increase of holding time. Therefore, the optimum mechanical properties of GSS1 and GSS2 were achieved by the holding time of 60 min.
SEM micrographs of etched fracture surface of the first layer of GSS1 with a thickness ratio of 0.3 sintered under 35 MPa at 1700 °C for (a) 20 min, (b) 40 min, (c) 60 min and (d) 80 min are shown in Fig. 9a–d. For a holding time of 20 min, a homogeneous microstructure with fine needle-like b-Si3N4 grains was observed in Fig. 9a. This kind of fine microstructure can reduce the mechanical properties [22]. For a holding time of 40 min and 60 min (Fig. 9b and c), a duplex distribution with small b-Si3N4 grains embedded in the matrix of large b-Si3N4 grains. The surface energy of the composite powder was increased because of the addition of nano-particles, leading to the increase in sintering driving force. Therefore the phase transformation (from a-Si3N4 to b-Si3N4), and densification of the composite were promoted through a rapid dissolution–transport–reprecipitation process. On the other hand, the difference in a-Si3N4 starting grain size led to the difference in b-Si3N4 grains size which can result in the formation of the interlocked duplex microstructure. The interlocked duplex microstructure can contribute to the improvement of flexural strength (Fig. 8a). In the microstructure after a holding time of 80 min, a few of b-Si3N4 grains had been coarsened, which resulted in a slight deterioration of the mechanical properties compared with that for a holding time of 60 min. Thus the holding time of GSS1 should not exceed 60 min. Consequently, the optimum mechanical properties of GSS1 with a thickness ratio of 0.3 were achieved by the sintering temperature at 1700 °C for 60 min. The mechanical properties of GSS1 are rf = 980 MPa, KIC = 9.54 MPa m1/2 and HV = 16.91 GPa, respectively. On the other hand, the optimum mechanical properties of GSS2 with a thickness ratio of 0.3 were achieved by the sintering
Fig. 9. SEM micrographs of etched fracture surface of the first layer of GSS1 with a thickness ratio of 0.3, sintered under 35 MPa at 1700 °C for (a) 20, (b) 40, (c) 60 and (d) 80 min.
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temperature at 1750 °C for 60 min. The mechanical properties of GSS2 are rf = 810 MPa, KIC = 9.33 MPa m1/2 and HV = 16.98 GPa, respectively. 3.2. Toughening and strengthening mechanisms Fig. 10 shows SEM micrographs of crack extension paths generated by Vicker’s indentation of the surface layer for (a) GSS1 with a thickness ratio of 0.3, sintered under 35 MPa at 1700 °C for 60 min and (b) GSS2 with a thickness ratio of 0.3, sintered under 35 MPa at 1750 °C for 60 min. In virtue of the thermal mismatch effect between the matrix Si3N4 (thermal expansion coefficient amatrix = 3.2 106 K1) and the TiC0.7N0.3 (aparticle = 8.6 106 K1) along with the presence of TiC0.7N0.3 particles, radial residual tensile stresses were formed in the matix around TiC0.7N0.3 particles during the fabricating process (cooling from the sintering temperature to room temperature), which resulted in transgranular fracture. Therefore, it was seen from Fig. 10 that transgranular fracture is the main way of crack extension. At the same time, intergranular fracture (region A) was also observed in Fig. 10a. So it is a mix of intergranular and transgranular fracture. Crack deflection can be observed clearly in the Vicker’s indentation crack extension path of the surface layer of GSS1 and GSS2. The crack propagated from the extension direction with a greater deflection angle. This made
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the newly ruptured surface areas become larger, and a greater amount of rupture energy was consumed. Hence, the fracture toughness increased. Additionally, crack bridging (region B) was observed, which provides a restraining force for the crack growth. Crack deflection and crack bridging can absorb additional amounts of fracture energy, which could significantly improve the fracture toughness of the composite. 4. Conclusions This work attempts to present a model for designing Sialon– Si3N4 graded nano-composite ceramic tool materials with symmetrical distribution, and a systematic experimental has been conducted to determine the optimal structural parameters and sintering parameters of the Sialon–Si3N4 graded ceramic tool materials. The following concluding remarks can be underlined: (1) GSS1 materials with optimum mechanical properties were sintered under a pressure of 35 MPa at 1700 °C for 60 min, while the thickness ratio e = 0.3. The optimum mechanical properties are a flexural strength of 980 MPa, a fracture toughness (surface layer) of 9.54 MPa m1/2 and a Vicker’s hardness (surface layer) of 16.91 GPa, which meet the requirements for cutting tool materials. (2) GSS2 materials with optimum mechanical properties were sintered under a pressure of 35 MPa at 1750 °C for 60 min, while the thickness ratio e = 0.3. The optimum mechanical properties are a flexural strength of 810 MPa, a fracture toughness (surface layer) of 9.33 MPa m1/2 and a Vicker’s hardness (surface layer) of 16.98 GPa, which meet the requirements for cutting tool materials. (3) The graded structure in Sialon–Si3N4 graded ceramic tool materials can induce residual stresses during fabrication. These residual stresses are compressive in the material surface layer. (4) The characterization revealed a typical duplex distribution, with small b-Si3N4 grains embedded in the matrix of large b-Si3N4 grains. Several strengthening and toughening mechanisms, such as a mix of intergranular and transgranular fracture, crack deflection and crack bridging, acted simultaneously in the surface layer of Sialon–Si3N4 graded ceramic tool materials. These mechanisms had a significant influence on the mechanical properties.
Acknowledgements This research is supported by the National Basic Research Program of China (2009CB724402), the National Natural Science Foundation of China (50875156) and the Specialized Research Fund for the Doctoral Program of Higher Education (20090131110030). References
Fig. 10. Vicker’s indentation crack extension paths of the surface layer of (a) GSS1 with a thickness ratio of 0.3, sintered under 35 MPa at 1700 °C for 60 min and (b) GSS2 with a thickness ratio of 0.3, sintered under 35 MPa at 1750 °C for 60 min.
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Composites: Part B journal homepage: www.elsevier.com/locate/compositesb
Fabrication and characterization of Sialon–Si3N4 graded nano-composite ceramic tool materials G.M. Zheng ⇑, J. Zhao, Y.H. Zhou, Z.J. Gao, X.B. Cui, A.H. Li Key Laboratory of High Efficiency and Clean Mechanical Manufacture of MOE, School of Mechanical Engineering, Shandong University, 17923 Jingshi Road, Jinan 250061, PR China
a r t i c l e
i n f o
Article history: Received 13 February 2011 Received in revised form 20 June 2011 Accepted 3 July 2011 Available online 14 July 2011 Keywords: A. Ceramic–matrix composites (CMCs) A. Layered structures B. Mechanical properties B. Microstructures
a b s t r a c t Sialon–Si3N4 graded nano-composite ceramic tool materials with five-layered symmetrical distribution were fabricated by using hot pressing technique. Mechanical property tests have been conducted to determine the optimal structural parameters and sintering parameters. The residual stresses in the surface layer of the graded ceramic tool materials were calculated by the indentation method. The experiment results showed that Sialon–Si3N4 graded ceramic tool materials with a thickness ratio of 0.3, which were sintered under a pressure of 35 MPa at a sintering temperature of 1700–1750 °C for 60 min, had optimum mechanical properties. And the graded structure can induce residual compressive stresses in the material surface layer. The characterization revealed a typical duplex distribution with small b-Si3N4 grains embedded in the matrix of large b-Si3N4 grains. This duplex microstructure can contribute to the improvement of flexural strength and fracture toughness. Additionally, a mix of intergranular and transgranular fracture, crack deflection and crack bridging in the material surface layer contributed to the strengthening and toughening mechanisms for Sialon–Si3N4 graded ceramic tool materials. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction The concept of functionally graded materials (FGMs) was raised in 1984 to develop heat-shielding structure materials for future space-plane program [1,2], with the composition of a FGM varying continuously from a pure metal to a pure ceramic. FGMs have the potential for use in a wide range of engineering applications, direct metal tools for industrial use, biomaterials used in artificial human implants, drug delivery devices with release rate control, armor and armament components for defense and many more [3,4]. For these applications, control of material composition is necessary both for improving a variety of properties such as toughness and strength, and for reducing interfacial stresses between dissimilar materials [5]. The introduction of the concept of FGM into the fabrication of ceramic cutting tool materials provided a new approach to improve their thermal and mechanical properties [6,7]. Functionally graded cutting ceramics with symmetrical structure have been developed to perform high speed intermittent machining of hard-to-cut materials, with higher tool lives than those of homogeneous cutting ceramics [8].
⇑ Corresponding author. Tel./fax: +86 531 88393904. E-mail address: [email protected] (G.M. Zheng). 1359-8368/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.compositesb.2011.07.007
The intrinsic drawbacks of ceramic cutting tools, such as lower strength, lower fracture toughness and lower thermal shock resistance usually make them more susceptible to excessive chipping or fracture when machining hardened materials, leading to a short tool life. According to the research [9,10], ceramic tool materials can be strengthened and toughened by adding nano-sized particles. Conventionally, Si3N4-based and Sialonbased ceramic tool materials were strengthened and toughened by the addition of nano-particles like SiC, TiC, TiN, etc. to improve the mechanical properties [11–14]. In addition, the matrix grains of the composite were refined by adding nano-Si3N4 particles. And the addition of nano-Si3N4 particles could promote the formation of the duplex distribution characteristic. The optimum flexural strength and fracture toughness were obtained when the volume fraction ratio of nano-sized Si3N4 to microsized Si3N4 is fixed at 1:3 [15]. In the present paper, Sialon–Si3N4 graded nano-composite ceramic tool materials were synthesized by hot pressing, which were designed by adding nano-sized particles and constructing a graded structure model. Scanning electron microscopy (SEM) was used to investigate microstructures. Phase identification was carried out by X-ray diffraction (XRD). The mechanical properties of the graded ceramic tool materials were tested. The residual stresses in the surface layer of the graded ceramic tool materials were calculated by the indentation method.
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G.M. Zheng et al. / Composites: Part B 42 (2011) 1813–1820 Table 2 Composite of each layer of five-layered FGM.
2. Experimental details 2.1. Model of graded nano-composite ceramic tool materials Considering the ease of both modeling and fabricating process, the compositional distribution of five-layered graded material with symmetrical structure was modeled in the present work (Fig. 1). The compositional distribution changes along the Z-axis. The thickness of the surface layer, the second layer, the third layer and the total thickness of the material, were a, b, c and H, respectively. Thickness ratio e = a/b = b/c, a structural parameter, was defined to determine the thickness of layers. For a given ceramic tool material, the value of its compressive strength is far higher than that of its tensile strength. From this point of view, formation of residual compressive stresses in the surface layer should be taken as the aim for material design to relax stresses resulting from external loadings in the cutting process. That is to say, the thermal expansion coefficient of the surface layer should be the lowest among all the layers of the material. In virtue of the thermal mismatch effect between the matrix Si3N4 (thermal expansion coefficient amatrix = 3.2 106 K1) and the TiC0.7N0.3 (aparticle = 8.6 106 K1), the layer with the highest volume fraction of TiC0.7N0.3 was put in the middle with the compositional distribution changing from the middle layer to the two surface layers. A symmetrical composition distribution was proposed to design the graded tool material, as shown in Fig. 2 (uTiC0.7N0.3 is the volume fraction of TiC0.7N0.3). Both the two opposite surfaces of an insert made by this means could be used as rake face.
Fig. 1. The model of five-layered graded material with symmetrical structure.
Fig. 2. Symmetrical composition distribution.
Code
1st (5th) layer
2nd (4th) layer
3rd layer
GSS1 GSS2
ST10 SAAT10
ST15 ST15
ST20 ST20
Table 1 shows the composition of different composites. For the composite SAAT10, b-Sialon could be produced by chemical reaction at high temperature between the major phases Si3N4, Al2O3 and AlN [16]. Table 2 shows the composite of each layer of five-layered graded material with symmetrical structure. 2.2. Materials preparation The starting materials were a-Si3N4 powders with average grain size of approximately 0.02 lm and 0.5 lm, purity 99% (Hefei Kai Nanometer Energy and Technology Co., Ltd., China), TiC0.7N0.3 particles with average grain size of approximately 0.5 lm, purity 99% (Beijing Xingrongyuan Technology Co., Ltd., China), a-Al2O3 powders with average grain size of approximately 0.1 lm, purity 99.6% (Shanghai TeamShare Nanotechnology Co., Ltd., China) and AlN powders with average grain size of approximately 0.5 lm, purity 99% (Hefei MoK Advanced Material Technology Co., Ltd., China). a-Al2O3 powders with average grain size of approximately 0.5 lm, purity 99.9% (Zibo Xinmeiyu Alumina Co., Ltd., China) and Y2O3 (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) were used as sintering additives to promote the densification of Si3N4 ceramics during the fabricating process. The surfactant polyethylene glycol (PEG, Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) and ethanol were used as dispersant and dispersing medium, respectively. The dispersion of Al2O3 nano-particles and Si3N4 nano-particles was conducted by measuring and adjusting zeta potentials and pH values (with PHS-3C pH meter, Shanghai Kangyi Instrument Co., Ltd., China) through sedimentation experiments, with the concentration of the suspensions being 2.0 vol.%. The dispersion process was assisted by ultrasonic stirring using ultra-sonic instrument (HS-10260D, China) for 15–20 min after the PEG was added. Then the pH value of the solution was adjusted with HCl (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) or NH3 (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) solutions. The relative sedimentation volume (RSV) was measured after the suspensions were sealed in 10 ml graduated cylinder for 7 days. The results indicated that when the amount of PEG (molecular weight is 4000) added is 1.5 wt.% (relative weight ratio to the particles to be dispersed) for Al2O3 nano-particles, the RSV was the smallest at pH 9.5–10. And the RSV was the smallest at pH 9.5–10 when the amount of PEG (molecular weight is 4000) added is 0.5 wt.% for Si3N4 nano-particles. Well reagglomerated and uniform suspension of Al2O3 nanoparticles and Si3N4 nano-particles were obtained. They were then mixed with micropowders of the same composite (Table 1). The mixed slurries were ball-milled for 48 h and then dried at 100 °C in a vacuum desiccator (Model ZK-40, China). The powder mixtures were sieved through a 120 mesh sieve. The composite powders
Table 1 Composition of different composites (vol.%). Composites
Si3N4 (0.5 lm)
Si3N4 (0.02 lm)
Al2O3 (0.5 lm)
Al2O3 (0.1 lm)
AlN (0.5 lm)
TiC0.7N0.3 (0.5 lm)
Y2O3
SAAT10 ST10 ST15 ST20
53.25 61.50 57.75 54.00
17.75 20.50 19.25 18.00
0 3.2 3.2 3.2
10 0 0 0
5 0 0 0
10 10 15 20
4.0 4.8 4.8 4.8
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Fig. 5. Cross-sectional surfaces of GSS2 with a thickness ratio of 0.3 sintered under 35 MPa at 1750 °C for 60 min.
with different mixture ratios were layered into the graphite mold one layer after another with a predetermined thickness ratio, according to the material design results listed in Table 2. The specimens were then sintered by a multifunctional hot pressing sintering furnace (Model ZRC85-25T, China) in a vacuum environment (the working vacuity is 6.75 102 Pa), according to predetermined sintering process. The concrete sintering procedure was described as follows: the heating rate was 75–80 °C/min as the temperature increases from room temperature to 1200 °C, a pressure of 8 MPa was added, and
Fig. 3. Effect of thickness ratio on (a) flexural strength, (b) fracture toughness (surface layer) and (c) Vicker’s hardness (surface layer) of GSS1 and GSS2 sintered under 35 MPa at 1750 °C for 60 min.
Fig. 4. XRD patterns of the surface layer of (a) GSS1 and (b) GSS2 sintered under 35 MPa at 1750 °C for 60 min.
Fig. 6. Effect of sintering temperature on (a) flexural strength, (b) fracture toughness (surface layer) and (c) Vicker’s hardness (surface layer) of GSS1 and GSS2 sintered under 35 MPa for 60 min.
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then the heating rate was 45–50 °C/min as the temperature increases from 1200 °C to sintering temperature (1650, 1700, 1750 or 1800 °C), the pressure was added from 8 MPa to 35 MPa. The holding time at the sintering temperature was 60 (20, 40 or 80) min. The cooling rate was 25–30 °C/min as the temperature decreases from sintering temperature to 1400 °C, and then cooling with furnace to room temperature. 2.3. Characterization Each sintered compact (with 42 mm diameter) was cut and machined into five bend bars with dimensions of 3 mm 4 mm 35 mm and then ground flat to a 10 lm finish. Flexural strength was measured using a three-point bending tester (Model WDW-50E, China) with a 20 mm span width at a loading crosshead rate of 0.5 mm/min. Fracture toughness measurement was performed using indentation method (IM) [17] at the top surface of the outer layer of the graded materials. Hardness measurements were performed by placing Vicker’s indentations on the top surface of the outer layer of the graded materials. The indentation load was 196 N and the holding time was 15 s. A minimum number of 10 specimens were tested for each experimental condition. The composition of the composites was characterized by XRD (D/max-rB, Japan Rigaku). The microstructures of specimens were analyzed with SEM (JSM-6380LA, Japan). The cross-sectional surfaces of specimens were observed by a Keyence VHX-600E optical microscope (Japan). The fractured specimen surfaces for SEM observations were eroded in melting NaOH at 400 °C for 60 s to reveal the grain boundaries.
3. Results and discussion 3.1. Mechanical properties and microstructure 3.1.1. Effect of thickness ratio Graded ceramic specimens were sintered under 35 MPa pressure at 1750 °C for 60 min. Fig. 3 shows the effect of thickness ratio on flexural strength (rf), fracture toughness of the surface layer (KIC) and Vicker’s hardness of the surface layer (HV), from which it can be seen that the thickness ratio has a great effect on the mechanical properties, especially the flexural strength and the fracture toughness. The flexural strength improves upon increasing the thickness ratio from 0.2 to 0.3, but the further increase of thickness ratio results in a decrease in flexural strength (Fig. 3a). The fracture toughness (surface layer) increases as the thickness ratio increases from 0.2 to 0.4 (Fig. 3b). A drop is caused by the increases of thickness ratio from 0.4 to 0.5, but it increases again with the further increase of thickness ratio. Similarly, the Vicker’s hardness (surface layer) increases as the thickness ratio increases from 0.2 to 0.3 (Fig. 3c). A slight drop is caused by the increases of thickness ratio from 0.3 to 0.5, but it increases again with the further increase of thickness ratio. Consequently, the optimum mechanical properties of the graded ceramic materials were achieved by the thickness ratio e = 0.3. The mechanical properties of GSS1 with a thickness ratio of 0.3 are rf = 865 MPa, KIC = 9.51 MPa m1/2 and HV = 16.41 GPa, respectively. The mechanical properties of GSS2 with a thickness ratio of 0.3 are rf = 810 MPa, KIC = 9.33 MPa m1/2 and HV = 16.98 GPa, respectively.
Fig. 7. SEM micrographs of etched fracture surface of the first layer of GSS2 with a thickness ratio of 0.3, sintered under 35 MPa for at (a) 1650 °C, (b) 1700 °C, (c) 1750 °C and (d) 1800 °C for 60 min.
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The XRD patterns of the surface layer of GSS1 and GSS2 are shown in Fig. 4a and b. The major phases in the surface layer of GSS1 are b-Si3N4 and TiC0.7N0.3 with no a-Si3N4 identified (Fig. 4a), revealing a complete phase transformation from a-Si3N4 to b-Si3N4. It can be seen from Fig. 4b, chemical reaction occurred between the major phases Si3N4, Al2O3 and AlN in this hot pressing sintering process, which resulted in the formation of b-Sialon (Si4Al2O2N6) phase. Sialon ceramic material is an excellent overall performance ceramic material with very high high-temperature hardness. Wear resistance and chemical stability of Sialon ceramic material are better than those of Si3N4 ceramic material. By comparing the flexural strength and the Vicker’s hardness (surface layer) (Fig. 3a and c), it also can be found that the flexural strength of GSS2 is lower than that of GSS1, but the Vicker’s hardness (surface layer) of GSS2 is higher than that of GSS1 because of the formation of Sialon phase in the surface layer of GSS2. Fig. 5 shows the cross-sectional surfaces of GSS2 observed by a VHX-600E optical microscope. The layered architectures can be clearly seen and are reasonably uniform. The layers are all highly densified without voids and the interfaces are straight. Moreover, the thickness of each layer is accordance with the design. The formation of residual compressive stresses in the material surface layer was taken as the aim for material design (Section 2.1). Indentation tests were performed on the top surface of the outer layer to measure residual stresses. The magnitude of surface compressive stress can be estimated from the difference in indentation crack length between a stressed surface and stress-free surface [18,19]. For a stress-free material the fracture toughness can be related to the indentation load and length of the relative cracks emanating from the corners of the impressions through the following equation
K IC ¼ XPc03=2
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The flexural strength of GSS1 increases as the sintering temperature increases from 1650 °C to 1700 °C. A drop is caused by the increases of sintering temperature from 1700 °C to 1750 °C, but it increases again with the further increase of sintering temperature (Fig. 6a). The fracture toughness (surface layer) of GSS1 is a little change with the increase of sintering temperature (Fig. 6b). The Vicker’s hardness (surface layer) of GSS1 improve upon increasing the sintering temperature from 1650 °C to 1700 °C, but the further increase of sintering temperature results in a decrease in Vicker’s hardness (surface layer) (Fig. 6c). It can be seen that the highest mechanical properties of GSS1 were obtained at the sintering temperature of 1700 °C. For a sintering temperature of less than 1750 °C, the flexural strength and fracture toughness (surface layer) of GSS2 increase with temperature (Fig. 6a and b). However, a further increase results in a decrease of the flexural strength and fracture toughness (surface layer). On the other hand, the highest Vicker’s hardness (surface layer) of GSS2 was obtained at 1650 °C (Fig. 6c). SEM micrographs of etched fracture surface of the first layer of GSS2 with a thickness ratio of 0.3 sintered under 35 MPa at (a) 1650 °C, (b) 1700 °C, (c) 1750 °C and (d) 1800 °C for 60 min are shown in Fig. 7a–d. It can be seen that there are very few equiaxed b-Sialon grains in Fig. 7a and d because of sintering temperature. Compared with the composites sintered at 1650 °C and 1700 °C
ð1Þ
where KIC is toughness of the material, P is indentation load, c0 is crack length in the stress free material, and X is a dimensionless constant. X is related to the magnitude of the residual stress field associated with the indentation. For many ceramics, the constant is approximated as 0.016(E/HV)1/2 where E is the Young’s modulus of the material and HV is the hardness of the material [19,20]. Young’s modulus E and hardness HV can be measured using the experimental method. In the presence of a uniform residual stress field, Eq. (1) becomes
K IC ¼ XPc13=2 þ Y rres c1=2 1
ð2Þ
where c1 is crack length in the stressed material, Y = 1.26 is a geometrical factor of crack (for a semicircular surface crack) [19], rres is residual stress, while the other symbols have the same meaning as above. Using the value for KIC (measured using the indentation method), P, X and c1 through Eq. (2) allows us to calculate rres. The calculation results of residual stress in the material surface layer of GSS1 and GSS2 with a thickness ratio of 0.3 are 494 MPa and 442 MPa, respectively. It can be seen that residual compressive stresses were introduced to the graded material surface layer, which is accordance with the design. The formation of residual compressive stresses in the material surface layer is contribution to an improvement of cutting performance [21]. 3.1.2. Effect of sintering temperature Graded ceramic specimens with a thickness ratio of 0.3 were sintered under 35 MPa pressure at different sintering temperatures (1650 °C, 1700 °C, 1750 °C and 1800 °C) for 60 min. Fig. 6 shows the effect of sintering temperature on the mechanical properties.
Fig. 8. Effect of holding time on (a) flexural strength, (b) fracture toughness (surface layer) and (c) Vicker’s hardness (surface layer) of GSS1 and GSS2 sintered under 35 MPa at 1700 °C and 1750 °C, respectively.
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(Fig. 7a and b), more equiaxed b-Sialon grains were formed on the fracture surface of the first layer sintered at 1750 °C (Fig. 7c). This indicates that b-Sialon grains coarsened quickly at the sintering temperature of 1800 °C (Fig. 7d). Furthermore, the TiC0.7N0.3 grains also grew to agglomerations. This was detrimental to the mechanical properties of the graded nano-composite ceramic materials. Thus, the optimum sintering temperature of GSS2 was 1750 °C.
3.1.3. Effect of holding time Then GSS1 with a thickness ratio of 0.3 was sintered at sintering temperature 1700 °C. GSS2 with a thickness ratio of 0.3 was sintered at sintering temperature 1750 °C. The effects of holding time on the mechanical properties of GSS1 and GSS2 are shown in Fig. 8a–c. For a holding time of less than 60 min, the flexural strength of GSS1 and GSS2 increases with time (Fig. 8a). However, a further increase results in a decrease of the flexural strength. For a holding time is less than 40 min, the fracture toughness (surface layer) of GSS1 increases with time (Fig. 8b), and then a further increase results in a decrease of the fracture toughness (surface layer). However, the fracture toughness (surface layer) of GSS2 increases with an increase in holding time. It also can be seen that the Vicker’s hardness (surface layer) of GSS1 and GSS2 decreases as the holding time increases from 20 min to 40 min (Fig. 8c). A slight rise is caused by the increases of holding time from 40 min to 60 min, but it decreases again with the further increase of holding time. Therefore, the optimum mechanical properties of GSS1 and GSS2 were achieved by the holding time of 60 min.
SEM micrographs of etched fracture surface of the first layer of GSS1 with a thickness ratio of 0.3 sintered under 35 MPa at 1700 °C for (a) 20 min, (b) 40 min, (c) 60 min and (d) 80 min are shown in Fig. 9a–d. For a holding time of 20 min, a homogeneous microstructure with fine needle-like b-Si3N4 grains was observed in Fig. 9a. This kind of fine microstructure can reduce the mechanical properties [22]. For a holding time of 40 min and 60 min (Fig. 9b and c), a duplex distribution with small b-Si3N4 grains embedded in the matrix of large b-Si3N4 grains. The surface energy of the composite powder was increased because of the addition of nano-particles, leading to the increase in sintering driving force. Therefore the phase transformation (from a-Si3N4 to b-Si3N4), and densification of the composite were promoted through a rapid dissolution–transport–reprecipitation process. On the other hand, the difference in a-Si3N4 starting grain size led to the difference in b-Si3N4 grains size which can result in the formation of the interlocked duplex microstructure. The interlocked duplex microstructure can contribute to the improvement of flexural strength (Fig. 8a). In the microstructure after a holding time of 80 min, a few of b-Si3N4 grains had been coarsened, which resulted in a slight deterioration of the mechanical properties compared with that for a holding time of 60 min. Thus the holding time of GSS1 should not exceed 60 min. Consequently, the optimum mechanical properties of GSS1 with a thickness ratio of 0.3 were achieved by the sintering temperature at 1700 °C for 60 min. The mechanical properties of GSS1 are rf = 980 MPa, KIC = 9.54 MPa m1/2 and HV = 16.91 GPa, respectively. On the other hand, the optimum mechanical properties of GSS2 with a thickness ratio of 0.3 were achieved by the sintering
Fig. 9. SEM micrographs of etched fracture surface of the first layer of GSS1 with a thickness ratio of 0.3, sintered under 35 MPa at 1700 °C for (a) 20, (b) 40, (c) 60 and (d) 80 min.
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temperature at 1750 °C for 60 min. The mechanical properties of GSS2 are rf = 810 MPa, KIC = 9.33 MPa m1/2 and HV = 16.98 GPa, respectively. 3.2. Toughening and strengthening mechanisms Fig. 10 shows SEM micrographs of crack extension paths generated by Vicker’s indentation of the surface layer for (a) GSS1 with a thickness ratio of 0.3, sintered under 35 MPa at 1700 °C for 60 min and (b) GSS2 with a thickness ratio of 0.3, sintered under 35 MPa at 1750 °C for 60 min. In virtue of the thermal mismatch effect between the matrix Si3N4 (thermal expansion coefficient amatrix = 3.2 106 K1) and the TiC0.7N0.3 (aparticle = 8.6 106 K1) along with the presence of TiC0.7N0.3 particles, radial residual tensile stresses were formed in the matix around TiC0.7N0.3 particles during the fabricating process (cooling from the sintering temperature to room temperature), which resulted in transgranular fracture. Therefore, it was seen from Fig. 10 that transgranular fracture is the main way of crack extension. At the same time, intergranular fracture (region A) was also observed in Fig. 10a. So it is a mix of intergranular and transgranular fracture. Crack deflection can be observed clearly in the Vicker’s indentation crack extension path of the surface layer of GSS1 and GSS2. The crack propagated from the extension direction with a greater deflection angle. This made
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the newly ruptured surface areas become larger, and a greater amount of rupture energy was consumed. Hence, the fracture toughness increased. Additionally, crack bridging (region B) was observed, which provides a restraining force for the crack growth. Crack deflection and crack bridging can absorb additional amounts of fracture energy, which could significantly improve the fracture toughness of the composite. 4. Conclusions This work attempts to present a model for designing Sialon– Si3N4 graded nano-composite ceramic tool materials with symmetrical distribution, and a systematic experimental has been conducted to determine the optimal structural parameters and sintering parameters of the Sialon–Si3N4 graded ceramic tool materials. The following concluding remarks can be underlined: (1) GSS1 materials with optimum mechanical properties were sintered under a pressure of 35 MPa at 1700 °C for 60 min, while the thickness ratio e = 0.3. The optimum mechanical properties are a flexural strength of 980 MPa, a fracture toughness (surface layer) of 9.54 MPa m1/2 and a Vicker’s hardness (surface layer) of 16.91 GPa, which meet the requirements for cutting tool materials. (2) GSS2 materials with optimum mechanical properties were sintered under a pressure of 35 MPa at 1750 °C for 60 min, while the thickness ratio e = 0.3. The optimum mechanical properties are a flexural strength of 810 MPa, a fracture toughness (surface layer) of 9.33 MPa m1/2 and a Vicker’s hardness (surface layer) of 16.98 GPa, which meet the requirements for cutting tool materials. (3) The graded structure in Sialon–Si3N4 graded ceramic tool materials can induce residual stresses during fabrication. These residual stresses are compressive in the material surface layer. (4) The characterization revealed a typical duplex distribution, with small b-Si3N4 grains embedded in the matrix of large b-Si3N4 grains. Several strengthening and toughening mechanisms, such as a mix of intergranular and transgranular fracture, crack deflection and crack bridging, acted simultaneously in the surface layer of Sialon–Si3N4 graded ceramic tool materials. These mechanisms had a significant influence on the mechanical properties.
Acknowledgements This research is supported by the National Basic Research Program of China (2009CB724402), the National Natural Science Foundation of China (50875156) and the Specialized Research Fund for the Doctoral Program of Higher Education (20090131110030). References
Fig. 10. Vicker’s indentation crack extension paths of the surface layer of (a) GSS1 with a thickness ratio of 0.3, sintered under 35 MPa at 1700 °C for 60 min and (b) GSS2 with a thickness ratio of 0.3, sintered under 35 MPa at 1750 °C for 60 min.
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