Effect of the nature of the intergranular phase on sliding-wear resistance of liquid-phase-sintered α-SiC

Effect of the nature of the intergranular phase on sliding-wear resistance of liquid-phase-sintered α-SiC

Scripta Materialia 57 (2007) 505–508 www.elsevier.com/locate/scriptamat Effect of the nature of the intergranular phase on sliding-wear resistance of ...

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Scripta Materialia 57 (2007) 505–508 www.elsevier.com/locate/scriptamat

Effect of the nature of the intergranular phase on sliding-wear resistance of liquid-phase-sintered a-SiC Oscar Borrero-Lo´pez,a Angel L. Ortiz,a,* Fernando Guiberteaua and Nitin P. Padtureb a

Departamento de Ingenierı´a Meca´nica, Energe´tica y de los Materiales, Escuela de Ingenierı´as Industriales, Universidad de Extremadura, 06071 Badajoz, Spain b Department of Materials Science and Engineering, The Ohio State University, Columbus, OH 43210, USA Received 9 May 2007; accepted 20 May 2007 Available online 20 June 2007

The effect of the nature of the intergranular phase on sliding-wear resistance of equiaxed-grained, liquid-phase-sintered (LPS) a-SiC ceramics was studied. The nature of the intergranular phase in LPS a-SiC ceramics was modified using Ar or N2 sintering atmospheres, while maintaining the same grain size and morphology. Sintering under N2 leads to the incorporation of nitrogen in the intergranular phase, and attendant solid-solution hardening of that phase. The solid-solution hardened intergranular phase results in LPS SiC ceramics with improved sliding-wear resistance.  2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: SiC; Sliding-wear; Liquid-phase-sintering; Microstructure

Owing to their high hardness and high stiffness, silicon carbide (SiC) ceramics are often used in contact-mechanical and tribological applications, such as bearings, wear parts, valves and seals [1–3]. Also, the excellent chemical inertness and thermal stability of SiC ceramics make them particularly suitable for use in harsh environments [1,4–6]. Fully dense SiC ceramics for such applications can be obtained at low cost by using liquid-phase sintering, aided by oxide additives. Liquid-phase-sintered (LPS) SiC ceramics are generally fabricated at lower temperatures, relative to solid-state sintered SiC [7], in an inert atmosphere of pure Ar or N2 gas. The resulting LPS SiC ceramics typically have a two-phase microstructure consisting of a distribution of SiC grains in a continuous oxide-phase matrix [8]. Sliding-wear behavior of LPS SiC ceramics has been studied recently [9–13], and it is found to be consistent with what is commonly observed in polycrystalline ceramics [14–17]: initial mild wear controlled by dislocation plasticity, followed by severe fracture-controlled wear and grain pullout, with a well-defined transition. Previous studies of sliding-wear in LPS SiC ceramics have focused on microstructural features such as size and shape of the SiC grains, and the content of the inter* Corresponding author. Tel.: +34 924 289600x6785; fax: +34 924 289601; e-mail: [email protected]

granular phase [9–13]. This has resulted in three main strategies for the microstructural design of sliding-wear resistant LPS SiC: (i) grain refinement [9,13], (ii) grain elongation [10,11,13] and (iii) reduction in intergranular phase content [9,13]. However, so far little attention has been paid to the nature of the intergranular phase. Here we have studied the sliding-wear resistance of two sets of LPS a-SiC ceramics which have identical grain size, grain morphology and intergranular phase content, but one has a harder intergranular phase than the other. Intergranular phase hardening was accomplished by sintering the LPS a-SiC in N2 atmosphere, instead of the commonly used Ar atmosphere, to incorporate nitrogen in the intergranular phase [18]. We show that an increase in the hardness of the intergranular phase results in improved sliding-wear resistance in LPS a-SiC ceramics. Starting a-SiC powder (UF-15, H.C. Starck., Goslar, Germany) with 4.29 wt.% Al2O3 (AKP-30, Sumitomo Chemical Company, New York, NY) and 5.71 wt.% Y2O3 (Fine Grade, H.C. Starck, Goslar, Germany) as additives were ball-milled for 24 h in ethanol using ZrO2 balls. This batch composition is expected to yield LPS a-SiC ceramics with 7.3 vol.% YAG after sintering. The slurry was dried, and the resulting powder deagglomerated and sieved. Compacts were made by uniaxial pressing (C, Carver Inc., Wabash, IN) at 50 MPa, followed by isostatic pressing (CP360, AIP, Columbus,

1359-6462/$ - see front matter  2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2007.05.021

O. Borrero-Lo´pez et al. / Scripta Materialia 57 (2007) 505–508

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OH) at 350 MPa. Pressureless sintering was performed (1000-3560-FP20, Thermal Technology Inc., Santa Rosa, CA) at 1950 C for 2 h in Ar atmosphere, or for 7 h in N2 atmosphere. The difference in the sintering durations gives LPS SiC ceramics with the same grain size and grain morphology [18]. Details of this processing procedure and of the microstructural evolution in LPS SiC sintered in Ar and N2 atmospheres are given elsewhere [8,18,19]. The resulting ceramics, which will be referred to as SiC–Ar and SiC–N2, were cleaned and surface layers were ground off. In order to measure the hardness of the intergranular phase, with and without nitrogen incorporation, LPS a-SiC ceramics with higher YAG content (23.2 vol.%) were also prepared using the above procedure. This increase in the YAG content has no effect on the hardness of the intergranular phase, but it allows room for making hardness indentations entirely within the intergranular phase. Densities of the sintered ceramics were measured using the Archimedes method, with distilled water as the immersion medium. Theoretical density of the LPS SiC ceramic was calculated using density values of 3.213 and 4.544 g cm 3 for SiC and YAG, respectively. Cross-sections of the SiC–Ar and SiC–N2 ceramics were polished to a 1 lm finish using routine diamondpaste polishing. The polished sections were plasmaetched (PT 1750, Fissions Instruments, East Sussex, UK), and were then observed in a scanning electron microscope (SEM) (S-3600N, Hitachi, Japan). Grain morphology analysis was performed from the SEM micrographs by image analysis on at least 300 grains per ceramic. The nitrogen content in SiC–N2 was measured using the inert gas (helium) fusion method (TNT-414, Leco Corporation, St. Joseph, MI). Sliding-wear testing was performed on polished LPS a-SiC specimens using a high-performance multi-specimen tribometer (Falex, Faville-Le Vally Corp., Sugar Grove, IL) configured in the ball-on-three-disks geometry, as described elsewhere [9]. In this testing procedure a commercially obtained Si3N4 ball (NBD 200, Cerbec, East Granby, CT) of radius 6.35 mm was rotated in contact with three flat disk specimens (thickness 2 mm; diameter 4 mm) aligned with their surface normals in tetrahedral coordination relative to the rotation axis. This was accomplished in a wear testing assembly to ensure equal distribution of the applied load on the three specimens. Paraffin oil (Heavy Grade, Fisher Scientific, Fair Lawn, NJ) with a viscosity of 3.4 · 10 5 m2 s 1 at 40 C was used as the lubricant to preclude any tribological effects such as friction-induced heating. The contact load on each disk was 70 N and the rotation speed was 100 rpm. The wear tests were interrupted at intervals, and the diameters of the circular wear scars on each disk were measured using the optical microscope (two

orthogonal measurements per disk, three disks per ceramic). After each interruption the specimens were placed back in the tribometer in exactly the same position using a precision fixture. The wear-scar diameter was used to quantify the extent of wear damage. Finally, the wear damage was observed in the SEM. Nanoindentation (Micromaterials, Wrexham, UK) experiments were performed using a Berkovich indenter (tip radius <100 nm) to measure the hardness of individual SiC grains and the YAG intergranular phase in the SiC–Ar and SiC–N2 ceramics with higher YAG content (23.2 vol.%). Tests were conducted on polished surfaces with variable loads, such that the maximum penetration depth was maintained constant at 20 nm. Hardness was evaluated from the indentation load–displacement curves, using standard formulae [20]. Care was taken to ensure that the hardness impressions were contained entirely within the SiC or the YAG phases. Conventional Vickers indentation tests (MV-1, Matsuzawa, Tokyo, Japan) were performed to evaluate the toughness (KIC) of the SiC–Ar and SiC–N2 ceramics. Tests were performed with 98 N load, and the toughness values were determined using the standard procedure and formula [21]; elastic modulus (E) values determined using Hertzian indentation were used in the KIC calculations [22]. Table 1 summarizes the characteristics of the two LPS a-SiC materials studied here. Figure 1a and b

Figure 1. SEM micrographs of polished and plasma-etched crosssections of: (a) SiC–Ar ceramic and (b) SiC–N2 ceramic. The dark and light regions are the SiC grains and the YAG phase, respectively. The core-shell substructure within the SiC grains is visible in the micrographs due to differential plasma etching.

Table 1. Microstructural parameters and mechanical properties of the SiC–Ar and SiC–N2 ceramics Ceramic

Density (%)

Grain size (lm)

Grain aspect ratio

Nitrogen content (wt.%)

Elastic modulus (GPa)

SiC hardness (GPa)

YAG hardness (GPa)

Indentation toughness (MPa m0.5)

SiC–Ar SiC–N2

100 100

0.8 0.8

1.4 1.4

0 0.648

390 391

31 31

19 23

2.9 2.6

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shows representative SEM micrographs of the microstructures of SiC–Ar and SiC–N2 ceramics, respectively. As seen in these micrographs, these two ceramics are fully dense, consistent with the density measurements presented in Table 1. The microstructures of both the ceramics appear very similar, with fine, equiaxed SiC grains embedded in a continuous crystalline YAG matrix [8,18]. The only difference between these two ceramics is that the SiC–N2 ceramic is found to contain 0.648 wt.% nitrogen, while no nitrogen was detected in the SiC–Ar ceramic. Since nitrogen has negligible solubility in the SiC grains (100 ppm [23]), it has been shown that most of the nitrogen in SiC–N2 material is in solid-solution with the YAG intergranular phase [18]. The nanoindentation hardness of individual SiC grains was found to be 31 GPa in both the SiC–Ar and SiC–N2 ceramics. On the contrary, the nanoindentation hardnesses measured for the intergranular YAG phase in the SiC–Ar and SiC–N2 ceramics were different (19 and 23 GPa, respectively). This 21% increase in hardness in the SiC–N2 material can be attributed to the incorporation of nitrogen in the intergranular phase [24]. Note that while all these nanoindentation-hardness values could be affected by the well-known size effect in hardness [25], here we will use the nanoindentationhardness measurements only for comparison purposes between SiC–Ar and SiC–N2 ceramics. Figure 2 shows sliding-wear data for SiC–Ar and SiC–N2 ceramics. As expected, both ceramics show the sliding-wear behavior of polycrystalline ceramics: initial mild wear with little material removal, followed by an abrupt transition to severe wear with extensive material removal [9–17]. From Figure 2, it is clear that, despite the similar microstructures, SiC–N2 ceramic has dramatically improved wear resistance over SiC–Ar ceramic: a lower mild-wear rate (by a factor of 2.7), a longer transition time (by a factor of 2.9) and a similar severe-wear rate. That marked improvement in the sliding-wear resistance with the incorporation of nitrogen in the intergranular phase is also evident in the SEM micrographs of the worm surfaces. Figure 3a and b shows

Figure 2. Wear scar diameter as a function of the sliding time for the SiC–Ar (close) and SiC–N2 (open) ceramics. Each datum point represents an average of three specimens tested per ceramic; error bars represent data dispersion. The solid lines are fits to the data, with the discontinuities in the lines indicating mild wear to severe wear transitions. The box on the left axis indicates the dimension of a single Hertzian contact in both ceramics at the same wear test contact load (70 N).

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Figure 3. SEM micrographs of the wear damage after 500 min of sliding in the: (a) SiC–Ar ceramic and (b) SiC–N2 ceramic. The existence of grain-boundary cracking as well as grain and intergranular phase pullout in the images is indicative of the fracture-controlled wear regime.

SEM micrographs taken at the center of wear scars in SiC–Ar and SiC–N2 ceramics, respectively, at the end of the wear tests (500 min). The higher severity of wear damage in the form of grain pullout in the SiC–Ar ceramic relative to the SiC–N2 ceramic is evident from these micrographs. These results can be analyzed using the mechanistic wear model of Cho et al. [15] as applied to the LPS SiC case, details of which are given elsewhere [9–13]. Briefly, during the mild-wear stage, stresses accumulate as a function of time at the SiC/YAG interfaces as a consequence of dislocation plasticity in both the SiC grains and the YAG intergranular phase. When the stress intensity factor due to these accumulating stresses on pre-existing grain-boundary flaws exceeds the grainboundary toughness, grain-boundary fracture and subsequent grain pullout takes place. This leads to rapid material removal in the severe-wear stage. In this context, the presence of harder phases in LPS SiC, in which dislocation activity is expected to be hindered, is likely to result in slower accumulation of stresses. This is expected to result in lower mild-wear rates and longer transition times. Since the microstructures of SiC–Ar and SiC–N2 ceramics are identical and the SiC grains have the same hardness in the two materials, the reduced mild-wear rate, and the attendant delayed wear transition in the SiC–N2 ceramic can only be attributed to the higher hardness of the intergranular phase in that ceramic. Note that a higher grain-boundary toughness could also result in delayed wear transition. However, the sintering atmosphere does not appear to influence the grain-boundary toughness, as attested by the similar indentation toughnesses of the two ceramics (Table 1) for intergranular indentation cracks. Such indentation toughness results also explain why the severe-wear rates are similar in SiC–Ar and SiC–N2 ceramics. The results presented in this work show that hardening of the intergranular phase is another approach to

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improving the sliding-wear resistance in LPS SiC ceramics. This approach is generic in nature, and could be extended to other LPS ceramics amenable to intergranular phase hardening. Thus, based on these results and previous studies [9–13], the following modified guidelines for the design of highly wear resistant LPS ceramics emerge: (i) grain refinement [9,13], (ii) grain elongation [10,11,13], (iii) reduction in intergranular phase content [9,13] and (iv) intergranular phase hardening. This work was supported by the Ministerio de Ciencia y Tecnologı´a (Government of Spain), and the Fondo Europeo de Desarrollo Regional (FEDER), under Grant Nos. CICYT MAT 2004-05971 and UNEX0023-013. [1] D.W. Richerson, Modern Ceramic Engineering, Marcel Dekker, New York, 1992. [2] N.P. Padture, J. Am. Ceram. Soc. 77 (1994) 519. [3] S.M. Hsu, M. Shen, Wear 256 (2004) 867. [4] R.P. Jensen, W.E. Luecke, N.P. Padture, S.M. Wiederhorn, Mater. Sci. Eng. A 282 (2000) 109. [5] L.S. Sigl, J. Eur. Ceram. Soc. 23 (2003) 1115. [6] A.L. Ortiz, A. Mun˜oz-Bernabe´, O. Borrero-Lo´pez, A. Domı´nguez-Rodrı´guez, F. Guiberteau, N.P. Padture, J. Eur. Ceram. Soc. 24 (2004) 3245. [7] K. Suzuki, M. Sasaki, Pressureless sintering of silicon carbide, in: S. Somiya, R.C. Bradt (Eds.), Fundamental Structural Ceramics, Terra Scientific Publishing Company, Tokyo, 1987, pp. 75–87. [8] H. Xu, T. Bhatia, S.A. Deshpande, N.P. Padture, A.L. Ortiz, F.L. Cumbrera, J. Am. Ceram. Soc. 84 (2001) 1578.

[9] O. Borrero-Lopez, A.L. Ortiz, F. Guiberteau, N.P. Padture, J. Am. Ceram. Soc. 88 (2005) 2159. [10] O. Borrero-Lopez, A.L. Ortiz, F. Guiberteau, N.P. Padture, J. Am. Ceram. Soc. 88 (2005) 3531. [11] O. Borrero-Lopez, A.L. Ortiz, F. Guiberteau, N.P. Padture, J. Am. Ceram. Soc. 90 (2007) 541. [12] O. Borrero-Lopez, A.L. Ortiz, F. Guiberteau, N.P. Padture, J. Eur. Ceram. Soc. 27 (2007) 2521. [13] O. Borrero-Lopez, A.L. Ortiz, F. Guiberteau, N.P. Padture, J. Eur. Ceram. Soc. 27 (2007) 3351. [14] S.-J. Cho, H. Moon, B.J. Hockey, S.M. Hsu, Acta Metall. Mater. 40 (1992) 185. [15] S.-J. Cho, B.J. Hockey, B.R. Lawn, S.J. Bennison, J. Am. Ceram. Soc. 72 (1989) 1249. [16] S.C. Thompson, A. Pandit, N.P. Padture, S. Suresh, J. Am. Ceram. Soc. 85 (2002) 2059. [17] X. Wang, N.P. Padture, H. Tanaka, A.L. Ortiz, Acta Mater. 53 (2005) 271. [18] A.L. Ortiz, T. Bhatia, N.P. Padture, G. Pezzotti, J. Am. Ceram. Soc. 85 (2002) 1835. [19] S.A. Deshpande, T. Bhatia, H. Xu, N.P. Padture, A.L. Ortiz, F.L. Cumbrera, J. Am. Ceram. Soc. 84 (2001) 1585. [20] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564. [21] G.R. Anstis, P. Chantikul, D.B. Marshall, B.R. Lawn, J. Am. Ceram. Soc. 64 (1981) 533. [22] B.R. Lawn, J. Am. Ceram. Soc. 81 (1998) 1977. [23] N.W. Jepps, T.F. Page, Polytypic transformation in silicon carbide, in: P. Krishna (Ed.), Crystal Growth and Characterization of Polytype Structures, vol. 7, Pergamon Press, Oxford, 1983, pp. 259–307. [24] T.E. Mitchell, A.H. Heuer, Mater. Sci. Eng. 28 (1977) 81. [25] W.D. Nix, H.J. Gao, J. Mech. Phys. Solids 46 (1998) 411.