Silicon Nitride Grain Boundary Glasses: Chemistry, Structure and

Key Engineering Materials Vol. 484 (2011) pp 46-51 Online available since 2011/Jul/04 at www.scientific.net © (2011) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/KEM.484.46

Silicon Nitride Grain Boundary Glasses: Chemistry, Structure and Properties Stuart Hampshire a and Michael J. Pomeroyb Materials and Surface Science Institute, University of Limerick, Limerick, Ireland a

b

[email protected], [email protected]

Keywords: silicon nitride, intergranular glass, oxynitride glass, Si3N4, liquid phase sintering, microstructural engineering, chemistry-microstructure-property relationships

Abstract. Silicon nitride is recognised as a high performance material for both wear resistant and high temperature structural applications. Oxide sintering additives such as yttrium oxide and alumina are used to provide conditions for liquid phase sintering, during which the additives react with surface silica present on the Si3N4 particles and some of the nitride to form an oxynitride liquid which allows densification and transformation of α- to β-Si3N4 and on cooling remains as an intergranular oxynitride glass. This paper provides an overview of liquid phase sintering of silicon nitride ceramics, grain boundary oxynitride glasses and the effects of chemistry and structure on properties. As nitrogen substitutes for oxygen in oxynitride glasses, increases are observed in glass transition and softening temperatures, viscosities, elastic moduli and microhardness. These property changes are compared with known effects of grain boundary glass chemistry in silicon nitride ceramics. Introduction – Silicon Nitride Ceramics Silicon nitride is one of the major structural ceramics that has been the subject of many years of intensive research effort [1-4]. It possesses high flexural strength, good fracture resistance, good creep resistance, high hardness and excellent wear resistance. These properties arise through the processing of the material which involves liquid phase sintering and the development of microstructures in which high aspect ratio grains and intergranular glass phase lead to excellent fracture toughness and high strength. Sintering of silicon nitride is carried out at 1750-1900°C under a nitrogen atmosphere (0.1 to 10 MPa). Sintering additives, such as yttria and alumina, are mixed with α-Si3N4 powder which has layers of silica on their particle surfaces. The oxides react with this silica and some of the nitride itself at sintering temperatures to form an oxynitride liquid which promotes liquid phase sintering and densification by solution-precipitation [1-5]. Various rare earth (RE) oxides [6-8], with either alumina or MgO, have also been investigated as sintering additives and this allows the development of different types of microstructures and, hence, properties by modifying the chemistry of the sintering liquid phase. A systematic study of pressureless sintering kinetics for silicon nitride ceramics [5] applied the Kingery liquid-phase sintering model [9] in which three stages are identified: (1) Particle Rearrangement following formation of the initial liquid phase, where the rate and the extent of shrinkage depend on both the volume and viscosity of the liquid; the volume depends on the total amount of oxide additives while the viscosity depends on additive (e.g.Y2O3:Al2O3) ratios; this stage is also the incubation period for the α→β transformation. (2) Solution-diffusion-reprecipitation, where, according to Kingery, shrinkage is given by: ∆V/Vo α tl/n (6) where t is time and n = 3 if solution into or precipitation from the liquid is rate controlling, as is the case for lower viscosity liquids (MgO, low Z REs), and n = 5 if diffusion through the liquid is ratecontrolling, as is the case for high Z REs, where diffusion through a more viscous oxynitride liquid All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of TTP, www.ttp.net. (ID: 193.1.104.2-03/01/12,12:48:53)

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is much slower. The α→β transformation begins in this stage and proceeds by dissolution of α and precipitation of Si and N onto existing β grains, which grow in the longitudinal direction as prismatic hexagonal rod-like crystals that eventually impinge on each other forming an interlocked microstructure. With lower viscosity liquid phases, the transformation proceeds when densification is almost complete and is enhanced by high N solubility within the liquid phase [5, 8]. (3) Elimination of closed porosity and grain growth, during which the liquid acts to further grow the elongated β grains but this critically depends on the liquid-phase composition, established by the additives used, and also the use of high α or high β content powders or β seeds, all of which have a significant impact on phase transformation, microstructural development and densification [6-8]. The liquid cools as an intergranular phase, usually a glass. Fig. 1 shows a scanning electron micrograph of silicon nitride densified with 6 wt.% yttria and 2 wt.% alumina [10]. The microstructure consists of β-Si3N4 grains, with high aspect (length to diameter) ratios, surrounded by a Y-Si-Al-O-N glass phase (white). Fig. 2 shows a high resolution transmission electron microscope (HRTEM) image of two β-Si3N4 grains separated by an intergranular glass film (IGF) leading to an oxynitride glass triple pocket. Grain-boundary thicknesses and compositions are characteristic of the metal oxide additive system, and film thickness (in the range 0.5–1.5 nm) depends strongly on chemical composition [11, 12]. The evolution of β-Si3N4 microstructures during sintering is influenced by the adsorption of RE cations at silicon nitride grain surfaces and by the viscosity of the intergranular phases. Theoretical and scanning transmission electron microscopy [8] show that RE atoms exhibit different tendencies to segregate from the liquid to grain surfaces and have different binding strengths at these surfaces. TP

β-Si3N4 grain

G F

β-Si3N4 grain

5nm

1µm

Figure 1. Scanning electron micrograph of silicon nitride (6 wt.% Y2O3 + 2 wt.% Al2O3) showing dark β-Si3N4 grains and bright YSiAlON glass

I

Figure 2. TEM micrograph of silicon nitride showing two β-Si3N4 grains, a triple point (TP) glass pocket and intergranular glass film (IGF)

Grain Boundary Glass Chemistry and Properties In silicon nitride ceramics, the amounts and ratios of the additives initially introduced determine the quantity and chemistry of the glass phase and this affects properties such as fracture toughness, ambient and high temperature strengths, creep resistance and oxidation resistance [6-8, 13-16]. Substitution of nitrogen for oxygen in alumino-silicate glasses induces greater coordination of the glass network due to the tri-coordinate bonding of nitrogen compared to the bi-coordinate bonding of oxygen, and this results in increases in elastic modulus, hardness, glass transition temperature and viscosity [17-21]. With respect to modifier cation type, evidence shows that as the cation field strength (valency / square of ionic radius) of the modifier increases then so do elastic modulus, hardness, glass transition temperature and viscosity [19].

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10.5

2

R = 0.997

10 150 9.5 2

140

R = 0.982

130

32Y:42Si:26Al: (100-y)O: yN glasses

120

9 8.5 8

0

5

10

15

20

25

nitrogen content (eq. %) E

microhardness

Figure 3. Effect of nitrogen content on Young’s modulus (E) and microhardness for glasses with cation composition (in eq.%): 37Y:42Si:21Al [21]

log(viscosity / Pa. s)

E (GPa)

160

microhardness (GPa)

11

170

15 14 13 12 11 10 9 8 7 0

10

20

30

nitrogen content / eq.% 20La 60Si 20Mg (875C) 30La 45Si 25Al (930 C)

28Y 56Si 16Al (930C) 30Y 45Si 25Al (925C)

Figure 4. Effect of nitrogen content on viscosity (log scale) for various glasses (cation compositions in eq.%) at 875, 925 and 930°C [22-25]

When cation ratios (M:Si:Al) are kept constant, linear correlations are typically observed between nitrogen content and Young’s modulus or microhardness of oxynitride glass. Fig. 3 shows a typical plot for Y-Si-Al-O-N glasses with Y:Si:Al = 37:42:21 (eq.%). Elastic Modulus is a function of bond energies, network compactness and cross-linking of the glass network and it has been shown that increases in elastic modulus with N can be related to increases in fractional glass compactness and decreases in molar volume [22, 26]. The effect of nitrogen substitution for oxygen on glass viscosity has not been studied widely. The carefully collected data which does exist indicates that glass viscosity increases with increasing nitrogen substitution and, like elastic modulus, hardness, glass transition temperature and softening temperatures, varies linearly with nitrogen content as shown in Fig. 4 [22-25]. The variation in viscosity values with nitrogen content depends on cation type and cation ratios. For an increase in nitrogen content from 0 to 20 eq.%, viscosity increases can vary between 2.0 and 3.6 orders of magnitude. Thus, for the Y- and La-Si-Al-O-N oxynitride glasses [17], viscosity increases by about 0.10 to 0.13 orders of magnitude per eq.% N [17, 24] while for 20La:60Si:20Mg (eq.%) oxynitride glasses [25], the increase is ~0.18 orders of magnitude per eq.% N which indicates that the use of Mg in conjunction with N and a rare earth modifier may result in Mg taking on the role of a network former as observed by Rouxel et al. [27]. Effects of Grain Boundary Glass Chemistry on Properties of Silicon Nitride Becher et al. [14] reported significant improvements in the fracture resistance of self-reinforced silicon nitride ceramics by tailoring the chemistry of the intergranular amorphous phase. They found that the steady-state fracture toughness values of these silicon nitrides increased with the Y:Al ratio of the oxide additives as shown in Fig. 5. The increased toughness was accompanied by a steeply rising R-curve and extensive interfacial debonding between the elongated β-Si3N4 grains and the intergranular glassy phase. Weak interfaces, which should be related to lower Young’s modulus of the glass, favour high fracture toughness. Fig. 6 shows the effect of Y:Al ratio on Young’s modulus of Y-Si-Al-O-N glasses which decreases with increasing Y:Al ratio, i.e. with easier interfacial debonding and increasing KIc in Si3N4. Compared to silicon nitrides with low Y:Al ratios, the high Y:Al ratio materials exhibited more extensive debonding at grain boundary interfaces, resulting in

11 10 9 8 7 6

49

170 E (GPa)

K Ic (MPa.m -1/2)


Easier interfacial de-bonding

Easier de-bonding

160 150

20N

140

10N

130 0

1 2 Y:Al molar ratio

3

Figure 5. Fracture toughness of silicon nitride sintered with different Y2O3:Al2O3 ratios. Molar ratio Y:Al 0.64 1.12 2.8 Weight ratio Y2O3:Al2O3 4:2.8 5:2 6.25:1

0

1

2

3

Y:Al molar ratio Figure 6. Young’s modulus of Y-Si-Al-O-N glasses with 10 and 20 eq.% N as a function of Y:Al ratio. Molar ratios of oxide additives as shown in Fig. 5 are superimposed.

increased intergranular fracture. Microstructural and chemistry characterization revealed that the Y:Al ratios in the additives influence the atomic bonding structure across the β-Si3N4/intergranular glass interface by altering the composition of the glassy phase and inducing different Al and O contents in the growth region of the elongated grains. In order to gain further understanding of the influence of intergranular glass on the fracture toughness of silicon nitride, the debonding behaviour of the interface between the prismatic faces of β-Si3N4 whiskers and oxynitride glasses was investigated in model systems based on various Si-(Al)-Y(Ln)-O-N (Ln = rare-earth) oxynitride glasses [28]. It was found that while the interfacial debonding strength increased when an epitaxial β'-SiAlON layer grew on the β-Si3N4 whiskers, the critical angle for debonding was lowered with increasing Al and O concentrations in the SiAlON layer showing that by tailoring the densification additives and hence the chemistry of the intergranular glass, it is possible to improve the fracture resistance of silicon nitride. The impact of various rare-earth and related doping elements (R = Lu, Sc, Yb, Y, Sm, La) on the grain growth anisotropy and the mechanical properties of polycrystalline silicon nitride ceramics has been studied using model experiments [29], in which Si3N4 particles are able to grow freely in an RE-Si-Mg-O-N glass matrix. With increasing ionic radius of the RE, grain anisotropy increases due to non-linear growth kinetics. Toughness and strength are affected by the rare-earth element. Samples of equivalent grain sizes and morphologies yield an increasing toughness with increasing ion size of the RE3+, reflecting an increasingly intergranular crack path. The choice of RE is essential to tailor microstructure, interfacial strength and mechanical properties. A first-principles model, the differential binding energy (DBE), was developed [29] to characterize the competition between RE and Si as they migrate to the β-Si3N4 grain surfaces. The theory predicts that, of the various rare earths, La should have the strongest and Lu the weakest preferential segregation to the grain surfaces. Elements with larger positive DBE values than Si prefer to reside in regions containing oxygen while those with negative values have a preference for the nitrogen-terminated Si3N4 grain surfaces, even more so than Si. Additional calculations defined the adsorption sites and their binding strengths for each of the rare earths on the prismatic plane of the β-Si3N4 grains. These predictions were confirmed by unique atomic-resolution images obtained by aberrationcorrected Z-contrast scanning transmission electron microscopy (STEM) [30]. The combined theoretical and STEM studies revealed that the elements that induce the greatest observed grain anisotropy are those with the strongest preferential segregation plus high binding strength to the prismatic grain surface.

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At temperatures exceeding 1000°C, strengths decrease due to the softening of the intergranular glass phase. Grain boundary chemistry, effective viscosity and volume fraction of the intergranular glass phase control creep rate and formation and growth of cavities in the amorphous phase [31]. Observations of intergranular films in silicon nitride [15] show that their width decreases with decrease in RE ion radius. Viscous flow of these films contributes to the initial stage of tensile creep deformation [31]. Smaller RE ions (Lu, Er) prefer to segregate at the triple points [30]. Larger RE ions have a preference for N and remain concentrated in the IG films. Creep behaviour is thus dependent on both intergranular film and triple point glass viscosities. Summary Oxide sintering additives such as yttrium oxide or lanthanide rare earth oxides and alumina are used to provide conditions for liquid phase sintering of silicon nitride, during which the additives react with surface silica present on the Si3N4 particles and some of the nitride to form an oxynitride liquid which allows densification and transformation of α- to β-Si3N4 and on cooling remains as an intergranular oxynitride glass. The presence of nitrogen in these types of glasses results in increases in elastic modulus, hardness, glass transition temperature and viscosity because of extra crosslinking of the glass network. These properties also increase with increasing cation field strength of the rare earth modifier. Properties such as elastic modulus and viscosity are also dependent on Y or RE:Al ratio. In silicon nitride ceramics, the amounts and ratios of the additives initially introduced determine the quantity and chemistry of the glass phase and this affects properties such as fracture toughness, ambient and high temperature strengths, creep resistance and oxidation resistance. Significant improvements in fracture resistance of self-reinforced silicon nitride ceramics can be achieved by tailoring the chemistry (Y or RE):Al ratio of the intergranular amorphous phase. The increased toughness is accompanied by a steeply rising R-curve and extensive interfacial debonding between the elongated β-Si3N4 grains and the intergranular glassy phase. Weak interfaces, which should be related to lower Young’s modulus of the glass, favour high fracture toughness. Silicon nitrides with different rare earth additives of equivalent grain sizes and morphologies exhibit increasing toughness with increasing ion size of the RE3+, reflecting an increasingly intergranular crack path. The choice of rare earth modifier is essential to optimise microstructure, interfacial bond strength and mechanical properties. Grain boundary glass chemistry, effective viscosity and volume fraction of the intergranular glass phase control creep rate and behaviour of silicon nitride ceramics at high temperatures. References [1] S. Hampshire: Metals Forum, Vol. 7 (1984) p. 162 [2] S. Hampshire: in: Structure and Properties of Ceramics, edited by M. Swain, Volume 11 of Materials Science and Technology - A Comprehensive Treatment, Chapter 3, VCH Verlagsgesellschaft, Weinheim (1994) [3] F. L. Riley: J. Am. Ceram. Soc., Vol. 83 (2000) p. 245 [4] F. F. Lange: J. Ceram. Soc. Japan, Vol. 114 (2006) p.873 [5] S. Hampshire, K. H. Jack: Proc. Brit. Ceram. Soc., Vol. 31 (1981) p. 37 [6] E. Y. Sun, P. F. Becher, K. P. Plucknett, C.-H. Hsueh, K. B. Alexander, S. B. Waters, K. Hirao, M. E. Brito: J. Am. Ceram. Soc., Vol. 81 (1998) p. 2831 [7] R. L. Satet, M. J. Hoffmann, R. M. Cannon: Mater. Sci. Eng. A, Vol. 422 (2006) p. 66


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Silicon Nitride Grain Boundary Glasses: Chemistry, Structure and Properties 10.4028/www.scientific.net/KEM.484.46 DOI References [7] R. L. Satet, M. J. Hoffmann, R. M. Cannon: Mater. Sci. Eng. A, Vol. 422 (2006) p.66. doi:10.1016/j.msea.2006.01.015 [8] P. F. Becher, N. Shibata, G. S. Painter, F. Averill, K. van Benthem, H. -T. Lin, S. B. Waters: J. Am. Ceram. Soc., Vol. 93 (2010) p.570. doi:10.1111/j.1551-2916.2009.03435.x [9] W. D. Kingery: J. Appl. Phys., Vol. 30 (1959) p.301. doi:10.1002/cber.19590921146 [10] S. Hampshire, M. J. Pomeroy: Int. J. Appl. Ceram. Tech., Vol. 5 (2008) p.155. doi:10.1111/j.1744-7402.2008.02205.x [14] P. F. Becher, E. Y. Sun, K. P. Plucknett, K. B. Alexander, C. -H. Hsueh, H. -T. Lin, S. B. Waters, C. G. Westmoreland, E. -S. Kang, K. Hirao, M. E. Brito: J. Am. Ceram. Soc., Vol. 81 (1998) p.2821. doi:10.1111/j.1151-2916.1998.tb02702.x [15] H. J. Kleebe, G. Pezzotti, G. Ziegler: J. Am. Ceram. Soc., Vol. 82 (1999) p.1857. doi:10.1002/chin.199938018 [16] S. M. Wiederhorn, R. F. Krause, F. Lofaj, U. Taffner: Key Eng. Mater., Vol. 287 (2005) p.381. doi:10.4028/www.scientific.net/KEM.287.381 [18] S. Hampshire, E. Nestor, R. Flynn, J.L. Besson, T. Rouxel, H. Lemercier, P. Goursat, M. Sabai, D.P. Thompson and K. Liddell: J. Eur. Ceram. Soc., Vol. 14 (1994) p.261. doi:10.1016/0955-2219(94)90095-7 [19] R. Ramesh, E. Nestor, M. J. Pomeroy, S. Hampshire: J. Eur. Ceram. Soc., Vol. 17 (1997) p. (1933). doi:10.1016/S0955-2219(97)00057-5

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