Grain bridging locations of monolithic silicon carbide by means of focused ion beam

milling technique

N. Al Nasiri1, E. Saiz, F. Giuliani and L.J. Vandeperre

Centre for Advanced Structural Ceramics, Department of Materials, Imperial College

London, South Kensington Campus, London SW7 2BP, United Kingdom

Abstract

A slice and view approach using a focused ion beam (FIB) milling technique was employed

to investigate grain bridging near the tip of cracks in four silicon carbide (SiC) based

materials with different grain boundary chemistries and grain morphologies. Using traditional

observations intergranular fracture behaviour and hence clear evidence of grain bridging was

found for SiC based materials sintered with oxide additives. More surprisingly, in large grain

materials, the FIB technique reveals evidence of grain bridging irrespective of the grain

boundary chemistry, i.e. also in materials which macroscopically fail by transgranular failure.

This helps to explain why the toughness of large grained materials is higher even if failure is

transgranular.

Keywords: SiC, FIB, grain bridging and grain boundary chemistry.

Introduction

The strong covalent bonding between silicon and carbon atoms gives SiC materials an

excellent oxidation resistance, strength retention to high temperatures, high wear resistance,

high thermal conductivity and good thermal shock resistance. Nowadays, SiC is recognized

as one of the leading structural ceramics for gas turbines [1], for mechanical seals and brakes

11 Corresponding author. Tel: +44 207 5895111 Email address: n.al-nasiri10@imperial.ac.uk

[2] and for bioengineering application (i.e. implants) [3, 4]. On the other hand, there is a

limitation to the usefulness of silicon carbide materials due to their low fracture toughness

and high strength dispersion, i.e. an inhomogeneous and large distribution of flaws and

defects.

During the last 30 years a lot of attention has been devoted to improving the toughness of

ceramic materials to ensure that the damage during processing and post-processing can be

acceptable without compromising the structure’s reliability. Ceramic materials have been

toughened using different approaches. A very successful approach consists of crack tip

shielding through grain bridging in the wake of the crack tip [5]. The toughness increase

arises because crack deflection leaves grains connected to both crack faces and these grains

transmit a closing force across crack walls trying to prevent the crack from further extension

by reducing the applied stress intensity at the crack tip [6]. The resistance to crack

propagation due to grain bridging in general increases with increasing crack length, which

leads to the so called Resistance-curve (R-curve) [7-9]. During the last 30 years, the R-curve

behaviour of ceramic materials has been extensively studied for long cracks. What may

happen for short cracks (< 100 µm) and its effect on toughness has until recently largely been

overlooked. A material with flat R-curve may still have a steep rise at the beginning of the

test as suggested by Fett et al. [10]. In materials with a steep rising R-curve, one may expect

that toughness will behave differently for short cracks than for long cracks. The

understanding of this concept is highly relevant as ceramic materials are generally produced

with small natural defects in the order of few microns [11]. While there is general agreement

that the action of grains a long distance behind the crack tip is a consequence of friction

between the grains and the matrix from which they are being pulled, the steep rise in R-curve

at the onset of failure is increasingly attributed to elastic bridging, i.e. where the grain

bridging the crack faces are still to some extent fully bonded to both fracture surfaces.

Different approaches were considered to study the effect of elastic bridging i.e. intact grains

still bonded to both crack faces bridging the crack near its tip. For example Foulk et al. [12]

proposed a 2D model of bridging by a single elongated grain based on finite element analysis

during the early stage of the rising R-curve. This increase in crack resistance in very short

cracks before forming grain bridges is in line with Fett et al. [10]. Fünfschilling and co-

workers [13] observed that silicon nitride materials sintered with different additives show

different R-curve behaviour ranging from flat, moderate and rising R-curve. However, the

main common factor between the studied microstructures is that they all have a steep rising

R-curve in the first 10 µm of crack extension. Still, direct evidence from observation of such

elastic grain bridging sites close to the tip of the crack has so far not been reported.

Experimental methods

In this work, four silicon carbide based materials were produced: two using solid state

sintering with boron and carbon (SS) and two using liquid phase sintering using alumina and

yttria as sintering additives (LP), see Figure 1. The processing details are described in [14].

The LP-Fine consists of homogenous fine equiaxed grains with average grain size of 1.0 ±

0.1 µm while LP-Coarse has elongated grains with a length of 16 ± 1 µm and a width of 5 ±

0.3 µm. For the SS-Fine, the grains have an average size of 4.0 ± 0.4 µm and the grains in the

SS-Coarse have a length of 18 ± 1 µm and a width of 5 ± 0.3  µm. The samples were ground

and polished up to 1 µm using diamond suspension. A crack is introduced by placing an

indent of 30 kg on the polished surface. Once the crack tip was found an area with

dimensions of 4 µm by 5 µm by 5 µm, Figure 2, was selected for investigation with a slice

and view approach, i.e. progressive milling in small steps (20 nm) along the X-axis and

taking micrographs in between each milling stage to create a 3D information set of the

microstructure and crack.

Figure 1: Scanning electron micrographs of chemically etched silicon carbide materials of this work.

Figure 2: Sketch showing the selected coordinates to identify planes and region of interest for the FIB slice and view process.

The R-curve measurements were carried out in situ in the SEM (vacuum level of 1 x 10-3 Pa)

using a constant moment double cantilevered test rig, built in line with the set-up proposed by

Sorensen et al.[15]. The silicon carbide specimens were 65 mm by 10 mm by 5 mm. One side

of the samples was polished to facilitate the observation of the crack tip. The speed of the

stage motor was 0.1 mm min-1. Every time crack propagation was observed, the applied load

was recorded and a SEM micrograph was taken.

Results and discussion

As expected in the LP-sintered materials, cracks are deflected at grain boundaries, whereas in

the SS-sintered materials, the cracks grow much straighter and do not appear to be influenced

strongly by grain boundaries, see Figure 3. However, a closer look at the coarse grained

materials revealed that close to the tip of the crack, fracture is slightly complicated. While

fracture continues to be along the grain boundaries in LP-Coarse and more cleavage-like in

SS-Coarse, ligaments of material which remain solidly bonded to both sides of the fracture

surface can be found in both materials.

Figure 3: FIB micrographs showing crack deflection for LP-Fine and SS-Fine and elastic bridging for LP-Coarse and SS-Coarse. Squares represent the area of interest, where elastic bridging takes place.

In LP-Coarse sequence it is clear that the main crack is found to change from propagating on

the right of a grain initially but changes to propagating on the left of a grain as the distance

from the crack tip increases as noted at the bottom box in Figure 3. For the SS-Coarse, it is

clear that the crack front has split at a grain and is moving forward either side of it, while

leaving a sliver of material connecting the two fracture surfaces.

The evolution of another elastic bridge in LP-Coarse is shown in Figure 4. It was not possible

to pin point the exact location of the crack tip during slice and view due to crack propagation

driven by a relaxation of the residual stresses when the material surrounding the crack was

milled away. However, by taking the actual crack tip location as a reference (X = 0 nm), it

can be stated that the crack propagated an unknown distance X0 beyond the first slice

obtained. The location of subsequent slices has been determined from the thickness of each

slice (20 nm) and as a function of the X0 distance, indicating the propagation of the crack up

to the crack tip. The crack starts splitting at a distance of X 0 + 2200 nm and it reaches its

maximum separation between its branches at a distance of X0, which is the closest slice

obtained from the crack tip.

Figure 4: FIB micrographs and corresponding sketches of the crack path at different distances approaching the crack tip of LP-Coarse.

At a distance of X0+2280 nm, the crack bumps into a grain, which is shown as a small curve

comparing to the rest of the crack path. Part of the crack path is not visible at a distance of

X0+2200 nm, because the grain is holding it from either side to prevent further propagation.

Until X0+1800 nm, the separation between the crack faces is small and occurs in a different

direction (perpendicular to the crack front). A noticeable big separation is observed at a

distance of X0+1500 nm. Further splitting of crack faces takes place, and the separation

between the crack segments is largest at a distance equal to X0 + 300 nm, hence near the crack

tip.

The lack of grain bridges near the crack tip in SS-Fine SiC relative to SS-Coarse can explain

why the long crack toughness of the latter is i.e. ~ 1 MPa m1/2 higher while both materials

have flat R-curves in the long crack regime (Figure 5). Similarly, the R-curve of LP-Coarse

SiC quickly rises to 5 MPa m1/2 in the first 100’s of micrometre of crack propagation, whereas

LP-Fine shows a slow rising R-curve starting from ~ 2 MPa m1/2. Hence, the elastic bridging

near the tip of the crack seems to require the presence of large grains. It is likely that this is

the result of residual stresses due to differential thermal expansion, leading to protective

compressive stresses which prevent the crack from penetrating the large grains, or the

presence of microcracks which help the crack front to split.

Figure 5: R-curves of the four SiC materials measured by DCB measurements under vacuum level of 1x10-3 Pa. Note the rising R-curves in LP-SiC compared to flat R-curves in SS-SiC.

Conclusion

The work provides evidence that elastic grain bridging occurs in coarse grained silicon

carbide regardless of the grain boundary chemistry, whereas no evidence of such elastin

ligaments was found in fine grained microstructures. The presence and absence of elastic

bridges as observed here correlates with toughness values as measures for short extension of

cracks elsewhere.

Acknowledgments

The authors would like to thank the Engineering and Physical Sciences Research Council

(EPSRC) of the UK for the funding of this project through grant EP/F033605/1.

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