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5vol.%SiC nanocomposites
Acta Materialia 51 (2003) 149–163 www.actamat-journals.com
Subsurface damage analysis by TEM and 3D FIB crack mapping in alumina and alumina/5vol.%SiC nanocomposites H.Z. Wu ∗, S.G. Roberts, G. Mo¨bus, B.J. Inkson Department of Materials, University of Oxford, Oxford OX1 3PH, UK Received 23 October 2001; received in revised form 17 July 2002; accepted 17 July 2002
Abstract TEM and 3D crack analysis by focused ion beam (FIB) cross sectioning have been used to quantify the subsurface damage beneath scratches made by a 120° cone indenter loaded to 1 N in monolithic polycrystalline alumina and alumina/5vol.%SiC nanocomposites. In the nanocomposite, an extensive plastic deformation zone was found under the scratch grooves, extending beyond the first layer of grains to a maximum plastic deformation depth of ~7 µm below the surface of the track. In the alumina, however, the plastically deformed region only extends to a maximum depth of ~4 µm and is contained within the first layer of grains adjacent to the groove surface. The 3D morphologies of the cracks under the scratches have been determined by FIB sectioning, showing that a high density of microcracks exists under the scratches in both ceramics. Differences between the plastic deformation and subsurface facture modes of the alumina and the nanocomposite are discussed. 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Subsurface damage; Transmission electron microscope (TEM); Focused ion beam (FIB); Dislocations; Microcracks; Alumina; Nanocomposite
1. Introduction In recent years, research has shown that Al2O3/SiC nanocomposites exhibit excellent contact erosion resistance and surface finishing properties. Davidge et al. [1] and Lawrence et al. [2] have reported significant improvement in resistance to erosive wear compared to polycrystalline alumina of similar grain size. A similar improvement in
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sliding wear resistance has also been reported by Rodriguez et al. [3]. Winn and Todd [4], Kara and Roberts [5], and Cock et al. [6] also found that Al2O3/SiC nanocomposite materials show faster polishing and give a smoother surface than polycrystalline alumina under identical polishing conditions. In previous research [7], we used a Hertzian indentation method to measure the surface residual stress produced by grinding in alumina and alumina 5%SiC nanocomposites. There was a much greater residual compressive stress in the ground surfaces of the nanocomposite (~650 MPa) compared with the stress with a similar grain size
1359-6454/03/$22.00 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. doi:10.1016/S1359-6454(02)00387-7
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polycrystalline alumina (~200 MPa) machined under the same conditions. Our recent cross-sectional TEM results [8,9] show that in all cases of machine grinding or polishing of alumina and the nanocomposites, there is a region of dislocation activity close to the surface. In the nanocomposites this region of dislocation activity extends further into the material with increasing diamond grit size, reaching to a depth of several grain diameters for the ground surfaces with 150 µm diamond grit. On alumina surfaces the plastic zone never extends more than a single grain diameter, and under the most severe grinding conditions during the residual deformation is mostly basal twinning. The estimated residual stress in the dislocated layer in both alumina and nanocomposite is ~1500 MPa [8]. Such a large compressive surface residual stress in the surface of machined nanocomposites was first proposed by Zhao et al. [10], who attributed the improvement in bend strength of the nanocomposite over alumina to this surface residual stress. There has been much research on fracture generation in brittle materials during scratching with the aim of understanding the material removal mechanisms in ceramic machining [11–16], as it is possible to control the load acting on the contacts and the geometry of the grooves. Contact forces due to scratching can cause a series of events that may include crushing, plastic flow, elastic recovery, generation of residual stress and radial cracking [17,18]. Plastic flow may be induced under the groove, to form an irreversible deformation zone [19]. The unloading process as the indenter moves off includes an important elastic recovery component, which has been supposed to result in many of the residual features near the contact tracks (i.e. lateral cracking, radial crack extension) as for quasi-static point indentations [19–21]. The 3D stress field around a scratch, however, is quite different to the stress field for an indentation in that it is asymmetric with a strong tensile stress near the surface behind the point contact [19]; fracture thus is expected to occur behind the point contact. The fracture pattern around scratches is highly complex. Microcracks and large radial or lateral cracks may all be present, and in polycrystalline alumina, the plastic deformation region near the
scratch involves the activation both of dislocations and of twins [11,12,14]. In this study we quantify the subsurface damage generated by single point scratching operation in alumina and Al2O3-5%vol.SiC nanocomposite using a new method of quantitative 3D focused ion beam (FIB) analysis to map the 3D crack distribution, and transmission electron microscopy (TEM) of cross-sectional specimens to map dislocation activity. The residual damage under the scratch is determined for both ceramics to assess the role of the nano-sized SiC particles.
2. Experimental procedures 2.1. Materials and specimen preparation The “nanocomposite” used in this study consists of an alumina matrix material containing 5 vol.% submicron SiC. The alumina powder used was AKP53 (Sumitomo, Japan) with submicron particle size and of reported chemical purity of 99.99% αAl2O3. The SiC particles were a commercial α-SiC powder UF 45 (Lonza, Germany), with a mean particle size of 90 nm. The materials fabrication process has been described in detail elsewhere [22,23], and is outlined briefly here. The SiC powder was first dispersed in distilled water using an ultrasonic probe. The dispersed SiC slurry was then mixed with the alumina powder, distilled water and a dispersing agent (Dispex A40, Allied Colloids, UK) in an attrition mill with zirconia milling media. After attrition milling, the homogeneous mixed slurry was freeze dried and passed through a 150 µm sieve. Pure alumina powder was prepared in the same way. All nanocomposites were hot pressed in a graphite die at 1650– 1680°C for 1 h under a pressure of 20–25 MPa in flowing argon. Alumina was hot pressed at 1500°C for 1 h under a pressure of 20 MPa, in order to achieve a reference material of grain size similar to that of the nanocomposite (3–4 µm). Both alumina and nanocomposite are fully densified under such processing conditions. The hot-pressed discs were ground to remove the top surface on both sides with an epoxy resin bonded diamond wheel (grit size 150 µm) to achi-
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eve a specimen thickness of about 2.5–3.0 mm. Specimens were cut directly from the ground disc and the ground surfaces were lapped using diamond grits in order 25, 8, 3 µm to a final 1 µm grit size. Each step of the polishing sequence was performed for long enough to eliminate all surface damage introduced by the previous step. The polishing behaviour of these materials is discussed in detail elsewhere [8]. 2.2. Single point scratch tests The single point scratching tests were carried out on 1 µm finished surfaces for both alumina and the nanocomposite. The specimen is traversed beneath a loaded conical diamond indenter (120° tip angle) by means of a motor-driven stage micrometer. The normal load used in this study was 1 N. The traverse speed of the specimen was 0.05 mm/s. Because of stiffness in the load-arm bearing on the machine, the accuracy of the applied load is approximately ±0.02 N. The diamond indenter was found to be substantially worn after a group of tests. To ensure that the alumina and nanocomposite were tested under the same conditions, the alumina and nanocomposite specimens were mounted together at the same horizontal level, and several scratches were made alternately on each specimen type at each progressively increasing load, using the same arm balance conditions throughout. 2.3. Fluorescence spectroscopy measurements of residual stress The residual stresses near to the scratch tracks in the alumina and the nanocomposite specimens were quantified using the technique of Cr3+/Al2O3 fluorescence spectroscopy [24]. For the Raman spectroscope machine used (Renishaw, UK), the laser probe has a spot size of about 2.5 µm. The shifts and broadening (full width at half maximum (FWHM) values) of the R1 and R2 fluorescence lines across the scratches were determined by fitting the raw data with GRAMS/32 (Galactic, US) using Gaussian and Lorentz functions. The shifts were measured relative to values determined from fracture surfaces, which we assume to be stressfree [25].
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2.4. Microstructural examination by TEM In order to quantify the subsurface plastic deformation, cross-section TEM specimens were prepared from the scratched nanocomposite and alumina specimens using the following procedure, as schematically shown in Fig. 1. The 1 µm finished alumina and nanocomposite specimens, with a width × length × thickness of about 2.5 × 5 × 1.3 mm, were aligned with their lengths parallel to the scratching direction. Six scratches were made on each surface by alternately scratching the alumina and nanocomposite samples in the same direction. The scratches on each specimen were separated laterally by about 250 µm. The scratched surfaces of alumina and nanocomposite were then bonded together with a thin layer of epoxy resin (M-Bond 610, Measurements Group Inc., USA). The two surfaces to be bonded were aligned so that scratches made immediately after one another on each surface faced each other across the bond. Pieces ~500 µm thick were cut from the glued body along the cross section normal to the epoxy bonded layer. These bonded pairs were then polished with 3 µm diamond slurry to 150–200 µm thickness. The four corners were removed by hand polishing using 8 µm diamond to make a roughly circular disc with a maximum diameter ⬍3 mm containing the epoxy-bonded interface across its middle. This was then bonded to a copper grid prior to thinning. A Gatan 656 (Gatan Inc., USA) dimple grinder was used to thin the central region of the specimen to a thickness of about 50 µm. Finally the dimpled region was thinned to electron transparency using a Gatan 600 ion beam miller with cooling stage. TEM examination was carried out using a JEM 200CX microscope and JEOL 4000EX. In order to prevent charging during observation, the specimens were coated with carbon prior to observation. 2.5. 3D mapping of subsurface cracks under scratches by FIB The 3D distribution of cracks generated under single point scratches on the surfaces of alumina and the nanocomposite was analysed using a new method of 3D FIB tomography which can be used
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Fig. 1. Schematic procedure for making cross section TEM foils of scratched Al2O3 and Al2O3/5vol.%SiC nanocomposite.
to determine the 3D morphology of cracks and grains with sub-micron resolution [26–28]. The analyses were carried out on a FEI FIB 200 microscope using the 3D mapping method discussed in detail elsewhere [26,27]. The method is outlined as follows: (a) First, a layer of gold was sputtered onto the sample surface to prevent charging during FIB processing. (b) A 1 µm thick platinum coating was deposited on the region of the scratch to be analysed, so as to prevent Ga+ implantation and sputter erosion of the top portion of the surface. (c) A rectangular trench was dug right across the scratch, adjacent to the platinum, using a 7000 pA 30 kV Ga+ ion beam. The trench was sufficiently large and deep such that the sample could be tilted and the scratch viewed in cross section, as schematically shown in Fig. 2(a). Positional reference marks were also drilled into one surface to enable subsequent 3D alignment of the 2D images. (d) A 2D cross section of the scratch track was then obtained by sputtering a clean surface perpendicular to the scratch using a 500 pA Ga+ ion beam.
(e) The 2D cross-sectioned surface was imaged down the direction used for sputtering using ion-induced secondary electrons (ISE) with a 50 pA Ga+ beam. The specimen was then tilted by 45°, and a second cross-sectional ISE image recorded, as shown schematically in Fig. 2(b). These two non-parallel images are sufficient to determine the 3D co-ordinates of features in the 2D cross-sectional plane. (f) A new 2D cross-section of the scratch track was obtained by sputtering a fresh surface parallel to the previous one displaced by 300–600 nm along the scratch axis. This new surface was then imaged from the top, and at q ⫽ 45° specimen tilt. (g) By iteratively repeating (d)–(f), that is sequentially drilling and imaging many parallel 2D cross-sections across the scratch track, a 3D data set of the microstructure is collected. The multiple 2D images are subsequently aligned by cross-correlation of reference markers, and analysed by computer [26–29]. The relative coordinate along the scratch axis (defined as the yaxis) for a given 2D x–z milled plane (z = Ga+ beam axis, x = sample rotation axis) is determined from the set of ISE images taken at q ⫽ 0°. The
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Fig. 2. Schematic FIB cross section imaging method. (a) Drilling deep trough and reference markers. (b) Tilting sample by q° around X-axis, and 2D imaging of cross-sectional plane x– y⬘ where y ⫽ y sin q.
ISE images taken at tilt angles q ⫽ 45° are stretched along the Ga+ beam z-axis by factor 1/sin(q) to compensate for the imaging angle, and aligned to enable the 3D co-ordinates of the crack and surface profiles to be extracted. All image processing was done with the voxel-based IDL image analysis package (Research Systems Inc, Boulder, CO, USA). In the aligned ISE images, the crack profiles are detected as dark lines due to reduced secondary electron yield compared to the surrounding matrix. Some cracks have partial bright edges due to
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enhanced secondary electron emission at edges (topographic contrast). The alumina surface in each ISE image is detectable as a line of bright contrast, due to the formation of a Schottky contact between the insulator and the overlying Au/Pt conducting metal. The crack and surface profiles were extracted from the surrounding background with ±2 pixel resolution by threshholding the contrast in the images (which can be enhanced using edge filters or manual edge contrast enhancement). Cracks could not be confused with other microstructural features such as grain boundaries and SiC particles since (i) no variation in the secondary electron yield is detectable between grains of different orientation (channeling contrast) due to the dominant voltage contrast contribution in these alumina samples [27], and (ii) the SiC particles gave no contrast in the ISE images obtained with low Ga+ exposures used for crack location. The errors in the 3D location of the crack and surface profiles are estimated as ⌬x ⫽ ± 2 pixels, ⌬z ⫽ ± 2√2 pixels in a given (x–z) 2D slice, plus ⌬x, ⌬y, ⌬z ⫽ ± 3 pixels error alignment in the positions of the 2D slices [27,28], giving minimum errors at 15K magnification of ⌬x ⫽ ± 115 nm, ⌬y ⫽ ± 70 nm and ⌬z ⫽ ± 135 nm. Bilinear interpolation was carried out to generate 3D reconstructions of the scratch surfaces. The positions of the alumina surface in the interpolated region between each slice have greater errors, with errors apparently equal to the distance from the adjacent 2D slice.
3. Experimental results 3.1. Surface morphology Typical morphologies of scratches in the alumina and the nanocomposites are shown in the optical micrographs in Fig. 3. No obvious difference in response to scratching, such as deformation pile-up, can be distinguished from the morphology of the scratches alone. Chipping occurs along both sides of the scratch tracks, appearing marginally more extensive along alumina scratches. Inside the grooves, both materials exhibit a smeared surface indicating possible yielding.
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Fig. 3. Optical morphology of single-point scratch tracks made with a 1 N normal load and a diamond cone indenter of 120° tip angle on surfaces of (a) Al2O3 and (b) Al2O3/5vol.%SiC nanocomposite.
3.2. Fluorescence spectroscopy measurements of residual stress Fig. 4 shows the measured R1 and R2 Cr3+ fluorescence line shifts (from reference fracture surfaces values) and line broadening (FWHM values) along tracks crossing the scratches in the alumina and the nanocomposite. The fluorescence shift can be related to an average stress level inside the probed volume using known piezo-spectroscopic coefficients [24,30]. The fluorescence broadening can result from inhomogeneous strains inside the probed volume [24]. This inhomogeneity can result from defects such as dislocations and cracks within the probed volume [25]. Fig. 4 shows that the fluorescence line shift from that for a stress-free surface is much greater in magnitude in the nanocomposite than in the alumina, and the shift extends laterally much wider. The shift for the alumina is so small that it is within the fluctuation range of the reference surface. This indicates that the scratches have generated a higher
average residual stress in the probed volume in the nanocomposite surface (biaxial compression of about 600 MPa) than in the alumina (biaxial compression of about 30 MPa). The stress field of the nanocomposite is also more extensive than that for the alumina; the shifts in frequency fall to zero at about 50–75 µm from the scratch track centre for the nanocomposite, but are confined to about 20 µm from the track for alumina. Fluorescence data from a wider range of scratched and indented surfaces are presented and discussed elsewhere [25]. Inside the groove, the fluorescence peaks for both materials were always broadened, i.e. a value of maximum FWHM to 18.27 cm⫺1 for R1 and 15.16 cm⫺1 for R2 in alumina and to 30.78 and 17.74 cm⫺1 in Al2O3-5vol.% SiC nanocomposite, compared to those from the respective reference fracture surfaces, where the average FWHM values are 15.74 ± 0.56 and 13.63 ± 0.63 cm⫺1 for R1 and R2 fluorescence lines in alumina and 15.21 ± 0.26 and 13.00 ± 0.26 cm⫺1 in the nanocomposite. The magnitude of broadening, however, is much larger in the nanocomposite than in alumina under the same probe conditions, suggesting a different residual defect distribution [25]. The defect distributions under the scratches were therefore investigated using FIB sectioning and crosssectional TEM. 3.3. Fracture pattern near the scratches Cracking damage under the single point scratch tracks was examined using 3D FIB cross sectioning. Fig. 5 shows 2D cross-sectional images and across scratch tracks in Al2O3 Al2O3/5vol.%SiC nanocomposite. At the start of ISE imaging, alumina images darkly compared to the overlaying protective metal due to dominant voltage contrast, and the SiC particles are not resolved. Cracks are resolved as dark lines, with partial bright edges due to enhanced secondary electron emission at edges (topographic contrast). With increased Ga+ exposure, differential surface modification (Ga+ implantation) reveals the SiC particles, as shown in Fig. 5(b). Repeated imaging or long exposures, however, may lead to some crack modification at the cross-section surface,
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Fig. 4. Shifts and FWHMs of fluorescence spectra across scratches made by 1 N normal load. (a-i), (b-i) peak shifts from stressfree fracture surfaces of Al2O3 and Al2O3/5vol.%SiC, respectively. (a-ii), (b-ii) FWHMs of Al2O3 and Al2O3/5vol.%SiC, respectively. The optical images of the probed scratches are superimposed on each chart at the same scale as the position axis.
such as preferential sputtering from the crack edges and sputtering into the finest cracks. Both alumina and the nanocomposite show residual damage from the scratching process with a fairly high density of micro-cracks under the scratch groove, even though there are very few cracks visible from the top surface (see Fig. 3). In order get a better quantification of the subsurface fracture as seen in Fig. 5, a series of parallel 2D sections across the scratches were cut and imaged with variable separations of 300–600 nm. Fig. 6 shows multiple parallel 2D cross sections through the scratch grooves, illustrating the non-uniform crack distribution in different 2D sections. These crack variations arise from the local variations in crystallography, grain boundary position and nanoparticle distribution. The 2D FIB images were aligned in 3D. Fig. 7 shows 3D reconstructions of the residual scratch
tracks, and the associated 3D crack zones viewed from inside the sample. The scratch tracks are shown as solid surfaces at the bottom of each image, and the cracks are shown in white on a black free space background. Microstructural regions (A)–(F) in Fig. 7 correspond to: (A) the sample surface adjacent to the scratch track; (B) the scratch edges where limited material pile-up and significant microcracking occurs; (C) the bottom of the residual scratch groove; (D) deep lateral cracks clusters propagating under the apparently undeformed surface (A); (E) deep radial cracks occurring under the scratch site; and (F) a zone where cracks are not resolved directly adjacent to scratch bottom (C). In both the alumina and nanocomposite the cracking extends ⬎7 µm below the scratch bottom (C) and ⬎10 µm from the central line of the scratch track. The subsurface cracking can be
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Fig. 5. FIB 2D cross-sections through scratch tracks made with a 1 N normal load. (a) Al2O3, (b) Al2O3/5vol.% nanocomposite. The arrows indicate subsurface cracks below the scratch site, and “CD” and “CA” indicate positions of crack deflection and crack arrest by SiC particles. The spots with dark contrast in (b) are SiC nanoparticles which become visible with long Ga+ exposure times. The viewing angle is 45° from the milled plane.
described as three zones ((I)–(III)) shown schematically in Fig. 8. Directly under the centre of the scratch tracks (which appear optically smooth) there is a zone (I) where no micro-cracks are resolved by the FIB analysis (regions (F) in Fig. 7(a) and (b)). Deeper under the scratch, cracks are initiated in a microcracking zone (II), which initially have very tortuous, interconnected branched morphologies. The microcracking zone (II) intersects the surface at the scratch edges (regions (B) in Fig. 7(a) and (b)), where optically the surface appears extremely rough and chipped. A handful of cracks extend over 7 µm below the scratch surface and over 10 µm from the scratch centre, as deep radial and lateral cracks in zone (III) (regions (D) and (E)) in Fig. 7(a) and (b)). The microcracking zone (II) extends downwards to link with the deep lateral and radial crack zone. Some of the lateral cracks extend under the outside of the residual scratch track, which optically
appears undeformed (regions (A) in Fig. 7(a) and (b)). The general 3D features of the crack zones described above appear very similar in the two analysed areas of the alumina and nanocomposite. As ISE imaging cannot resolve the grain boundaries, it is impossible to determine definitively which cracks are intragranular and which intergranular. However, the branched crack morphologies, and the fact that many cracks are separated by distances smaller than the average grain size of 3–4 µm, indicate that intragranular cracking occurs in both materials. Some cracks in the nanocomposite interact with SiC particles: (i) deflecting on a scale comparable to the size of nano-particle (see the cracks arrowed as “CD” in Fig. 5(b)); and (ii) exhibiting local crack-tip arrest by a relatively large particle (see particle arrowed as “CA” in Fig. 5(b)). Crack deflection on the micron scale occurs in both ceramics; this might be a result of interactions between cracks and grain boundaries. 3.4. Plastic deformation zone under the scratches Fig. 8 shows a TEM cross section containing two scratches, one in the alumina and one in the nanocomposite. By measuring the epoxy glue thickness in the cross section, it can be seen that in the vicinity of the scratches about 2 µm thick of material was removed from the original top surface at each side during TEM specimen preparation by ion milling. Even though the original scratch surface has not been preserved, the edge geometry (maximum scratch depth and width) is similar to the true profiles of the scratch surfaces as measured by FIB (Table 1), as expected if the material removal is nearly constant along the sample edge. A number of TEM cross sections show that beneath the scratches, a plastically deformed zone forms due to dislocation and twin activation in both the alumina and the nanocomposite. The measured maximum extent of the residual deformation from the centre of the groove surfaces was observed to be ~4 µm for the alumina and ~7 µm for the nanocomposite under the same scratching conditions. Residual plastic deformation observed under a scratch in alumina is shown in Fig. 9. Slip is acti-
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Fig. 6. Examples of multiple parallel 2D FIB cross-sections (spacing schematic) showing subsurface fracture damage under the scratch grooves in (a) Al2O3, (b) Al2O3/5vol.% nanocomposite. The viewing angle is 45° from the milled plane. The arrow indicates limited pile-up of material above the original surface height at the scratch groove edges.
vated only within the surface layer of grains. The residual dislocation density is extremely high, making it difficult to image individual dislocations and producing significant strain contrast. Some microcracks are observed in the TEM sections around the grain boundaries. In the nanocomposite, there is a much more extensive residual deformation zone, extending below the surface layer of grains. Similar to the alumina, the first layer of grains adjacent to the scratch exhibits such a high dislocation density and residual strain that it is almost impossible to characterise individual defects (grains G1 and G2 in Fig. 10(a)). In addition to dislocation activation, twinning was observed. These twins are confirmed by diffraction to be basal twins, with the (0001) planes oriented close to the normal loading direction of scratching. Twinning penetration might be partially limited by SiC nanoparticles. In grain G2
of Fig. 10(a), the growth front of the arrowed twin locally stops at the site of a SiC particle. Although the twinning partials will only be pinned at some points along their length, and can extend around the particles in 3D, it is likely that the pinning sites are effective in retarding the growth of the twins. There are about six twins visible in the upper part of the G2 grain, only one of which extends into the lower part of the G2 grain. The morphology of the basal twins is very high aspect ratio platelets, with (0001) habit plane and thickness in the [0001] direction of only about 20 nm measured using HREM. This platelet morphology is comparable to that of basal twins formed in sapphire [31]. Fig. 10(b) shows residual plastic deformation in the second and third layers of grains under the centre of the groove at depth of more than ~4 µm. The defect density decreases with increasing distance from the groove surface, and varies from
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surface grain containing a section of the plastic– elastic boundary for the scratch in Fig. 10. Adjacent grains have dislocation content indistinguishable from the bulk. Cracks were observed beneath the scratches in both the alumina and nanocomposite TEM samples. In alumina (Fig. 9), a crack is present close to the boundary between the first and second layer of grains. In the nanocomposite (Fig. 10), a large crack is present at the boundary between the first and second layer of grains, and ends at the limit of the residual plastic deformation zone in Fig. 11. It is not clear if all these cracks are formed during the scratch operation, or whether some at least arise and coalesce as a mechanism for stress relief during thinning of the TEM samples to electron transparency.
4. Discussion
Fig. 7. 3D reconstructions of the crack morphologies around a 1 N load scratch track (solid surface at the bottom of each image), viewed “inverted” from inside the sample. The scratch track is shown in grey, with the cracks white on a black free space background. (a) Polycrystalline alumina: reconstructed dimension 19.1 µm across scratch × 7.2 µm along scratch. (b) Polycrystalline Al2O3/5vol.% nanocomposite: reconstructed dimension 16.8 µm across scratch × 4.3 µm along scratch. Residual cracks around the scratch site penetrate more than 7 µm into sample from groove bottom in (a) and more than 10 µm in (b). Labels (A)–(F) denote microstructural regions discussed in the text.
grain to grain. The grains either side of the scratch groove also contain significant residual plastic deformation. For the scratch in Fig. 10, the maximum lateral extent of the residual deformation was ~6 µm from the groove centre in a plane 2 µm below the original surface. Fig. 11 shows a
In contrast to studies of ground, abraded, polished or eroded surfaces [12] where there are many overlapping adjacent contact events, these observations of single point scratches enable the plastic deformation and subsurface fracture patterns arising from a relatively simple contact geometry to be quantified. In contact with a sharp indenter, the elastic stress field is in principle singular about the indenter point, which inevitably leads to plastic deformation in most ceramics [20]. The activation of plastic deformation is controlled by the large components of shear and hydrostatic compression in the contact field. The TEM and FIB cross-sectional examinations of single point scratches in alumina and the nanocomposite show that the scratch operation results in dislocation activation, twin nucleation and microcracking. At the load used here (1 N), in the alumina, plastic deformation is only observed within the first layer here of grains contacted by the moving indenter (i.e. to a maximum depth of ~3 µm). The dominant plastic deformation modes in alumina found by other researchers are basal twinning and 1/3⬍112¯ 0⬎ dislocations gliding on basal and non-basal slip planes [31–36]. Some ⬍10 1¯ 0⬎ and pyramidal Burgers vectors may also be activated [31,34,35,37]. The defects appear to orig-
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Fig. 8. Cross-sectional TEM of scratches in alumina and nanocomposite. Dashed lines indicate the position of the un-scratched surface plane in alumina and nanocomposite, which was assessed by measuring the thickness of the epoxy filling gap near this regime. Table 1 Measured scratch depth and width from FIB cross-section image Material
Scratch depth (µm)
Scratch width (µm)
Alumina Nanocomposite
1.6 1.6
10.4 11.5
inate from the contact zone [31,33,35,36]. Inkson [31] found that under prismatic and basal surfaces of single crystal alumina abraded with 30 µm diamond grit, dislocation penetration was to about 5 µm below the surface. Penetration of basal twinning was found to have a crystallographic dependency [31,34–36], with penetration into prismatic surfaces occurring to ⬎10 µm [31]. In the polycry-
stalline alumina examined here, the average grain size is about 3–4 µm. The glide of dislocations and basal twins appears to have been stopped by the grain boundaries. In the nanocomposite, dislocation and twin activation was observed to a rather greater extent than in alumina, penetrating to below the surface layer of grains (Fig. 10(b)) to a depth of about 10 µm. This may be due to (a) defect propagation across the grain boundaries, and (b) activation of subsurface dislocation sources. Considering possible subsurface dislocation sources, dislocation nuclei might be formed by local thermal residual stress around the SiC particles due to the difference in thermal expansion coefficients of alumina (~8.3 × 10⫺6 K⫺1) and silicon carbide (~5.1 × 10⫺6 K⫺1). During cooling
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Fig. 9. Plastic deformation beneath the scratch groove in monolithic Al2O3.
from the processing temperature (1650–1680°C), residual stresses will build up to a level of near 2 GPa near the Al2O3/SiC interface (SiC particle in compression, Al2O3 matrix in tension) in the room temperature [38,39]; at intermediate temperatures (about 400–1100°C), dislocations with basal and non-basal Burgers vectors may be activated [40– 42]. In addition, dislocations may be activated at grain boundaries which have not undergone stress relief by cracking. Under the contact stress field, these dislocations may glide further and/or act as sources for further dislocation generation. The 3D FIB mapping of crack morphologies around the residual scratch resolves microcracks under the scratch track in both alumina and the nanocomposite (Fig. 7). We assume that the act of milling a hole across the scratch track does not change the local stress field sufficiently to cause marked alteration in the original crack profiles, for example by crack propagation towards the new free surface. The stress relief in the FIB sectioning, with one free surface and surrounding constraining material, should be much less than that in a TEM foil. The 3D crack maps exhibit very similar crack morphologies for both materials, which may be broadly described as consisting of three zones, as shown in Fig. 12:
Fig. 10. Plastic deformation in the grains beneath the centre of the scratch groove in Al2O3/5vol%SiC nanocomposite. Two adjacent grains (G1 and G2) have plastically deformed by the activation of dislocations in grain G1 and a mixture of dislocation and basal twin in grain G2 where selected area diffraction shows basal twins; (b) slip originating from a grain boundary in a grain under grain G2; grain (G3) under grain G1 is heavily strained with a high density of dislocations.
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Fig. 11. Plastic deformation features to one side of the scratch. The plastic/elastic boundary can be seen near the centre of the micrograph.
Fig. 12. Schematic diagram of types of subsurface cracking observed beneath scratches. The cracked volume is composed of three zones: zone I (crack deficient), zone II (microcracking) and zone III (deep lateral and median cracking).
(I)
A crack deficient zone directly under the centre of the scratch groove. (II) A microcracking zone below and around zone (I), which extends up to the surface at the two groove edges. (III) A zone below zones (I) and (II), with deep lateral and median cracks which do not intersect the surface. The maximum depth of cracking (⬎7 µm) is comparable to the depth of the dislocation zone (~7 µm) in the nanocomposite, but almost twice the depth of the dislocation zone (~4 µm) in the alumina. Considering first the crack deficient zone (I) directly under the centre of the scratch groove, this zone has a very high residual dislocation density in both the nanocomposite and the alumina. This
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zone will have a high compressive residual stress after the indenter moves off [19]. This appears to inhibit both the nucleation of cracks and the propagation of subsurface cracks up to the groove surface. With no crack intersection and associated material chipping, the groove bottom appears optically smooth after scratching in both materials. This contrasts with grinding and polishing, where nanocomposites may exhibit much better surface qualities than pure aluminas [5,6,8,9]. Surrounding the compressive zone, microcracking occurs (zone (II)). The microcracking extends up to the surface at the groove edges, where material removal by crack linking and chipping occurs. Microcracking is expected to be observed where the stress field starts to be dominated by tensile components [43]. Close to the compressive zone, the microcracks have a heavily branched morphology, and in the sampled regions there appeared to be a marginally higher density of microcracks in the nanocomposite than in the alumina. In both the alumina and nanocomposite, the microcracking zone (II) then evolves into a zone (III) of deep lateral and median cracks. The nucleation and propagation of cracks within the tensile stress field in both the alumina and nanocomposite appears to be comparable, since no major difference in the residual crack morphologies was observed in this study. This is consistent with measurements of the fracture toughness Klc for the two materials, which are found to be similar [7,10,22,23], During the scratching process, the activation of dislocations will change the local stress concentration distribution around the tip. In the alumina the dislocations are activated within the top layer of grains, which corresponds to the crack-free compressive zone (I) and the microcracking zone (II). In the nanocomposite, dislocations are also activated in zone III below the first layer of grains where deep lateral and median cracks are observed. Dislocation activation may affect crack nucleation and propagation by: (i)
Causing directional stress concentrations due to difficulties in initiation of enough independent slip systems to accommodate shape change (for example, dislocations with non-
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(ii)
(iii)
(iv)
(v)
(vi)
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basal Burgers vector components), leading to the nucleation of cracks on grain boundaries or inside grains [31,33,35–37]. Generating basal twins, the (0001) interfaces of which are preferential crack paths [31]. Microcracking initiated by the intersection of rhombohedral twins as found by Heuer [44] was not observed in these specimens. Causing stress concentrations at dislocation pile-ups within grains and at grain boundaries, which may encourage microcrack nucleation [31,34,37]. In the nanocomposite, the distribution of SiC nanoparticles may additionally affect crack nucleation and propagation by: Providing subsurface dislocation sources, generated during cooling to accommodate thermal misfit as discussed above. Directly causing spontaneous crack initiation at the SiC–Al2O3 interfaces to relieve thermal mismatch. Some large SiC particles or agglomorates may in principle act as crack initiation sites; however, no such cracks were observed around SiC particles in the bulk material away from the scratch zone. Causing crack deflection or pinning at nanoparticles. FIB imaging shows the deflection of some cracks at SiC particles⬎30 nm in size (smaller particles are below the resolution limit). The observation of cracks terminating/beginning at SiC particles in the 2D cross sections was rare and only observed at the larger particles. With only 5 vol.% SiC, these pinning events may simply be coincidental juxtapositions. This is different from the frequent and property-dominating crack/second phase interactions in large particle, high vol.% ceramic composites [45].
The difference in depth of dislocation activation between the alumina and nanocomposite, under the present loading conditions, appears not to alter significantly crack initiation and propagation by mechanisms (i)–(vi) under the loading conditions used. Both materials have similar subsurface crack morphologies and a compressive stress zone inhibiting crack propagation to the surface at the
centre of the groove. This results in the scratch track morphologies appearing very similar. The measurements of residual stress across the scratches in alumina and the nanocomposite by Cr3+/Al2O3 fluorescence spectroscopy are consistent with the microstructural observations by TEM and FIB. The much larger broadening of the fluorescence lines observed in the nanocomposite arises from a greater average density of dislocations inside the probed volume, due to the greater depth of the dislocated layer [25]. Stress inhomogeneities due to microcracking inside the groove will also contribute to the broadening [25], but the observed similar microcracking patterns under the grooves in alumina and the nanocomposite will give rise to similar broadening contributions. The fluorescence shift from the value for a stress-free surface (i.e. a fracture surface of the same ceramic) represents an average stress level inside the probed volume. The shift is much greater in the nanocomposite than in alumina (Fig. 4); this can also be attributed to the more extensive residual dislocation zone in the nanocomposite than in the alumina, giving rise to greater and more extensive elastic stresses outside this sub-scratch plastic zone, and is consistent with measurements of surface residual stresses on machine ground surfaces [7]. 5. Summary 1. Scratch tracks made with a 1 N normal load in the alumina and an alumina/5% silicon carbide nanocomposite appear similar by optical microscopy. 2. Cross-sectional TEM shows that dislocation activation in the alumina was in the first layer of grains only (~3 µm depth), and in the nanocomposite, penetrated to below the first layer (to ~10 µm). 3. FIB sectioning shows that both alumina and the nanocomposite exhibit similar 3D crack patterns around the scratches: (I) a crack-free compressive zone under the smooth groove base (~3 µm), (II) a microcrack zone below and around zone (I) (~3 µm), (III) a deep lateral and median crack zone below (I) and (II). Microcracks extend to the surface at the groove edges and produce chipping at the groove edges.
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Residual stress levels (measured by Cr+ fluorescence) across the scratches are higher and more extensive in the nanocomposite. Line broading near the scratch is also greater in the nanocomposite. This is consistent with the TEM observations of a larger plastic zone in the nanocomposite.
Acknowledgements We gratefully acknowledge the support of the EPSRC under grant GR/L95908 (H.Z.W.), and of the Royal Society (B.J.I.). We thank Prof R. Young at the Manchester Materials Centre for providing facilities for fluorescence measurements, and Prof. B. Derby (also at MMC) for helpful discussions.
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[17] Imanaka O, Fujino S, Mineta S. The science of ceramic machining and surface finish, NBS Special Publication 348; 1972. p. 7 [18] Larchuk TJ, Conway Jr JC, Kirchner HP. J. Am. Ceram. Soc. 1985;68:209. [19] Swain MV. Proc. R. Soc.(Lond.) 1979;A336:575. [20] Lawn BR, Howes VR. J. Mater. Sci. 1981;16:2745. [21] Marshall DB, Lawn BR, Evans AG. J. Am. Ceram. Soc. 1982;65:561. [22] Sternitzke M, Dupas E, Twigg P, Derby B. Acta mater. 1997;45:3963. [23] Carroll L, Sternitzke M, Derby B. Acta metall. mater. 1996;44:4543. [24] Ma Q, Clarke DR. J. Am. Ceram. Soc. 1993;76:1433. [25] Wu HZ, Roberts SG, Derby B, in preparation. [26] Inkson BJ, Steer TJ, Mo¨ bus G, Wagner T. J. Microsc. 2001;201:256. [27] Inkson BJ, Wu HZ, Steer TJ, Mo¨ bus G. Fundamentals of Nanoindentation & Nanotribology, MRS Proc. 2001;649(7):7. [28] Inkson BJ, Mulvihill M, Mo¨ bus G. Scripta mater. 2001;45:753. [29] Prewett PD, Mair GLR. Focused ion beams from liquid metal sources. Tauntory, UK: Research Studies Press Ltd, 1991. [30] Munro RG, Piermarini T, Block S, Holzapfel WB. J. Appl. Phys. 1985;57:165. [31] Inkson BJ. Acta mater. 2000;48:1883. [32] Lagerlof KPD, Heuer AH, Castaing J, Riviere JP, Mitchell TE. J. Am Ceram. Soc. 1994;77:385. [33] Conrad H. J. Am. Ceram. Soc. 1965;48:195. [34] Hockey BJ. J. Am. Ceram. Soc. 1971;54:223. [35] Hockey BJ. In: Bradt RC, Tressler RE, editors. Deformation of ceramic materials. New York: Plenum Press; 1975. p. 167–79. [36] Chan HM, Lawn BR. J. Am. Ceram. Soc. 1988;71:29. [37] Barceinas-Sanchez JDO, Rainforth WM. Acta mater. 1998;46:6475. [38] Todd RI, Bourke MAM, Borsa CE, Brook RJ. Acta mater. 1997;45:1791. [39] Levin I, Kaplan WD, Brandon DG. Acta metall. mater. 1994;42:1147. [40] Cadoz J, Riviere JP, Castaing J. In: Tressler RE, Bradt RC, editors. Deformation of ceramic materials II. New York: Plenum Press; 1983. [41] Castaing J, Cadoz J, Kirby SH. J. Am. Ceram. Soc. 1981;64:504. [42] Jiao SL. D. Phil. thesis, Oxford University; 1997. [43] Ahn Y, Farris TN, Chandrasekar S. Mech. Mater. 1998;29:143. [44] Heuer AH. Phil. Mag. A 1966;13:379. [45] Faber KT, Evans AG. Acta metall. 1983;31:565.
Subsurface damage analysis by TEM and 3D FIB crack mapping in alumina and alumina/5vol.%SiC nanocomposites H.Z. Wu ∗, S.G. Roberts, G. Mo¨bus, B.J. Inkson Department of Materials, University of Oxford, Oxford OX1 3PH, UK Received 23 October 2001; received in revised form 17 July 2002; accepted 17 July 2002
Abstract TEM and 3D crack analysis by focused ion beam (FIB) cross sectioning have been used to quantify the subsurface damage beneath scratches made by a 120° cone indenter loaded to 1 N in monolithic polycrystalline alumina and alumina/5vol.%SiC nanocomposites. In the nanocomposite, an extensive plastic deformation zone was found under the scratch grooves, extending beyond the first layer of grains to a maximum plastic deformation depth of ~7 µm below the surface of the track. In the alumina, however, the plastically deformed region only extends to a maximum depth of ~4 µm and is contained within the first layer of grains adjacent to the groove surface. The 3D morphologies of the cracks under the scratches have been determined by FIB sectioning, showing that a high density of microcracks exists under the scratches in both ceramics. Differences between the plastic deformation and subsurface facture modes of the alumina and the nanocomposite are discussed. 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Subsurface damage; Transmission electron microscope (TEM); Focused ion beam (FIB); Dislocations; Microcracks; Alumina; Nanocomposite
1. Introduction In recent years, research has shown that Al2O3/SiC nanocomposites exhibit excellent contact erosion resistance and surface finishing properties. Davidge et al. [1] and Lawrence et al. [2] have reported significant improvement in resistance to erosive wear compared to polycrystalline alumina of similar grain size. A similar improvement in
Corresponding author. Fax: +44 1865 273768. E-mail address: [email protected] (H.Z. Wu). ∗
sliding wear resistance has also been reported by Rodriguez et al. [3]. Winn and Todd [4], Kara and Roberts [5], and Cock et al. [6] also found that Al2O3/SiC nanocomposite materials show faster polishing and give a smoother surface than polycrystalline alumina under identical polishing conditions. In previous research [7], we used a Hertzian indentation method to measure the surface residual stress produced by grinding in alumina and alumina 5%SiC nanocomposites. There was a much greater residual compressive stress in the ground surfaces of the nanocomposite (~650 MPa) compared with the stress with a similar grain size
1359-6454/03/$22.00 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. doi:10.1016/S1359-6454(02)00387-7
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polycrystalline alumina (~200 MPa) machined under the same conditions. Our recent cross-sectional TEM results [8,9] show that in all cases of machine grinding or polishing of alumina and the nanocomposites, there is a region of dislocation activity close to the surface. In the nanocomposites this region of dislocation activity extends further into the material with increasing diamond grit size, reaching to a depth of several grain diameters for the ground surfaces with 150 µm diamond grit. On alumina surfaces the plastic zone never extends more than a single grain diameter, and under the most severe grinding conditions during the residual deformation is mostly basal twinning. The estimated residual stress in the dislocated layer in both alumina and nanocomposite is ~1500 MPa [8]. Such a large compressive surface residual stress in the surface of machined nanocomposites was first proposed by Zhao et al. [10], who attributed the improvement in bend strength of the nanocomposite over alumina to this surface residual stress. There has been much research on fracture generation in brittle materials during scratching with the aim of understanding the material removal mechanisms in ceramic machining [11–16], as it is possible to control the load acting on the contacts and the geometry of the grooves. Contact forces due to scratching can cause a series of events that may include crushing, plastic flow, elastic recovery, generation of residual stress and radial cracking [17,18]. Plastic flow may be induced under the groove, to form an irreversible deformation zone [19]. The unloading process as the indenter moves off includes an important elastic recovery component, which has been supposed to result in many of the residual features near the contact tracks (i.e. lateral cracking, radial crack extension) as for quasi-static point indentations [19–21]. The 3D stress field around a scratch, however, is quite different to the stress field for an indentation in that it is asymmetric with a strong tensile stress near the surface behind the point contact [19]; fracture thus is expected to occur behind the point contact. The fracture pattern around scratches is highly complex. Microcracks and large radial or lateral cracks may all be present, and in polycrystalline alumina, the plastic deformation region near the
scratch involves the activation both of dislocations and of twins [11,12,14]. In this study we quantify the subsurface damage generated by single point scratching operation in alumina and Al2O3-5%vol.SiC nanocomposite using a new method of quantitative 3D focused ion beam (FIB) analysis to map the 3D crack distribution, and transmission electron microscopy (TEM) of cross-sectional specimens to map dislocation activity. The residual damage under the scratch is determined for both ceramics to assess the role of the nano-sized SiC particles.
2. Experimental procedures 2.1. Materials and specimen preparation The “nanocomposite” used in this study consists of an alumina matrix material containing 5 vol.% submicron SiC. The alumina powder used was AKP53 (Sumitomo, Japan) with submicron particle size and of reported chemical purity of 99.99% αAl2O3. The SiC particles were a commercial α-SiC powder UF 45 (Lonza, Germany), with a mean particle size of 90 nm. The materials fabrication process has been described in detail elsewhere [22,23], and is outlined briefly here. The SiC powder was first dispersed in distilled water using an ultrasonic probe. The dispersed SiC slurry was then mixed with the alumina powder, distilled water and a dispersing agent (Dispex A40, Allied Colloids, UK) in an attrition mill with zirconia milling media. After attrition milling, the homogeneous mixed slurry was freeze dried and passed through a 150 µm sieve. Pure alumina powder was prepared in the same way. All nanocomposites were hot pressed in a graphite die at 1650– 1680°C for 1 h under a pressure of 20–25 MPa in flowing argon. Alumina was hot pressed at 1500°C for 1 h under a pressure of 20 MPa, in order to achieve a reference material of grain size similar to that of the nanocomposite (3–4 µm). Both alumina and nanocomposite are fully densified under such processing conditions. The hot-pressed discs were ground to remove the top surface on both sides with an epoxy resin bonded diamond wheel (grit size 150 µm) to achi-
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eve a specimen thickness of about 2.5–3.0 mm. Specimens were cut directly from the ground disc and the ground surfaces were lapped using diamond grits in order 25, 8, 3 µm to a final 1 µm grit size. Each step of the polishing sequence was performed for long enough to eliminate all surface damage introduced by the previous step. The polishing behaviour of these materials is discussed in detail elsewhere [8]. 2.2. Single point scratch tests The single point scratching tests were carried out on 1 µm finished surfaces for both alumina and the nanocomposite. The specimen is traversed beneath a loaded conical diamond indenter (120° tip angle) by means of a motor-driven stage micrometer. The normal load used in this study was 1 N. The traverse speed of the specimen was 0.05 mm/s. Because of stiffness in the load-arm bearing on the machine, the accuracy of the applied load is approximately ±0.02 N. The diamond indenter was found to be substantially worn after a group of tests. To ensure that the alumina and nanocomposite were tested under the same conditions, the alumina and nanocomposite specimens were mounted together at the same horizontal level, and several scratches were made alternately on each specimen type at each progressively increasing load, using the same arm balance conditions throughout. 2.3. Fluorescence spectroscopy measurements of residual stress The residual stresses near to the scratch tracks in the alumina and the nanocomposite specimens were quantified using the technique of Cr3+/Al2O3 fluorescence spectroscopy [24]. For the Raman spectroscope machine used (Renishaw, UK), the laser probe has a spot size of about 2.5 µm. The shifts and broadening (full width at half maximum (FWHM) values) of the R1 and R2 fluorescence lines across the scratches were determined by fitting the raw data with GRAMS/32 (Galactic, US) using Gaussian and Lorentz functions. The shifts were measured relative to values determined from fracture surfaces, which we assume to be stressfree [25].
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2.4. Microstructural examination by TEM In order to quantify the subsurface plastic deformation, cross-section TEM specimens were prepared from the scratched nanocomposite and alumina specimens using the following procedure, as schematically shown in Fig. 1. The 1 µm finished alumina and nanocomposite specimens, with a width × length × thickness of about 2.5 × 5 × 1.3 mm, were aligned with their lengths parallel to the scratching direction. Six scratches were made on each surface by alternately scratching the alumina and nanocomposite samples in the same direction. The scratches on each specimen were separated laterally by about 250 µm. The scratched surfaces of alumina and nanocomposite were then bonded together with a thin layer of epoxy resin (M-Bond 610, Measurements Group Inc., USA). The two surfaces to be bonded were aligned so that scratches made immediately after one another on each surface faced each other across the bond. Pieces ~500 µm thick were cut from the glued body along the cross section normal to the epoxy bonded layer. These bonded pairs were then polished with 3 µm diamond slurry to 150–200 µm thickness. The four corners were removed by hand polishing using 8 µm diamond to make a roughly circular disc with a maximum diameter ⬍3 mm containing the epoxy-bonded interface across its middle. This was then bonded to a copper grid prior to thinning. A Gatan 656 (Gatan Inc., USA) dimple grinder was used to thin the central region of the specimen to a thickness of about 50 µm. Finally the dimpled region was thinned to electron transparency using a Gatan 600 ion beam miller with cooling stage. TEM examination was carried out using a JEM 200CX microscope and JEOL 4000EX. In order to prevent charging during observation, the specimens were coated with carbon prior to observation. 2.5. 3D mapping of subsurface cracks under scratches by FIB The 3D distribution of cracks generated under single point scratches on the surfaces of alumina and the nanocomposite was analysed using a new method of 3D FIB tomography which can be used
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Fig. 1. Schematic procedure for making cross section TEM foils of scratched Al2O3 and Al2O3/5vol.%SiC nanocomposite.
to determine the 3D morphology of cracks and grains with sub-micron resolution [26–28]. The analyses were carried out on a FEI FIB 200 microscope using the 3D mapping method discussed in detail elsewhere [26,27]. The method is outlined as follows: (a) First, a layer of gold was sputtered onto the sample surface to prevent charging during FIB processing. (b) A 1 µm thick platinum coating was deposited on the region of the scratch to be analysed, so as to prevent Ga+ implantation and sputter erosion of the top portion of the surface. (c) A rectangular trench was dug right across the scratch, adjacent to the platinum, using a 7000 pA 30 kV Ga+ ion beam. The trench was sufficiently large and deep such that the sample could be tilted and the scratch viewed in cross section, as schematically shown in Fig. 2(a). Positional reference marks were also drilled into one surface to enable subsequent 3D alignment of the 2D images. (d) A 2D cross section of the scratch track was then obtained by sputtering a clean surface perpendicular to the scratch using a 500 pA Ga+ ion beam.
(e) The 2D cross-sectioned surface was imaged down the direction used for sputtering using ion-induced secondary electrons (ISE) with a 50 pA Ga+ beam. The specimen was then tilted by 45°, and a second cross-sectional ISE image recorded, as shown schematically in Fig. 2(b). These two non-parallel images are sufficient to determine the 3D co-ordinates of features in the 2D cross-sectional plane. (f) A new 2D cross-section of the scratch track was obtained by sputtering a fresh surface parallel to the previous one displaced by 300–600 nm along the scratch axis. This new surface was then imaged from the top, and at q ⫽ 45° specimen tilt. (g) By iteratively repeating (d)–(f), that is sequentially drilling and imaging many parallel 2D cross-sections across the scratch track, a 3D data set of the microstructure is collected. The multiple 2D images are subsequently aligned by cross-correlation of reference markers, and analysed by computer [26–29]. The relative coordinate along the scratch axis (defined as the yaxis) for a given 2D x–z milled plane (z = Ga+ beam axis, x = sample rotation axis) is determined from the set of ISE images taken at q ⫽ 0°. The
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Fig. 2. Schematic FIB cross section imaging method. (a) Drilling deep trough and reference markers. (b) Tilting sample by q° around X-axis, and 2D imaging of cross-sectional plane x– y⬘ where y ⫽ y sin q.
ISE images taken at tilt angles q ⫽ 45° are stretched along the Ga+ beam z-axis by factor 1/sin(q) to compensate for the imaging angle, and aligned to enable the 3D co-ordinates of the crack and surface profiles to be extracted. All image processing was done with the voxel-based IDL image analysis package (Research Systems Inc, Boulder, CO, USA). In the aligned ISE images, the crack profiles are detected as dark lines due to reduced secondary electron yield compared to the surrounding matrix. Some cracks have partial bright edges due to
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enhanced secondary electron emission at edges (topographic contrast). The alumina surface in each ISE image is detectable as a line of bright contrast, due to the formation of a Schottky contact between the insulator and the overlying Au/Pt conducting metal. The crack and surface profiles were extracted from the surrounding background with ±2 pixel resolution by threshholding the contrast in the images (which can be enhanced using edge filters or manual edge contrast enhancement). Cracks could not be confused with other microstructural features such as grain boundaries and SiC particles since (i) no variation in the secondary electron yield is detectable between grains of different orientation (channeling contrast) due to the dominant voltage contrast contribution in these alumina samples [27], and (ii) the SiC particles gave no contrast in the ISE images obtained with low Ga+ exposures used for crack location. The errors in the 3D location of the crack and surface profiles are estimated as ⌬x ⫽ ± 2 pixels, ⌬z ⫽ ± 2√2 pixels in a given (x–z) 2D slice, plus ⌬x, ⌬y, ⌬z ⫽ ± 3 pixels error alignment in the positions of the 2D slices [27,28], giving minimum errors at 15K magnification of ⌬x ⫽ ± 115 nm, ⌬y ⫽ ± 70 nm and ⌬z ⫽ ± 135 nm. Bilinear interpolation was carried out to generate 3D reconstructions of the scratch surfaces. The positions of the alumina surface in the interpolated region between each slice have greater errors, with errors apparently equal to the distance from the adjacent 2D slice.
3. Experimental results 3.1. Surface morphology Typical morphologies of scratches in the alumina and the nanocomposites are shown in the optical micrographs in Fig. 3. No obvious difference in response to scratching, such as deformation pile-up, can be distinguished from the morphology of the scratches alone. Chipping occurs along both sides of the scratch tracks, appearing marginally more extensive along alumina scratches. Inside the grooves, both materials exhibit a smeared surface indicating possible yielding.
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Fig. 3. Optical morphology of single-point scratch tracks made with a 1 N normal load and a diamond cone indenter of 120° tip angle on surfaces of (a) Al2O3 and (b) Al2O3/5vol.%SiC nanocomposite.
3.2. Fluorescence spectroscopy measurements of residual stress Fig. 4 shows the measured R1 and R2 Cr3+ fluorescence line shifts (from reference fracture surfaces values) and line broadening (FWHM values) along tracks crossing the scratches in the alumina and the nanocomposite. The fluorescence shift can be related to an average stress level inside the probed volume using known piezo-spectroscopic coefficients [24,30]. The fluorescence broadening can result from inhomogeneous strains inside the probed volume [24]. This inhomogeneity can result from defects such as dislocations and cracks within the probed volume [25]. Fig. 4 shows that the fluorescence line shift from that for a stress-free surface is much greater in magnitude in the nanocomposite than in the alumina, and the shift extends laterally much wider. The shift for the alumina is so small that it is within the fluctuation range of the reference surface. This indicates that the scratches have generated a higher
average residual stress in the probed volume in the nanocomposite surface (biaxial compression of about 600 MPa) than in the alumina (biaxial compression of about 30 MPa). The stress field of the nanocomposite is also more extensive than that for the alumina; the shifts in frequency fall to zero at about 50–75 µm from the scratch track centre for the nanocomposite, but are confined to about 20 µm from the track for alumina. Fluorescence data from a wider range of scratched and indented surfaces are presented and discussed elsewhere [25]. Inside the groove, the fluorescence peaks for both materials were always broadened, i.e. a value of maximum FWHM to 18.27 cm⫺1 for R1 and 15.16 cm⫺1 for R2 in alumina and to 30.78 and 17.74 cm⫺1 in Al2O3-5vol.% SiC nanocomposite, compared to those from the respective reference fracture surfaces, where the average FWHM values are 15.74 ± 0.56 and 13.63 ± 0.63 cm⫺1 for R1 and R2 fluorescence lines in alumina and 15.21 ± 0.26 and 13.00 ± 0.26 cm⫺1 in the nanocomposite. The magnitude of broadening, however, is much larger in the nanocomposite than in alumina under the same probe conditions, suggesting a different residual defect distribution [25]. The defect distributions under the scratches were therefore investigated using FIB sectioning and crosssectional TEM. 3.3. Fracture pattern near the scratches Cracking damage under the single point scratch tracks was examined using 3D FIB cross sectioning. Fig. 5 shows 2D cross-sectional images and across scratch tracks in Al2O3 Al2O3/5vol.%SiC nanocomposite. At the start of ISE imaging, alumina images darkly compared to the overlaying protective metal due to dominant voltage contrast, and the SiC particles are not resolved. Cracks are resolved as dark lines, with partial bright edges due to enhanced secondary electron emission at edges (topographic contrast). With increased Ga+ exposure, differential surface modification (Ga+ implantation) reveals the SiC particles, as shown in Fig. 5(b). Repeated imaging or long exposures, however, may lead to some crack modification at the cross-section surface,
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Fig. 4. Shifts and FWHMs of fluorescence spectra across scratches made by 1 N normal load. (a-i), (b-i) peak shifts from stressfree fracture surfaces of Al2O3 and Al2O3/5vol.%SiC, respectively. (a-ii), (b-ii) FWHMs of Al2O3 and Al2O3/5vol.%SiC, respectively. The optical images of the probed scratches are superimposed on each chart at the same scale as the position axis.
such as preferential sputtering from the crack edges and sputtering into the finest cracks. Both alumina and the nanocomposite show residual damage from the scratching process with a fairly high density of micro-cracks under the scratch groove, even though there are very few cracks visible from the top surface (see Fig. 3). In order get a better quantification of the subsurface fracture as seen in Fig. 5, a series of parallel 2D sections across the scratches were cut and imaged with variable separations of 300–600 nm. Fig. 6 shows multiple parallel 2D cross sections through the scratch grooves, illustrating the non-uniform crack distribution in different 2D sections. These crack variations arise from the local variations in crystallography, grain boundary position and nanoparticle distribution. The 2D FIB images were aligned in 3D. Fig. 7 shows 3D reconstructions of the residual scratch
tracks, and the associated 3D crack zones viewed from inside the sample. The scratch tracks are shown as solid surfaces at the bottom of each image, and the cracks are shown in white on a black free space background. Microstructural regions (A)–(F) in Fig. 7 correspond to: (A) the sample surface adjacent to the scratch track; (B) the scratch edges where limited material pile-up and significant microcracking occurs; (C) the bottom of the residual scratch groove; (D) deep lateral cracks clusters propagating under the apparently undeformed surface (A); (E) deep radial cracks occurring under the scratch site; and (F) a zone where cracks are not resolved directly adjacent to scratch bottom (C). In both the alumina and nanocomposite the cracking extends ⬎7 µm below the scratch bottom (C) and ⬎10 µm from the central line of the scratch track. The subsurface cracking can be
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Fig. 5. FIB 2D cross-sections through scratch tracks made with a 1 N normal load. (a) Al2O3, (b) Al2O3/5vol.% nanocomposite. The arrows indicate subsurface cracks below the scratch site, and “CD” and “CA” indicate positions of crack deflection and crack arrest by SiC particles. The spots with dark contrast in (b) are SiC nanoparticles which become visible with long Ga+ exposure times. The viewing angle is 45° from the milled plane.
described as three zones ((I)–(III)) shown schematically in Fig. 8. Directly under the centre of the scratch tracks (which appear optically smooth) there is a zone (I) where no micro-cracks are resolved by the FIB analysis (regions (F) in Fig. 7(a) and (b)). Deeper under the scratch, cracks are initiated in a microcracking zone (II), which initially have very tortuous, interconnected branched morphologies. The microcracking zone (II) intersects the surface at the scratch edges (regions (B) in Fig. 7(a) and (b)), where optically the surface appears extremely rough and chipped. A handful of cracks extend over 7 µm below the scratch surface and over 10 µm from the scratch centre, as deep radial and lateral cracks in zone (III) (regions (D) and (E)) in Fig. 7(a) and (b)). The microcracking zone (II) extends downwards to link with the deep lateral and radial crack zone. Some of the lateral cracks extend under the outside of the residual scratch track, which optically
appears undeformed (regions (A) in Fig. 7(a) and (b)). The general 3D features of the crack zones described above appear very similar in the two analysed areas of the alumina and nanocomposite. As ISE imaging cannot resolve the grain boundaries, it is impossible to determine definitively which cracks are intragranular and which intergranular. However, the branched crack morphologies, and the fact that many cracks are separated by distances smaller than the average grain size of 3–4 µm, indicate that intragranular cracking occurs in both materials. Some cracks in the nanocomposite interact with SiC particles: (i) deflecting on a scale comparable to the size of nano-particle (see the cracks arrowed as “CD” in Fig. 5(b)); and (ii) exhibiting local crack-tip arrest by a relatively large particle (see particle arrowed as “CA” in Fig. 5(b)). Crack deflection on the micron scale occurs in both ceramics; this might be a result of interactions between cracks and grain boundaries. 3.4. Plastic deformation zone under the scratches Fig. 8 shows a TEM cross section containing two scratches, one in the alumina and one in the nanocomposite. By measuring the epoxy glue thickness in the cross section, it can be seen that in the vicinity of the scratches about 2 µm thick of material was removed from the original top surface at each side during TEM specimen preparation by ion milling. Even though the original scratch surface has not been preserved, the edge geometry (maximum scratch depth and width) is similar to the true profiles of the scratch surfaces as measured by FIB (Table 1), as expected if the material removal is nearly constant along the sample edge. A number of TEM cross sections show that beneath the scratches, a plastically deformed zone forms due to dislocation and twin activation in both the alumina and the nanocomposite. The measured maximum extent of the residual deformation from the centre of the groove surfaces was observed to be ~4 µm for the alumina and ~7 µm for the nanocomposite under the same scratching conditions. Residual plastic deformation observed under a scratch in alumina is shown in Fig. 9. Slip is acti-
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Fig. 6. Examples of multiple parallel 2D FIB cross-sections (spacing schematic) showing subsurface fracture damage under the scratch grooves in (a) Al2O3, (b) Al2O3/5vol.% nanocomposite. The viewing angle is 45° from the milled plane. The arrow indicates limited pile-up of material above the original surface height at the scratch groove edges.
vated only within the surface layer of grains. The residual dislocation density is extremely high, making it difficult to image individual dislocations and producing significant strain contrast. Some microcracks are observed in the TEM sections around the grain boundaries. In the nanocomposite, there is a much more extensive residual deformation zone, extending below the surface layer of grains. Similar to the alumina, the first layer of grains adjacent to the scratch exhibits such a high dislocation density and residual strain that it is almost impossible to characterise individual defects (grains G1 and G2 in Fig. 10(a)). In addition to dislocation activation, twinning was observed. These twins are confirmed by diffraction to be basal twins, with the (0001) planes oriented close to the normal loading direction of scratching. Twinning penetration might be partially limited by SiC nanoparticles. In grain G2
of Fig. 10(a), the growth front of the arrowed twin locally stops at the site of a SiC particle. Although the twinning partials will only be pinned at some points along their length, and can extend around the particles in 3D, it is likely that the pinning sites are effective in retarding the growth of the twins. There are about six twins visible in the upper part of the G2 grain, only one of which extends into the lower part of the G2 grain. The morphology of the basal twins is very high aspect ratio platelets, with (0001) habit plane and thickness in the [0001] direction of only about 20 nm measured using HREM. This platelet morphology is comparable to that of basal twins formed in sapphire [31]. Fig. 10(b) shows residual plastic deformation in the second and third layers of grains under the centre of the groove at depth of more than ~4 µm. The defect density decreases with increasing distance from the groove surface, and varies from
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surface grain containing a section of the plastic– elastic boundary for the scratch in Fig. 10. Adjacent grains have dislocation content indistinguishable from the bulk. Cracks were observed beneath the scratches in both the alumina and nanocomposite TEM samples. In alumina (Fig. 9), a crack is present close to the boundary between the first and second layer of grains. In the nanocomposite (Fig. 10), a large crack is present at the boundary between the first and second layer of grains, and ends at the limit of the residual plastic deformation zone in Fig. 11. It is not clear if all these cracks are formed during the scratch operation, or whether some at least arise and coalesce as a mechanism for stress relief during thinning of the TEM samples to electron transparency.
4. Discussion
Fig. 7. 3D reconstructions of the crack morphologies around a 1 N load scratch track (solid surface at the bottom of each image), viewed “inverted” from inside the sample. The scratch track is shown in grey, with the cracks white on a black free space background. (a) Polycrystalline alumina: reconstructed dimension 19.1 µm across scratch × 7.2 µm along scratch. (b) Polycrystalline Al2O3/5vol.% nanocomposite: reconstructed dimension 16.8 µm across scratch × 4.3 µm along scratch. Residual cracks around the scratch site penetrate more than 7 µm into sample from groove bottom in (a) and more than 10 µm in (b). Labels (A)–(F) denote microstructural regions discussed in the text.
grain to grain. The grains either side of the scratch groove also contain significant residual plastic deformation. For the scratch in Fig. 10, the maximum lateral extent of the residual deformation was ~6 µm from the groove centre in a plane 2 µm below the original surface. Fig. 11 shows a
In contrast to studies of ground, abraded, polished or eroded surfaces [12] where there are many overlapping adjacent contact events, these observations of single point scratches enable the plastic deformation and subsurface fracture patterns arising from a relatively simple contact geometry to be quantified. In contact with a sharp indenter, the elastic stress field is in principle singular about the indenter point, which inevitably leads to plastic deformation in most ceramics [20]. The activation of plastic deformation is controlled by the large components of shear and hydrostatic compression in the contact field. The TEM and FIB cross-sectional examinations of single point scratches in alumina and the nanocomposite show that the scratch operation results in dislocation activation, twin nucleation and microcracking. At the load used here (1 N), in the alumina, plastic deformation is only observed within the first layer here of grains contacted by the moving indenter (i.e. to a maximum depth of ~3 µm). The dominant plastic deformation modes in alumina found by other researchers are basal twinning and 1/3⬍112¯ 0⬎ dislocations gliding on basal and non-basal slip planes [31–36]. Some ⬍10 1¯ 0⬎ and pyramidal Burgers vectors may also be activated [31,34,35,37]. The defects appear to orig-
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Fig. 8. Cross-sectional TEM of scratches in alumina and nanocomposite. Dashed lines indicate the position of the un-scratched surface plane in alumina and nanocomposite, which was assessed by measuring the thickness of the epoxy filling gap near this regime. Table 1 Measured scratch depth and width from FIB cross-section image Material
Scratch depth (µm)
Scratch width (µm)
Alumina Nanocomposite
1.6 1.6
10.4 11.5
inate from the contact zone [31,33,35,36]. Inkson [31] found that under prismatic and basal surfaces of single crystal alumina abraded with 30 µm diamond grit, dislocation penetration was to about 5 µm below the surface. Penetration of basal twinning was found to have a crystallographic dependency [31,34–36], with penetration into prismatic surfaces occurring to ⬎10 µm [31]. In the polycry-
stalline alumina examined here, the average grain size is about 3–4 µm. The glide of dislocations and basal twins appears to have been stopped by the grain boundaries. In the nanocomposite, dislocation and twin activation was observed to a rather greater extent than in alumina, penetrating to below the surface layer of grains (Fig. 10(b)) to a depth of about 10 µm. This may be due to (a) defect propagation across the grain boundaries, and (b) activation of subsurface dislocation sources. Considering possible subsurface dislocation sources, dislocation nuclei might be formed by local thermal residual stress around the SiC particles due to the difference in thermal expansion coefficients of alumina (~8.3 × 10⫺6 K⫺1) and silicon carbide (~5.1 × 10⫺6 K⫺1). During cooling
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Fig. 9. Plastic deformation beneath the scratch groove in monolithic Al2O3.
from the processing temperature (1650–1680°C), residual stresses will build up to a level of near 2 GPa near the Al2O3/SiC interface (SiC particle in compression, Al2O3 matrix in tension) in the room temperature [38,39]; at intermediate temperatures (about 400–1100°C), dislocations with basal and non-basal Burgers vectors may be activated [40– 42]. In addition, dislocations may be activated at grain boundaries which have not undergone stress relief by cracking. Under the contact stress field, these dislocations may glide further and/or act as sources for further dislocation generation. The 3D FIB mapping of crack morphologies around the residual scratch resolves microcracks under the scratch track in both alumina and the nanocomposite (Fig. 7). We assume that the act of milling a hole across the scratch track does not change the local stress field sufficiently to cause marked alteration in the original crack profiles, for example by crack propagation towards the new free surface. The stress relief in the FIB sectioning, with one free surface and surrounding constraining material, should be much less than that in a TEM foil. The 3D crack maps exhibit very similar crack morphologies for both materials, which may be broadly described as consisting of three zones, as shown in Fig. 12:
Fig. 10. Plastic deformation in the grains beneath the centre of the scratch groove in Al2O3/5vol%SiC nanocomposite. Two adjacent grains (G1 and G2) have plastically deformed by the activation of dislocations in grain G1 and a mixture of dislocation and basal twin in grain G2 where selected area diffraction shows basal twins; (b) slip originating from a grain boundary in a grain under grain G2; grain (G3) under grain G1 is heavily strained with a high density of dislocations.
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Fig. 11. Plastic deformation features to one side of the scratch. The plastic/elastic boundary can be seen near the centre of the micrograph.
Fig. 12. Schematic diagram of types of subsurface cracking observed beneath scratches. The cracked volume is composed of three zones: zone I (crack deficient), zone II (microcracking) and zone III (deep lateral and median cracking).
(I)
A crack deficient zone directly under the centre of the scratch groove. (II) A microcracking zone below and around zone (I), which extends up to the surface at the two groove edges. (III) A zone below zones (I) and (II), with deep lateral and median cracks which do not intersect the surface. The maximum depth of cracking (⬎7 µm) is comparable to the depth of the dislocation zone (~7 µm) in the nanocomposite, but almost twice the depth of the dislocation zone (~4 µm) in the alumina. Considering first the crack deficient zone (I) directly under the centre of the scratch groove, this zone has a very high residual dislocation density in both the nanocomposite and the alumina. This
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zone will have a high compressive residual stress after the indenter moves off [19]. This appears to inhibit both the nucleation of cracks and the propagation of subsurface cracks up to the groove surface. With no crack intersection and associated material chipping, the groove bottom appears optically smooth after scratching in both materials. This contrasts with grinding and polishing, where nanocomposites may exhibit much better surface qualities than pure aluminas [5,6,8,9]. Surrounding the compressive zone, microcracking occurs (zone (II)). The microcracking extends up to the surface at the groove edges, where material removal by crack linking and chipping occurs. Microcracking is expected to be observed where the stress field starts to be dominated by tensile components [43]. Close to the compressive zone, the microcracks have a heavily branched morphology, and in the sampled regions there appeared to be a marginally higher density of microcracks in the nanocomposite than in the alumina. In both the alumina and nanocomposite, the microcracking zone (II) then evolves into a zone (III) of deep lateral and median cracks. The nucleation and propagation of cracks within the tensile stress field in both the alumina and nanocomposite appears to be comparable, since no major difference in the residual crack morphologies was observed in this study. This is consistent with measurements of the fracture toughness Klc for the two materials, which are found to be similar [7,10,22,23], During the scratching process, the activation of dislocations will change the local stress concentration distribution around the tip. In the alumina the dislocations are activated within the top layer of grains, which corresponds to the crack-free compressive zone (I) and the microcracking zone (II). In the nanocomposite, dislocations are also activated in zone III below the first layer of grains where deep lateral and median cracks are observed. Dislocation activation may affect crack nucleation and propagation by: (i)
Causing directional stress concentrations due to difficulties in initiation of enough independent slip systems to accommodate shape change (for example, dislocations with non-
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(ii)
(iii)
(iv)
(v)
(vi)
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basal Burgers vector components), leading to the nucleation of cracks on grain boundaries or inside grains [31,33,35–37]. Generating basal twins, the (0001) interfaces of which are preferential crack paths [31]. Microcracking initiated by the intersection of rhombohedral twins as found by Heuer [44] was not observed in these specimens. Causing stress concentrations at dislocation pile-ups within grains and at grain boundaries, which may encourage microcrack nucleation [31,34,37]. In the nanocomposite, the distribution of SiC nanoparticles may additionally affect crack nucleation and propagation by: Providing subsurface dislocation sources, generated during cooling to accommodate thermal misfit as discussed above. Directly causing spontaneous crack initiation at the SiC–Al2O3 interfaces to relieve thermal mismatch. Some large SiC particles or agglomorates may in principle act as crack initiation sites; however, no such cracks were observed around SiC particles in the bulk material away from the scratch zone. Causing crack deflection or pinning at nanoparticles. FIB imaging shows the deflection of some cracks at SiC particles⬎30 nm in size (smaller particles are below the resolution limit). The observation of cracks terminating/beginning at SiC particles in the 2D cross sections was rare and only observed at the larger particles. With only 5 vol.% SiC, these pinning events may simply be coincidental juxtapositions. This is different from the frequent and property-dominating crack/second phase interactions in large particle, high vol.% ceramic composites [45].
The difference in depth of dislocation activation between the alumina and nanocomposite, under the present loading conditions, appears not to alter significantly crack initiation and propagation by mechanisms (i)–(vi) under the loading conditions used. Both materials have similar subsurface crack morphologies and a compressive stress zone inhibiting crack propagation to the surface at the
centre of the groove. This results in the scratch track morphologies appearing very similar. The measurements of residual stress across the scratches in alumina and the nanocomposite by Cr3+/Al2O3 fluorescence spectroscopy are consistent with the microstructural observations by TEM and FIB. The much larger broadening of the fluorescence lines observed in the nanocomposite arises from a greater average density of dislocations inside the probed volume, due to the greater depth of the dislocated layer [25]. Stress inhomogeneities due to microcracking inside the groove will also contribute to the broadening [25], but the observed similar microcracking patterns under the grooves in alumina and the nanocomposite will give rise to similar broadening contributions. The fluorescence shift from the value for a stress-free surface (i.e. a fracture surface of the same ceramic) represents an average stress level inside the probed volume. The shift is much greater in the nanocomposite than in alumina (Fig. 4); this can also be attributed to the more extensive residual dislocation zone in the nanocomposite than in the alumina, giving rise to greater and more extensive elastic stresses outside this sub-scratch plastic zone, and is consistent with measurements of surface residual stresses on machine ground surfaces [7]. 5. Summary 1. Scratch tracks made with a 1 N normal load in the alumina and an alumina/5% silicon carbide nanocomposite appear similar by optical microscopy. 2. Cross-sectional TEM shows that dislocation activation in the alumina was in the first layer of grains only (~3 µm depth), and in the nanocomposite, penetrated to below the first layer (to ~10 µm). 3. FIB sectioning shows that both alumina and the nanocomposite exhibit similar 3D crack patterns around the scratches: (I) a crack-free compressive zone under the smooth groove base (~3 µm), (II) a microcrack zone below and around zone (I) (~3 µm), (III) a deep lateral and median crack zone below (I) and (II). Microcracks extend to the surface at the groove edges and produce chipping at the groove edges.
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Residual stress levels (measured by Cr+ fluorescence) across the scratches are higher and more extensive in the nanocomposite. Line broading near the scratch is also greater in the nanocomposite. This is consistent with the TEM observations of a larger plastic zone in the nanocomposite.
Acknowledgements We gratefully acknowledge the support of the EPSRC under grant GR/L95908 (H.Z.W.), and of the Royal Society (B.J.I.). We thank Prof R. Young at the Manchester Materials Centre for providing facilities for fluorescence measurements, and Prof. B. Derby (also at MMC) for helpful discussions.
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