Nanoindentation and fracture toughness of nanostructured zirconia/multi-walled carbon nanotube composites

Elsevier Editorial System(tm) for Ceramics International Manuscript Draft Manuscript Number: CERI-D-14-04126R1 Title: Nanoindentation and fracture toughness of nanostructured zirconia/multi-walled carbon nanotube composites Article Type: Full Length Article Keywords: Ceramic-matrix composites (CMCs); Carbon nanotubes; Mechanical properties; Fracture toughness; short pulse laser machining. Corresponding Author: Ms. Latifa Melk, Ph.D. Corresponding Author's Institution: Universitat Politècnica de Catalunya (UPC) First Author: Latifa Melk, Ph.D. Order of Authors: Latifa Melk, Ph.D.; Joan Josep Roa Rovira; Fernando García-Marro; Marta-Lena Antti; Ben Milsom; Michael Reece; Marc Anglada Abstract: Multi-walled carbon nanotubes (MWCNTs)/3 mol% yttria-doped tetragonal zirconia (3YTZP) composites were produced using spark plasma sintering (SPS) with MWCNT content ranging within 0-2 wt%. In the present paper, it was shown that the addition of MWCNTs results in a refinement of the composites microstructure. Moreover, nanoindentation tests were performed in order to monitor the change in elastic modulus and hardness with MWCNT content and it was found that both properties decrease with the addition of MWCNT content. A novel method was used to measure the true fracture toughness of the composites by producing a shallow surface sharp notch machined by ultra-short pulsed laser ablation on the surface of beam specimens. The true fracture toughness obtained on this laser machined single edge V-notch beam (SEVNB) specimens tested in four point bending was compared to the indentation fracture toughness measured using a Vickers indenter. It was found that the indentation fracture toughness increases with increasing MWCNT content, while the true fracture toughness determined with SEVNB was practically independent of the composition. Finally, it was concluded that the increase in the resistance to indentation cracking of the composites with respect to 3Y-TZP matrix cannot be associated to higher true fracture toughness. The results were discussed in terms of transformation toughening, damage induced in front of the notch tip, microstructure of the composites, and fracture toughness of 3Y-TZP. We thank the referees for their comments which have been helpful to increase the quality of the manuscript. Please find below the answers to the referee comments. We have also made slight changes in the manuscript in red colour. Response on Referee comments: 1) Beginning of section 3.1 describes density of the prepared materials. How big is the porosity? Also, it would be good if the pores were characterized in terms of size and shape, and possibly locations (I guess they would be connected to the MWCNTs). This might be of a consequence for the crack propagation. The drop in density of the composites seen in Table 1 is related to the presence of porosity connected to the CNT agglomerates. Fig. 1 shows the fracture surface of the composite with 2 wt% CNT content where some pores related to the clusters of MWCNTs could be observed. In order to measure with better accuracy the grain size of the composites and to see clearly the grain boundaries, the composites were etched by a heat treatment at high temperature in air and the locations of the burned out CNTs were revealed. See fig. 2 of the manuscript. (See changes in red colour in the manuscript) Fig. 1 Fracture surface showing the presence of porosity in the composite with 2 wt% CNT content. Detailed Response to Reviewers 2) Section 3.2 and Fig.3 describe the nanoindentation results. For most of the materials the values stabilize for depths <100 nm, it is only for the one with 0.5 % CNTs that they increase very slowly. What is the reason? Is it some real problem of the material or the tip, or it can be just because of some really bad measurement which should have been eliminated from the statistics? The scatters are not too large, so it seems it could be an artifact coming from tip geometry. Could the authors comment on that? As the reviewer comments, it is true that the scatters for the 0.5% wt. CNT is not too large, however, the hardness trend does not stabilize until 1000 nm of displacement into surface as it can clearly be observed in Fig. 3a. This phenomenon cannot be attributed to an artifact coming from the tip geometry because the indenter tip was calibrated using fused silica and all the samples were tested in the same day under the same experimental conditions. Furthermore, all the different tests performed for this condition presented the same trend. Therefore, the behavior of 0.5 wt% CNT could be attributed to a real problem of the material. However, this does not change the trend of the effect of CNT on 3Y-TZP presented in the manuscript. No additional changes have been introduced in the reviewed manuscript 3) Page 14, 1st paragraph - the increase in contact damage resistance measured by indentation is attributed to the weak interfacial bonding between the matrix and CNTs. It is difficult to judge, since no photographs of crack lines and no information on porosity are presented. Is it possible that the cracks are arrested by voids or microcracks? There are slight differences in morphology of the crack path (see figure below), but we understand that they are small and do not contribute significantly to the indentation fracture toughness. We rather associate the enhanced contact-damage resistance related to a reduction in the driving force for crack growth related to change in plastic deformation behaviour below the indentation. The intense confined-shear under Vickers indentation in the presence of shear-deformable MWCNTs in the small agglomerates increases irreversible deformation by this mechanism as well as by compaction of porosity below the indenter instead of by plastic deformation. The series of indentation fracture toughness of the literature are based on the assumption of plastic deformation below the indenter but here they are not fully applicable: as the CNT concentration increases there are additional mechanisms to change shape below the indenter (shear of bundles of nanotubes and porosity compaction) which induce a smaller driving force for crack growth. (See changes in red colour in the manuscript). Fig. Indentation cracks in the composites with a) 2 wt% CNT and with b) 0% wt% CNT (Red and blue lines were used to measure crack length and diagonals of residual imprint respectively). This picture is not incorporated in the manuscript 4) Page 14, section "Damage" claims that no slip lines were observed (Fig. 7). The size of the indents is simply too big, resolution of the images in Fig.7 and quality of the surfaces is completely inadequate to observe the slip lines. While I agree that it is likely that the dominant process is the (nano) grain sliding and not their plastic deformation or transformation, the evidence as presented is insufficient. We completely agree with this comment. The resolution of this image is completely inadequate. This was a mistake by our part for which we apologize. We have modified the text like that: No fracture events like radial cracks at the corners of the imprints were detected. (See changes in red colour in the manuscript). 5) It would be interesting to see the morphology of the laser machined notch in cross section. Can the authors provide such a picture? We have incorporated in the manuscript a new figure (Fig. 5) in which the notch for the composite with 0 wt% CNT content is presented. In order to get a close overview of the notch, the top surface of the notch (surface on which the notch was induced) was grinded step by step, then polished till approximately the mid depth of the notch is reached. The damage zone below the notch is hardly seen before a) b) fracture. However, the damage could easily be detected after fracture. See fig. 6 of the manuscript. (See changes in red color in manuscript). Fig. 5 View of the morphology of the notch induced by UPLA on the specimen with 0 wt% CNT content. 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 $Corresponding author at: Department of Materials Science and Engineering/ETSEIB Diagonal 647. 08028-Barcelona Tel. +34 934016701 Fax. +34 934016706 E-mail: [email protected] 1 Nanoindentation and fracture toughness of nanostructured zirconia/multi-walled carbon nanotube composites Latifa Melk ‡,§,$ , Joan J. Roa Rovira‡, Fernando García-Marro‡, Marta-Lena Antti§, Ben Milsom*,¶, Michael J. Reece*,¶, and Marc Anglada‡† ‡Department of Materials Science and Engineering, Universitat Politècnica de Catalunya, Barcelona, 08028, Spain §Department of Engineering Sciences and Mathematics, Luleå University of Technology, Luleå, Sweden *Department of Materials, Queen Mary College, University of London, London E1 4NS, UK ¶Nanoforce Technology Limited and School of Engineering and Materials Science, Queen Mary, University of London, London E1 4NS, UK. *Manuscript Click here to view linked References 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 2 Abstract Multi-walled carbon nanotubes (MWCNTs)/3 mol% yttria-doped tetragonal zirconia (3Y-TZP) composites were produced using spark plasma sintering (SPS) with MWCNT content ranging within 0-2 wt%. In the present paper, it was shown that the addition of MWCNTs results in a refinement of the composites microstructure. Moreover, nanoindentation tests were performed in order to monitor the change in elastic modulus and hardness with MWCNT content and it was found that both properties decrease with the addition of MWCNT content. A novel method was used to measure the true fracture toughness of the composites by producing a shallow surface sharp notch machined by ultra-short pulsed laser ablation on the surface of beam specimens. The true fracture toughness obtained on this laser machined single edge V-notch beam (SEVNB) specimens tested in four point bending was compared to the indentation fracture toughness measured using a Vickers indenter. It was found that the indentation fracture toughness increases with increasing MWCNT content, while the true fracture toughness determined with SEVNB was practically independent of the composition. Finally, it was concluded that the increase in the resistance to indentation cracking of the composites with respect to 3Y-TZP matrix cannot be associated to higher true fracture toughness. The results were discussed in terms of transformation toughening, damage induced in front of the notch tip, microstructure of the composites, and fracture toughness of 3YTZP. Keywords: Ceramic-matrix composites (CMCs); Carbon nanotubes; Mechanical properties; Fracture toughness; short pulse laser machining. 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 3 1. Introduction In the last decades there has been growing interest in developing ceramic materials with high fracture toughness (KIc) and strength for structural applications. In the specific case of 3 mol% yttria-doped tetragonal zirconia (3Y-TZP), KIc can be increased by promoting phase transformation from tetragonal (t) to monoclinic (m) phase in front of a propagating crack tip (so called transformation toughening) [1]. However, as stronger is the tendency for stress induced transformation, as higher is the risk for premature spontaneous t-m transformation on the external surface in contact with moisture in the environment. This effect is referred to as hydrothermal degradation or low temperature degradation (LTD) and it is accompanied by surface microcracking. This phenomenon is the main drawback for the wider use of 3Y-TZP. It is more severe at higher temperatures, for larger grain sizes, and under the presence of tensile residual stress or low concentrations of yttria [1]. The resistance to LTD can be strongly increased by reducing the grain size into the nanoscale. However, this will reduce the transformation toughening capability, so that the fracture toughness will diminish [2]. One strategy to counteract LTD is the incorporation of a second phase as a toughening mechanism in the form of particles, grains, whiskers or fibres into the nano-grain zirconia matrix. Recent investigations have been focused on the addition of carbon nanotubes (CNTs) as reinforcements into a ceramic matrix because of their high aspect ratio and outstanding mechanical properties [3]. Moreover, it was reported that the addition of CNT to TZP-based composites improves the fracture toughness without affecting the resistance to LTD [1]. 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 4 On the other hand, there is still a debate about the efficiency of CNTs in the mechanical properties of the ceramic matrix nanocomposites. Zhan et al. [4] reported an important improvement in indentation KIC (KIC 9.7 MPa.m1/2) with addition of 10 vol.% SWNTs to Al2O3 matrix, but Wang et al. [5] found a much smaller improvement (only 3 %) when KIC was measured by the single edge V-notch beam (SEVNB) method on the same nanocomposites. A key observation made on these alumina composites is that they can be more contact-damage resistant than alumina, as it was shown by the lack of crack formation during indentation tests, and, in the same time, they were brittle as dense Al2O3 with low KIc measured by the SEVNB method [5]. Therefore, it was concluded that for the alumina composites studied, a high tolerance to indentation loading does not necessarily represent an increase in KIc. Some studies about the incorporation of multi-walled carbon nanotubes (MWCNTs) to 3Y-TZP matrix reported an increase in indentation KIc using different techniques for the processing of 3Y-TZP/CNT composites [6–9]. Values up to 13 MPa.m1/2 have been reported for the addition of 1 wt. % SWNTs [10]. Other authors have found that the increase in indentation KIc only occurs for the addition of relatively small CNTs content: 0.5 wt. % ([11,12]), 1.0 wt. % [11]. However, an increase of KIc was also reported for the addition of 4 wt% of functionalized CNTs [12]. While, a decrease in indentation KIc was usually observed for concentrations higher than 2 wt% CNT [11–15] and explained by the poor dispersion of CNTs and the weak bonding between CNTs and zirconia matrix. Moreover, a decrease in KIc was reported in 1.07 wt% CNF/CNT [16,17] as well as by adding a small fraction of 0.69 wt. % carbon nanofibers (CNF) [18]. Regarding the influence of CNT on the Vickers hardness to 3Y-TZP, most of the studies have concluded that with the addition of either small CNT content ( 0.5. %) 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 11 0 up to 2 wt% CNT respectively (see Table 1). The decrease of grain size with adding MWCNTs content has been attributed to the presence of MWCNTs in grain boundaries which slows [9] or hinders the grain growth of zirconia grains [7]. 3.2 Hardness and elastic modulus Fig. 43 shows the Berkovich hardness and the elastic modulus of all composites in terms of penetration depth. It is clearly seen that the values are stabilised for penetration depth larger than 400 nm, while for lower penetration depths, these properties are affected by surface defects and roughness. Moreover, both HBerk and EBerk of all composites are smaller than for monolithic 3Y-TZP (see Table 1). On the other hand, the Meyer hardness of 3Y-TZP/CNT diminishes in a similar way to HBerk, but with less scatter (see Table 1). Therefore, the trend is the same whether a small or relatively large surface area is tested. Since the ratio (EBerk/H)1/2 practically does not change with MWCNT content and the length of the indentation cracks is shorter with increasing MWCNT amount, the indentation KIc increases (see Table 1). The value of KIc depends on whether the equation of Anstis et al. [26] or Niihara et al. [25] is used. Values in the range of 3.5– 5.0 MPam1/2 were found using Anstis et al. [26] (see Ttable 1), while the values were about 20 % higher if KIc was calculated using the equation of Niihara et al. [25] Therefore, it was concluded that the indentation KIc increases for the compositions studied independently of the equation used. With respect to the 3Y-TZP matrix, the observed contribution of 0.5 wt% MWCNT to the indentation KIc, is in a good agreement with similar composition investigated by Mazaheri et al. [7]. In the same direction, Ukai et al. [13] showed that the addition of 0.5 wt% CNTs improves indentation KIc in about 0.5 MPam1/2, similar to the present 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 12 results. However, our absolute values of indentation KIc are noticeably smaller even for the starting 3Y-TZP matrix. 3.3 True fracture toughness Fig. 54 shows the morphology of the notch induced by UPLA in the composite with 0 wt% CNT content. In order to get a close overview of the notch, the top surface of the notch (surface on which the notch was induced) was grinded step by step, then polished till approximately the mid depth is reached. The damage zone below the notch is hardly seen before fracture. However, the damage could easily be detected after fracture. See Fig. 6. Fig. 6 shows the fracture surface of the notched in four point bending specimens for 0, 0.5 and 2 wt% MWCNT, where three different regions are well differentiated. The first one (region A), at the top, corresponds to the notch, region B has a length of about  24 μm with a different fracture surface appearance, and region C at the bottom has the usual common surface fracture appearance of zirconia. The transition between regions B and C is clearly defined by a relatively sharp straight border normal to the crack advance (see Fig. 75). The extension of region B increases slightly with the MWCNT content. This corresponds to the laser ablation damage region in front of the notch tip [29]. The crack length used in the calculation of KIc is easily identified as the notch length (region A) plus the length of the microcracked damaged zone (region B) clearly observed on the fracture surface. The effective length of the crack was taken as the length of the notch plus the length of region B where the damage was observed. Therefore, the different appearance of region B is associated, in the current study, to the coalescence of the 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 13 damage induced by laser ablation in front of the notch during fracture. The damage is restricted to the length of region B and to a maximum average depth of 2 μm below the fracture surface. This can be appreciated in Fig. 86 where sections of region B (Fig. 86a) and of the border between regions B and C (Fig. 86b) are shown. On the left (region B) microcracking and porosity can be observed below the surface while just to the right in region C no trace of microcracking is observed. Similar observations using a notch tip induced by UPLA in conventionally sintered 3Y-TZP have been observed [29]. The true KIc values of the notched Sspecimens using UPLA are given in Table 1, where it can be appreciated that KIC values hardly change with the addition of MWCNT content. Moreover, the true KIc of monolithic SPSed 3Y-TZP is found to be smaller (KIc=2.7 MPam1/2) than the conventionally sintered monolithic 3Y-TZP of larger grain size (330 nm) where a KIc slightly higher than 4 MPam1/2 was reported previously by the present authors [29]. It is interesting to notice that the higher values of KIc reported in the literature using the standard SEVNB method for the same composites include also a high KIc value for the starting monolithic 3Y-TZP matrix of about 6 MPa MPam1/2 [20]. It should be noticed that this is a very high value for 3Y-TZP with a grain size of 145 nm. Since the starting powder, grain size and density of SPS 3Y-TZP reported in [20] are very similar to the present study, the difference regarding the KIc reported here could be associated to the sharpness of the starting notches. The present low KIc for 177 nm grain size 3Y-TZP (2.7 MPam1/2) is in line with the results of Eichler et al.[33] where they reported values between 2.8 and 3.1 MPam1/2 for 3Y-TZP with grain sizes in the range between 110 and 210 nm. Therefore, it may be concluded that at least part of the low KIc values measured 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 14 for the composites with respect to [20] is related to the matrix and possibly to the sharpness of the notch and length of microcracks induced in front of the notch. As is shown above, here the damage in the form of microcracks in front of the notch is clearly revealed on the fracture surface. It is much longer than the notch radius, and it ends abruptly at some distance of the notch tip. This is the reason for taking the length of the initial crack as the length of the notch plus the length of the microcracked zone as it has been shown in [29]. If only the notch length was considered as the crack length, then even much lower values of KIc would be found. On the other hand, it is unlikely that the low value of KIc is related to the short depth of the shallow starting notch since the R-curve in nano grain size 3Y-TZP is relatively weak and very steep [33]. As seen in Table 1, the true KIc does not practically increase with MWCNT content. However, the contact damage resistance measured by indentation KIc increases. There were slight differences in the morphology of the crack paths with the addition of MWCNTs. However, the crack lengths are small and do not contribute significantly to the indentation fracture toughness. We rather associate the enhanced contact-damage resistance related to a reduction in the driving force for indentation crack growth related to change in plastic deformation behavior below the indentation. The intense confinedshear under Vickers indentation in the presence of shear-deformable MWCNTs in the small agglomerates increases irreversible deformation by this mechanism as well as by compaction of porosity below the indenter instead of by plastic deformation. The reason for this behaviour is probably related to the weak interfacial bonding between MWCNTs and zirconia matrix. 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 15 Damage No fracture events like radial cracks at the corners of the imprints were detected No damage events slip lines or transformations were observed just under the Berkovich indentations (see Fig.97). This indicates that the presence of fine grain microstructure (see Fig.32) improves the damage tolerance under contact loading and it blocks the deformation mechanisms induced during the indentation process, which is similar to the behaviour reported in particle-dispersion-toughening of ceramic based nanocomposites [34]. The amount of t-m phase transformation below the indenter has not been measured in the present study. It was reported in a previous work [6], that t-m transformation is reduced in 3YTZP/CNT composites because of the small grain size. With the addition of 2 vol% MWCNT with grain size similar to the present study, t-m transformation at the fracture plane of Vickers indentation cracks was less than 5% at distances less than  2.0 µm below the fracture surface. Therefore less transformation would be expected below the indenter since the main stress state is compression. 4. Conclusions True KIc of spark plasma sintered 3Y-TZP/CNT composites was successfully calculated using a novel method based on inducing a very sharp notch by UPLA. The true KIc for the SPSed 3Y-TZP matrix studied is smaller as compared to presureless conventional sintered 3Y-TZP. Moreover, the true KIc hardly increases with the addition of MWCNT while the indentation KIc of the starting 3Y-TZP matrix is higher and it increases with MWCNTs content. 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 16 It is concluded that the resistance to indentation cracking of the composites by adding MWCNTs to 3Y-TZP matrix does not indicate higher true fracture toughness. It is suggested that this is related to the drop in hardness and elastic modulus with the addition of MWCNTs content as it was detected using Berkovich nanoindentation. Acknowledgements The authors gratefully acknowledge the financial support given by the “Ministerio de Economía Ciencia y Competitividad” of Spain through research grant MAT201123913. L. M. acknowledges the fellowship received from the Joint Doctoral Programme in Materials Science and Engineering (DocMASE). 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Y. 48 (2010) 2994–2996. [33] J. Eichler, J. Rödel, U. Eisele, M. Hoffman, Effect of Grain Size on Mechanical Properties of Submicrometer 3Y-TZP: Fracture Strength and Hydrothermal Degradation, J. Am. Ceram. Soc. 90 (2007) 2830–2836. [34] H. Awaji, S.-M. Choi, E. Yagi, Mechanisms of toughening and strengthening in ceramic-based nanocomposites, Mech. Mater. 34 (2002) 411–422. 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 61 62 63 64 65 20 Table captions: Table 1. Properties of the 3Y-TZP/CNT nanocomposites. Figure captions: Fig. 1. Fracture surface showing the presence of porosity in the composite with 2 wt% CNT content. Fig. 21. Fracture surfaces of the sintered nanocomposites; (a) 0.5 wt% CNT, (b) 1 wt% CNT, (c) 2 wt% CNT composites. Fig.32. Microstructure of (a) 0 wt% CNT, (b) 0.5 wt% CNT, (c) 1 wt% CNT, (d) 2 wt% CNT composites after etching in air at high temperature. Fig. 43. Berkovich hardness and elastic modulus of all composites in terms of penetration depth. Fig. 5. View of the notch induced by UPLA on specimen with 0 wt% CNT content. Fig. 64. Fracture surface of notched specimens: (a) 0 wt% CNT, (b) 0.5 wt% CNT and (c) 2 wt% CNT composites. A is the surface of the notch, B the microcracked region in front of the notch and C is the final fracture. Fig. 75. High Mmagnification of the transitions between regions A, B and C. Fig. 86. (a) FIB cross section of the fracture surface below region B: (b) fracture surface with a trench at the border of regions B and C. The border coincides with the disappearance of damage below the surface. Figure 5 Click here to download high resolution image Figure 6 Click here to download high resolution image Figure 7 Click here to download high resolution image Figure 8 Click here to download high resolution image Figure 9 Click here to download high resolution image