Transmission electron microscopy characterization and sculpting of sub-1 nm Si-O-C freestanding nanowires grown by electron beam induced deposition

Thin Solid Films 518 (2010) 4890–4897

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Transmission electron microscopy characterization of TiN/SiNx multilayered coatingsplastically deformed by nanoindentation

M. Parlinska-Wojtan a,b,⁎, S. Meier c, J. Patscheider a

a Laboratory for Nanoscale Materials Science, Empa, Überlandstr. 129, CH-8600 Dübendorf, Switzerlandb Centre for Electron Microscopy, Empa, Überlandstr. 129, CH-8600 Dübendorf, Switzerlandc Electronics/Metrology/Reliability, Empa, Überlandstr. 129, CH-8600 Dübendorf, Switzerland

⁎ Corresponding author. Centre for ElectronMicroscop8600 Dübendorf, Switzerland. Tel.: +41 (44) 823 4778.

E-mail address: [email protected] (M. P

0040-6090/$ – see front matter © 2010 Elsevier B.V. Adoi:10.1016/j.tsf.2010.02.064

a b s t r a c t

Article history:Received 12 August 2009Received in revised form 18 February 2010Accepted 26 February 2010Available online 6 March 2010

Keywords:Multilayer thin filmsTitanium nitrideShear slidingSevere plastic deformation (SPD)Transmission electron microscopy (TEM)

Plastic deformation of TiN5 nm/SiN0.5 nm multilayers by nanoindentation was investigated by transmissionelectron microscopy in order to identify deformation mechanisms involved in film failure resulting fromsevere plastic deformation. The TiN layers exhibited a crystalline fcc structure with a [002] preferentialorientation; further crystal growth was interrupted by the amorphous SiNx layers. After severe plasticdeformation collective vertical displacement of slabs of several TiN/SiNx-bilayers, which resulted from shearsliding at TiN/TiN grain boundaries, was observed. They are, together with horizontal fractures along the SiNx

layers, vertical cracks under the indenter tip following the TiN grain boundaries and delamination from thesubstrate, the predominant failure mechanisms of these coatings. The deformation behaviour of these filmsprovides an experimental support for the absence of dislocation activity in grains of 5 nm size.

y, Empa, Überlandstr. 129, CH-

arlinska-Wojtan).

ll rights reserved.

© 2010 Elsevier B.V. All rights reserved.

1. Introduction

Titanium nitride was introduced in the early 1980s as protectivecoating material for cutting and forming tools. To enhance itshardness, coatings with nanosized grains are designed either in theform of nanocomposites, composed of crystalline TiN nanograinssurrounded by an amorphous SiN matrix [1–5], which can bedeposited by different techniques[6–10], or alternatively in the formof multilayers [11–15]. The nanoscaled multilayers are constituted bytwo hard transition metal nitrides with individual layer thickness ofthe order of several nanometers. Such coatings exhibit excellentproperties like high hardness and highwear resistance [16–18], whichdepend critically on the thickness of the individual layers [18,19] aswell as on the nature of the interface [20].

The most popular method to measure the hardness of coatings isnanoindentation [21], however it is a destructive method, whichinduces plastic deformation into the coatings, and allows thus toobserve the occurring deformation mechanisms and the damagesgenerated in the material. The knowledge of these mechanisms willhelp in a targeted design of coatings with better hardness properties.The deformation in nanocrystalline materials such as nanocompositesoccurs by shear sliding at grain boundaries, grain boundary rotationand collective displacement of non-deformed nanocrystals [22–24]and as observed by Cairney et al. [25] via cracking at the nanocrystal-

line boundaries under local tension in a quasi-plastic manner.Moreover it has to be pointed out that in nanocomposites the grainsize is too small to be deformed by dislocation activity [23].

The nanomultilayers built of alternating layers of nc-TiN andamorphous SiNx allow an independent variation of the single layerthickness and an unambiguous visualization of the amorphous phase,both as prepared and after being plastically deformed. The interca-lating SiNx layers act as markers to help retrieve the original positionof the multilayer segment prior to deformation. Such nanosizedmultilayers with alternating layers of TiN and SiNx were described byChen et al. [26] with a hardness maximum found for a SiNx thicknessof 0.5 nm. Another study on the hardening mechanism present inmultilayered Si3N4/TiN coatings was performed by Xu et al. [27]. Theyfound that the hardness of the multilayers is affected not only by themodulation period, but also by the layer thickness ratio anddeposition temperature. These authors suggested that the alternatingstress field caused by the mismatch of the thermal expansioncoefficients of Si3N4 and TiN is one of the main reasons for thehardness enhancement in this multilayered system. Recently Söder-berg et al. [12–15] investigated nanomultilayers of TiN and SiNx andfound that the mechanical properties, while varying the individuallayer thickness, yielded similar values as encountered in isotropicallynanostructured nanocomposites.

In the present study the deformation mechanisms induced byindentation in nanomultilayered TiN/SiNx coatings were analyzed byTEM on lamellae cut by focused ion beam (FIB) through indentsproduced at different loads. We provide an experimental visualizationof the damages caused to the coating during the different stages of

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deformation. It is shown which deformation processes happen atincreasingly higher loads causing plastic deformation.

2. Experimental details

TiN/SiNx multilayered films were deposited by reactive unbal-anced magnetron sputtering from elemental titanium and silicontargets. 100 repetitions of 5 nm TiN and 0.5 nm SiNx were preparedresulting in a total layer thickness of 550 nm. The operating pressurewas 0.5 Pa at a sputter gas composition of Ar/N2 of 4:1. The depositiontemperature was kept at 200 °C; a negative substrate bias voltage of70 Vwas used to provide for sufficient surfacemobility of the growingfilm. For improved adhesion on the silicon substrates an interlayer of50 nm TiN was applied. In order to enhance the interface smoothnessat these low temperatures helium (amounting to 10% of the total gasflow) was added to the gas feed. The film quality in terms of singlelayer thickness and interface sharpness was verified by X-rayreflectometry. These results confirmed that the intended layersequence and individual layer thickness was deposited. The TiNlayers in the films have a pronounced [200] texture.

The TiN/SiNx films were indented at 100 nm, 250 nm, 500 nm and1000 nm (initial penetration depth of the indenter tip) correspondingto 20%, 50%, 100% and 200% of the total film thickness by aNanoindenter XP system with a Berkovich indenter tip in order toproduce permanent deformation. The tip rounding was of the order of50 nm. The indentation depths used in this study exceed by far thoseused for hardness determinations as it is intended to deliberatelydestroy the film in order to study the deformation mechanisms.

The microstructure of the as-deposited and indented TiN/SiNx

multilayeredfilmswas studied by transmission electronmicroscopy ina conventional Philips CM30 equipped with a LaB6 cathode and thehigh resolution transmission electron microscopy (HRTEM) imageswere obtained in a Philips CM300 with a field emission cathode bothinstruments operating at 300 keV. A focused ion beam instrument(FIB-Dual Beam FEI STRATADB235)was used to prepare TEM lamellaethrough the indent in the multilayered coating as shown in Fig. 1(a).The obtained lamella with final dimensions of 20 μm×5 μm×80 nm,

Fig. 1. (a) Schematic view of the lamella cut by FIB through the indent; (b) SEM imageof the ready lamella cut through the indent: the delaminating coating and the crack inthe substrate are clearly visible.

shown in Fig. 1(b) has the electron transparent area in its central part,and the indented film is clearly seen. It was transferred under anoptical microscope onto a 3 mmCu grid coveredwith a carbon film forTEM observations.

3. Results

Fig. 2(a) shows a dark field image taken with the (002) reflectionof the TiN/SiNx multilayer sample composed of 100 repetitions of5 nm TiN and 0.5 nm SiNx undeformed layers. The film is composed ofcrystalline TiN grains; the multilayering is evident, and it can beclearly seen that the SiNx layers inhibit the TiN crystallite growth tomore than the intended 5 nm thickness or to form long columnarstructures throughout the film thickness. Consequently it is assumedthat no heteroepitactic relation is present between the TiN grains ofadjacent layers. The selected area electron diffraction pattern inset inFig. 2(a) contains strong reflections in the [200] direction, indicatingtexturing of the TiN, as already observed by X-ray diffraction (notshown here). Additionally, satellite reflections originating from themultilayering are visible in the central spot of the diffraction pattern.

The absence of heteroepitaxy across the SiNx layer is confirmed bythe HRTEM image in Fig. 2(b) which shows differently oriented,equiaxed, crystalline nanograins with a diameter of 5 nm composingthe TiN layers separated by regions having lighter contrastcorresponding to amorphous 0.5 nm thick SiNx layers. The FastFourier Transform (FFT) patterns taken from different grains atpositions 1 to 4 confirm the crystalline structure corresponding to thelattice spacing of fcc-TiN. Fig. 3(a) shows a TEM bright field image of across-section of the top part and the interfacewith the substrate of theTiN/SiNx coating indented to a depth of 100 nm. Directly under theindent apex a well-defined, triangularly shaped area is observed, inwhich the multilayering cannot be distinguished any longer as seen inFig. 3(b). The Berkovich indenter tip angle is marked with the whitedashed line. After tilting the sample by about 8° (see Fig. 3c), themultilayer structure is again visible and seems not to be altered: nocompression of the layers was observed. The multilayers, which werevisible in the intact region of the sample before titling, are then out ofvisibility under this angle, because these undeformed multilayers donot lie anymore in the observation direction. In the region of theindent onset several areas with seemingly double SiNx layers, markedbywhite arrows in Fig. 3(d), were observed and are discussed in detailfurther below.

Fig. 4(a) illustrates the cross-section through a 250 nm indent. Theprotective platinum top-coat was accidentally removed by the ionbeam during the preparation of the TEM lamella by FIB and thus thearea directly under the imprint is damaged by the Ga+ ion beam.Nevertheless, the important features, such as open cracks in the lowerpart of the film shown in insets (b) and (c) in Fig. 4, and underneaththese cracks some delamination from the substrate, indicated bywhite arrows in Fig. 4(a) are clearly visible.

The cross-section through the 500 nm deep indent is depicted inFig. 5(a). In addition to the delamination of the film from the substrate,large, completely open lateral cracks at a depth of about one quarter ofthefilm thickness from the substrate are observed. Amedian crackwasgenerated below the indent tip as shown in Fig. 5(a). This crack doesnot run straight but follows the TiN grain boundaries as indicated bythe white arrows in Fig. 5(b). Around this crack some bright lines arevisible, which start from within the SiNx layers, seeming to penetratethe TiN crystallites or to double the SiNx layers. They are also found inthe surroundings of the large lateral cracks at the bottomof the film, asshown by the white arrows in Fig. 5(c).

Fig. 6(a) shows a TEM bright field image of the cross-sectionthrough the TiN/SiNx film indented to a depth of 1000 nm with amedian crack in the Si wafer appearing at the bottom of the image.Conversely to the 500 nm indent, the film indented at 1000 nm was

Fig. 2. TEM micrographs of the undeformed TiN5 nm/SiN0.5 nm film: (a) dark field imageobtained using the (200) reflection with the electron diffraction pattern inset. Thewhite arrow indicated a line of satellite spots around the transmitted spot originatingfrom the multilayering; (b) HRTEM image showing crystalline TiN grains separated byamorphous layers of SiNx with the corresponding FFTs numbered 1–4.

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almost completely delaminated from the substrate and only very tinylateral cracks opened at the film bottom, Fig. 6(b).

The vertical crack at the magnified right onset of the indentimprint in Fig. 6(c) follows a tortuous path along the TiN grainboundaries andwas generated prior to delamination as a consequenceof the high stress originating from the very deep indentation. Thisissue will be discussed further below. No median crack under theindent tip was observed as shown in Fig. 6(d).

Fig. 7 shows again a TEM bright field image: this time it is a cross-section through a 100 nm indent performed, for comparison purposes,in a TiN/SiNx coating having a different layer thickness i.e. TiN=1 nmand SiNx=1 nm.As canbe clearly seen in the top left corner of Fig. 7(a)the multilayering underneath the indent imprint is tilted out ofvisibility conditions. Throughout the whole film thickness column-shaped zones, with hardly visible multilayering are present. A closerlook at these zones, (Fig. 7(b)), reveals that the multilayering is splitthere similarly to the 100 nm indent shown in Fig. 3(d), which will bediscussed in detail in the following paragraphs. Fig. 8(a) and (b) showsmagnified views of the deformed regions in the vicinity of the 100 nmindent where the “double layering” is visible. Parts of the multilayers,which are cleaved parallel to the image plane and shifted by about halfof the TiN layer thickness with respect to the pristine multilayers, areclearly observed. Also edges of split SiNx layers are particularly wellvisible next to the displaced TiN grains.

4. Discussion

The first consequence of the plastic deformation generated byindentation in the TiN/SiNx multilayered sample, which was observedin all samples indented to different depths, is that the multilayerstructure in the area affected by the plastic deformation, i.e. theindenter imprint, is not visible as shown in Fig. 3(b). However, tiltingthe TEM sample by about 8° for the 100 nm indent and 20° for the1000 nm indent around the horizontal axis in plane with the layersbrings the multilayering into visibility conditions, Fig. 3(c), showingthat it was not destroyed by the plastic deformation occurring duringindentation. This is in agreement with other authors, who observed asimilar rotation of the multilayers with respect to the original surface,caused by the plastic deformation of the substrate by indentation, inTiN coatings with large columnar grains [28] and single crystalline TiNcoatings [29]. In both studies the rotation of the crystallographicplanes was 8° around the [001] zone axis (deduced from thediffraction pattern of the single crystalline TiN). Such phenomenonof rotation of the crystallographic planes has been reported previouslyby Molina-Aldareguia et al. [29], and was interpreted as a mean ofplastic yield of the crystalline layer. In our studies the diffractionpattern had a polycrystalline character, due to the fiber texture of theTiN grains; thus it was impossible to determine exactly the degree ofrotation. Therefore the values of 8° for the 100 nm indent and ∼20° forthe 1000 nm indent have been deduced from the tilt of the sampleholder.

The stress upon indentation is mediated by deformation of thesilicon substrate which is much softer (H=10.7 GPa) than the coating(HTiN/SiNx

=18 GPa). As a result large lateral cracks at areas ofmaximum stress concentration located at the bottom of the filmwere induced. The lateral cracks are due to shear stresses exceedingthe shear strength of the interfaces between TiN and SiNx; they wereobserved in all indented samples. From the TEM images, it isimpossible to determine if the fully open cracks are generated insidethe SiNx layer or at the interface between TiN and SiNx. Due to thelower Young's modulus of SiNx (210 GPa) compared to that of TiN(590 GPa) [34] one would expect SiNx to be more elasticallydeformable at the same strain and therefore to be more resistant tocrack formation than the crystalline TiN with its higher Young'smodulus. The parallel cracking occurs at the weakest part of themultilayering which in this case is thus the TiN/SiNx interface.

Fig. 3. (a) Bright field TEM cross-section through the 100 nm deep indent in the TiN5 nm/SiN0.5 nm film; (b) and (c) magnified views of the triangularly shaped area under the indentapex with (b) invisible multilayering and (c) multilayering in imaging conditions after tilting the lamella by about 8° around a horizontal axis in plane with the multilayering;(d) magnified view of the left indent onset. White arrows indicate the “double layering”. See text for details.

Fig. 4. (a) Bright field TEM image of TiN/SiNx film cross-section through the 250 nm deep indent. (b) and (c)Magnified views of the fully open cracks at the bottom of the film in areasof maximum stress concentration left and right from the indent imprint, respectively.

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Fig. 5. (a) General view of the 500 nm deep indent; detailed views of: (b) the 100 nm long median crack generated at the indent tip following the grain boundaries of the TiNnanocrystallites; (c) large lateral crack at the bottom of the film resulting from debonding between SiNx and TiN layers.

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The median crack, induced under the 500 nm indent, Fig. 5(a),propagates perpendicularly to the multilayering, following the TiNgrain boundaries in an apparent zig-zag path. From studies ondeformation of nanocrystalline metals [24] it is known that it isimprobable to induce dislocations into grains smaller than 10 nm. Forthese reasons we assume that no dislocations will be active insidecrystalline grains of 5 nm size such as the TiN grains in this study.Therefore it is easier to generate a crack between two grains, whichthus follows the incidental location of the TiN grain boundariesinstead of a straight crack, which would have to go through the grains.Söderberg [15] observed also in indented TiN/SiN multilayers cracksnucleating and propagating downwards through the film in a zig-zagfashion, following TiN–SiNx interfaces but also crossing TiN layers inthe vertical direction. However in his study, the cracks subsequentlyjoined and formed a larger crack further down in the film.

The median crack directly under the 500 nm indent results fromexcessive tensile stress levels exceeding the fracture toughness, arisingin the coating, due to the strong adhesion to the substrate, Fig. 5(a).Indeed, as long as the film adheres to the substrate during indentation,all observed types of cracks such as themedian crack, the vertical cracksat both sides of the indent onset formed for the 1000 nm indent and thehorizontal cracks at the film bottom, will continue to open. Thisassumption is confirmed by the images taken from the 1000 nm indentthat exhibits severe delamination of the multilayer coating, mediatingshear stress und indentation. These delaminations leave the whole filmvirtually intactwithout cracks longer thana few tensofnanometers. As aresult the cracks parallel to the multilayering found in coatings withlower indentation depths are almost completely closed. Also the verticalcrack underneath the indenter tip is no more visible, most probablybecause the areas adjacent to the onsets away from the indent are liftedup thus causing the closing of the crack underneath the indenter tip. Atexcessively high loads the delamination of the coating becomes thedominant stress relaxation mechanism. This indicates that the filmunder these excessive loads delaminates from the substrate due to thehigh substrate deformation before the lateral cracks can form and splitopen, although the existence of cracks that closed upon stress relaxationcannot be ruled out.

Another feature, which was observed in the TEM cross-sections ofthe deformed films, appears like “double” SiNx layers as evidenced inFig. 3(d), Fig. 5(c), Fig. 6(c) and Fig. 7(b), respectively. This doublelayeringmust be a consequence of the plastic deformation since it wasobserved neither in the TEM images of the as-deposited TiN/a–SiNx

films (cf. Fig. 2a) nor in undeformed areas. For this double layering wepropose an explanation related to the deformation mechanism thatwas not addressed in the available studies.

The double layering visible in the TEM images presented in thispaper, resulting from the indentation process, originates from thedisplacement of pieces of one or several TiN/SiNx layers, whichfracture perpendicular to the layers at TiN/TiN grain boundaries. Thedirect evidence of individual grains undergoing this displacementmechanism, involving fracture at the vertical grain boundariesbetween the TiN grains is shown in Fig. 8(a)–(b), and is furtherreferred to as “shear sliding at grain boundaries” [25]. The cartoon inFig. 8(c) schematically illustrates this situation (the indicated viewingdirection is perpendicular to that on the TEM images in Fig. 8(a) and(b)): a piece consisting of three SiNx layers alternated with two TiNlayers, each containing several TiN grains is shifted downwards byhalf of themultilayer period creating thus a double layering of the SiNx

layers superimposed on the pristine multilayers in the direction ofobservation. The cleaving plane lies perpendicular to the observationdirection and is indicated in the drawn scheme. This “double layering”is particularly well visualized experimentally for the 100 nm indent inFig. 3(d) where it is observed mainly under the indented area and forthe 500 nm indent around the lateral cracks in Fig. 5(c).

The above described mechanism of shear sliding at grainboundaries is also observed in coatings with a different TiN/SiNx

layer thickness, as clearly shown in Fig. 7(a) illustrating a cross-section through a 100 nm indent performed in TiN=1 nm and SiNx=1 nm coating. Seemingly the TiN/TiN boundaries of 1 nm length canalso undergo this shear sliding mechanism, which is obviouslyindependent on the TiN crystallite size, as it occurs as well incolumnar TiN with column heights reaching several hundreds ofnanometers [25,28,30–33] as well as at the nanometer scale down tocrystallite sizes of 1 nm shown here. This means that although in

Fig. 6. (a) TEM Cross-section through the 1000 nm indent: (b) magnified view of thefine lateral cracks along the TiN/SiN interfaces at the film bottom; (c) vertical crackunder the left onset, following the TiN/TiN grain boundaries; (d) the deformed areaunder the indent tip where the multilayering is still visible.

Fig. 7. Cross-section through a 100 nm deep indent performed in a TiN/SiN multilayer witlayering” resulting from cracks at the TiN/TiN interface perpendicular to the multilayering,

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nanoscaled multilayers the amorphous SiNx layers inhibit the growthof large TiN columns, the deformation of nanosized TiN grains underindentation occurs unambiguously by grain boundary shear sliding.

The deformation sequence of the studied TiN/SiNx films, as theindenter load in increased, is summarized below.

At initial indentationdepth to 100 nm,which corresponds to the20%of the film thickness, the first sign of failure is the appearance of adiscontinuity of themultilayering, which results from the displacementof entire pieces of multilayers by about half a period. This is caused bythe shear sliding at theTiN/TiNgrainboundaries and is visible as “doublelayering” in the TEM images due to the overlapping of the broken pieceswith the intactmultilayers. This failuremode is independent on the TiN/SiN layer thickness, as demonstrated for two sampleswith different TiNand SiNx layer thicknesses (see Figs. 3 and 7) and this mechanism isobserved for all indentation depths i.e. 100 nm, 250 nm, 500 nm and1000 nm.

From the indentation depth of 250 nm on, lateral cracks at the sitesof highest stress concentration, located at about a quarter of the filmthickness from the substrate, occur, (cf. Fig. 4). Simultaneously,beginning delamination is observed at the coating–substrate interfacein the direct vicinity of the cracks.

Higher loads i.e. indentation to 500 nm, lead to the generation of amedian crack under the indenter apex, following a zig-zag like pathalong the TiN/TiN boundaries, to the opening of the cracks at thelower half of the coating and to enhanced delamination.

For the 1000 nm indent severe delamination of the multilayercoating is observed and vertical zig-zag like cracks at the indent onsetsare formed. The median crack and the cracks in the lower part of thecoating are not visible anymore, most probably because they wereclosed up due to the delamination, which at this high loads is thedominating stress relaxation mechanism.

h a layer thickness of 1 nm and 1 nm, respectively; (b) magnified view of the “doublewhich seems not to depend on the TiN/SiN layer thickness.

Fig. 8. (a) and (b) Magnified views of deformed regions under the 100 nm indent exhibiting double layering: white arrows show individual grains, which underwent shear sliding atgrain boundaries by approximately half of the TiN layer thickness i.e. 2.5 nm; (c) schematic representation of the deformation mechanism in the TiN/SiNx multilayer systemidentified as collective shear sliding at grain boundaries. A very thin piece consisting of three SiNx layers alternated with two TiN layers, each containing several TiN nanograins wasshifted by half of the multilayer period. Thus the double layering comes into sight when looking in the direction of observation.

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5. Conclusions

Themicrostructure of TiN/SiNxmultilayer coatings before and afterdeep indentation was investigated in order to determine themechanisms of plastic deformation occurring during the indentationprocess. The coatings consisted of equiaxed 5 nm crystalline TiNgrains completely separated by 0.5 nm of SiNx. The first failureevidence after indentation is an apparent “double layering” which isattributed to shear sliding of a few nanometers at TiN/TiN grainboundaries perpendicular to the substrate of pieces of multilayersconsisting of several TiN/SiNx layers. The TEM images show pieces ofmultilayers with a size of about 20 nm that are collectively displacedin the plane of the lamella without being internally destroyed. Thisdestruction mechanism is independent on the multilayer periodicityand indentation depth as it was observed in all indented samples. Theobservation that pieces being several bilayers thick, crack at TiN/TiNboundaries strongly suggests that the presence of the amorphous SiNx

separating layer does not prohibit shear siding at grain boundaries. Atall indentation depth no compression of the multilayers was observedindicating the absence of dislocation activity. Depending on theindentation depth different types of cracks were observed: horizontalones along the multilayering at the TiN/SiNx interface and verticalones under the indentation imprint or at the onset. The cracksperpendicular to the substrate were never straight, independently on

their location with respect to the indent, but always followed atortuous path along the vertical TiN/TiN grain boundaries. Partialdelamination of the coating from the substrate was observed forinitial indentation depths exceeding the 20% of the coating thickness.Higher indentation depths caused increasing delamination whichresults in relaxation of the transient indentation stress, closing of theinduced cracks and hence less deformation damage within thecoating.

Acknowledgements

The authors are grateful to the Centre of Microscopy at EPFL,Lausanne, Switzerland for providing the access to the HRTEM. Manythanks toMr. J.C. Cancio for the deposition of the TiN/SiNmultilayeredsamples.

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