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Applied Surface Science 256 (2010) 7583–7590

Contents lists available at ScienceDirect

Applied Surface Science

journa l homepage: www.e lsev ier .com/ locate /apsusc

robing mechanical properties of thin film and ceramic materials in micro- andano-scale using indentation techniques

ostas A. Charitidis ∗

ational Technical University of Athens, School of Chemical Engineering, 9 Heroon Polytechniou St., Zographos, Athens GR-157 80, Greece

r t i c l e i n f o

rticle history:eceived 24 November 2009eceived in revised form 27 May 2010ccepted 3 June 2010vailable online 9 June 2010

a b s t r a c t

In this study, we report on the mechanical properties, failure and fracture modes in two cases of engi-neering materials; that is transparent silicon oxide thin films onto poly(ethylene terephthalate) (PET)membranes and glass-ceramic materials. The first system was studied by the quazi-static indentationtechnique at the nano-scale and the second by the static indentation technique at the micro-scale.Nanocomposite laminates of silicon oxide thin films onto PET were found to sustain higher scratch

eywords:ractureailureanoindentation

induced stresses and were effective as protective coating material for PET membranes. Glass-ceramicmaterials with separated crystallites of different morphologies sustained a mixed crack propagationpattern in brittle fracture mode.

© 2010 Elsevier B.V. All rights reserved.

anoscratchilicon oxide filmlass-ceramic material

. Introduction

Applying a hard coating to a surface can improve the resis-ance to wear and environmental degradation for mechanical andribological applications. However, there is a high risk of crack for-

ation through the film or failure at the interface between thelm and the substrate. To study failure mechanisms and to char-cterize mechanical properties of coated systems, nanomechanicalechniques such as nanoindentation and nanoscratching have beenccepted as effective experimental methods. The measurementsre usually controlled by the coating material itself, by the sub-trate upon which the coating is deposited, and by the interfaceshich bond the system together. The experimental procedures and

nalytical models to extract the film properties such as hardness,lastic modulus as well as interfacial fracture and fracture tough-ess have been described in several papers [1–3]. Indentation ofoated systems showed that soft substrate yields the formation ofnnular primary cracks, while hard substrate favors the formationf radial corner cracks [4].

The role of nano- and microcracks on mechanical properties areot yet well established. The failure mechanism of coatings during

cratch tests is very complex and depends on many factors, such ashe hardnesses of the coating and the substrate, the coating thick-ess, the scratch indenter radius, the deposition techniques, theubstrate material, and the scratch method [5].

∗ Tel.: +30 210 772 4046; fax: +30 210 772 2339.E-mail address: charitidis@chemeng.ntua.gr.

169-4332/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.apsusc.2010.06.006

The failure of coatings is also determined by the deformation ofboth the substrate and the coating. When the coating is not thickenough, a high compressive stress field caused by the indenter dur-ing a scratch test may be transferred to the substrate, and hencethe scratch response is controlled by plastic deformation of thesubstrate. Coatings with the same hardness may exhibit differentresistances to plastic deformation, and the deformation of a hardcoating is strongly dependent upon the combination of its hardness(H) and elastic modulus (E).

The task of packaging is to protect the packaged goods againstthe impact of the surrounding environment and to prevent the lossof the constituents from the packaged food. In the first case the pen-etration of oxygen and water vapour causes mostly impact, becauseof their important effect on the quality reduction of the packed food.At present, polymeric materials are primarily used in the packagingindustry. Polymeric materials are flexible, lighter, and less materialcan be used for the production of the packaging. Polymer materialsare often also transparent and cost-effective. The common disad-vantage of polymeric materials is their lower barrier propertiesagainst gases and vapours such as oxygen, moisture or differentorganic compounds. At the same time, their barrier properties canbe improved using inorganic barrier layers.

Silicon oxide films have been proven to be useful in improv-ing the gas barrier properties of plastic packaging, often reducing

the permeation of oxygen and water through polymer film by 100times or more [6]. Such composites have several advantages overtraditional metallised polymer films (Al) including transparencyand microwave compatibility. To achieve a long-term barrier to gaspermeation, these coatings must be resilient and durable enough

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o that they can withstand the stresses that are applied to the com-osite during its lifecycle [7]. To this end a number of studies haveeen performed to assess the mechanical properties of these filmsnd the adhesion to the substrate [8–12]. A number of depositionechniques have been used to deposit SiOx films onto polymeric

aterials with the required functional properties, with cost to becritical parameter for industrial use. The development of elec-

ron beam (e-beam) evaporation allows refractory materials likeilicon oxide and aluminium oxide to be deposited. At present theechnique of e-beam evaporation is widely used in manufactur-ng of the transparent barrier films for packaging [13]. Regardinglasma-enhanced chemical vapour deposited (PECVD) SiOx films7–150 nm thick) it has been found that such layers are under min-mal compressive internal stress and that they display excellentdhesion to PET films. Thicker coatings have been found to be morerone to failure than thinner ones at low strain [12] due to the

ncreased probability of large defects being present. Yanaka et al.14] found that for evaporated silicon oxide, thicker coatings haveower tensile strength and are under less compressive strength dueo stress induced cracking.

Nanocomposite laminate silicon oxide films/PET have excellentarrier properties when x is up to 1.8 [15–18]. It is possible thathe silicon dioxide atomic network contains spaces in its struc-ure through which the gas molecules can be diffused [15]. Inontrast, when x < 2, dangling bonds on silicon sites deform the net-ork, connecting and tightening the structure and thus reduce the

as diffusion [15]. SiOx with x ∼ 1 have yellowish tone making itsse in food packaging not favourable. Consequently, the commer-ial layers are produced with an elementary ratio x ∼ 1.8, so thathe barrier properties and full transparency in VIS range can berovided. Furthermore, the film thickness strongly influences thearrier properties. A thin layer, which does not cover completelyhe substrate roughness, cannot guarantee sufficient barrier prop-rties. A very thick layer is, on the other hand, easily breakable.racking and peeling are generally observed for layer thickness150 nm [15,18].

The SiOx/PET barrier system meets the demands regardingcratch-resistance, wettability, biocompatibility, or friction, main-aining the bulk characteristics of PET (such as easy processing,ood mechanical properties, reasonably low permeability to oxy-en and carbon dioxide gases, and good chemical coupling withntibacterial coatings). In addition, the SiOx/PET barrier system issed in flexible organic electronics [19] and flexible solar cells [20].

To study failure modes in hard thin films, SiO1.8 thin films ontoET substrates were subjected to indentation stressing and nano-cratching. This case is of high interest due to the mismatch ofhe hardness and elastic modulus of the SiO1.8 film and PET sub-trate. The objective of the present work is to provide an improvednderstanding of the mechanical response of the SiO1.8/PET system.

Fracture, being the local separation of a piece of material, underhe action of stress, comprises one of the mechanisms of energybsorption, for instance during indentation, impact or wear. Thebility to study the fracture has long been conceived to be a highlyon-trivial issue. The toughness is not a constitutive parameter,ince T-curves and R-curves were discovered. In order to gainnsight into the mechanical behaviour of a material several param-ters, sometimes interrelated, should be considered at the sameime [21,22]. Toughness, strength, brittleness and extend of plas-icity at the crack tip area are few examples of such parameters. Aighly important, if not conclusive, step towards a complete under-tanding of the mechanical behaviour of brittle solids would be the

onnection between the mechanics of fracture with the physicalechanisms of crack growth [23].Cracking can be undesirable when it causes catastrophic failure

ut also necessary for specific applications, such as armour plates.n brittle solids, like ceramics or glass-ceramics, the development of

ence 256 (2010) 7583–7590

microcracking, instead of formation and propagation of long crackshas been found to reduce brittleness and increase plasticity. Duringthis brittle to quazi-plastic transition a long crack is diffused in acumulated damage zone, beneath the indentation area, where amicrocrack cloud is formed mainly due to intergrain microcracking.Such transformation increases the wear resistance of brittle solids[24] and is caused by coarsening of the grain structure [25].

In this study, we report on the fracture modes in micro-scale,namely transgranular and intergranular modes. In the former casethe crack paves its way irrespectively of the direction of the grainboundaries, i.e., the interfaces between the different phases. In thelatter case, the crack preferentially follows them, i.e., debonds theinterfaces [26]. These fracture modes have been extensively stud-ied for ceramic materials [27–29], while they have not receivedthe same attention in the case of glass-ceramic materials. Theintergranular fracture is connected to high toughness, while trans-granular to high strength [30].

2. Experimental

2.1. SiO1.8/PET system

The examined PET membranes Hostaphan® RNK (thickness∼12 �m) were industrially supplied, and had been treated usingmechanical uniaxial stretching that is usually applied in industrialscale. This treatment resulted in the formation of a crystalline-likelayer on the top of the membrane of unknown thickness [31]. Itis a highly transparent, uniaxially oriented, coextruded film withboth sides having a standard topography. A lab scale e-beam coaterA260 E/B1 from Applied films and Leybold Optics GmbH was usedto deposit SiOx films 50 nm thick on one side of the PET web byroll to roll processes. The vacuum chamber of the coater is 550 mmhigh and 350 mm wide. The coater is equipped with a microwavegenerator for the applications of plasma pre-treatment and reac-tive evaporation. The maximum power of the microwave generatoris 2.7 kW. For plasma pre-treatment a variety of gases can be usede.g. O2, N2, NH3 or CO2. The gas inlet is controlled by mass flowcontrollers MKS PR 3000. The evaporator is equipped with a water-cooled rotary mount and crucible plates which can be exchangedseparately. The crucible rotate continuously for uniform depositionof material, however the source material for deposition is againand again heated and evaporated. During the coating process, thedeposition rate, which is the amount of the evaporated material intime, is monitored by an Inficon deposition controller based on theoscillating quartz crystals.

The vacuum system consists of the water-cooled rootspumpFM 12S and turbomolecular pump–Turbovac 1000 with pumpingspeed ∼1000 l/s (N2). The ultimate pressure of the turbomolecularpump is <1.10−6 mbar (>1.10−4 Pa) in 40 h of pumping. The typicalpressure during deposition process is 1 × 10−2 Pa. Three pressuremeasuring devices are installed in the system: two Termovac TM210 S with the pressure measurement limit from 10 Pa down to10−1 Pa and Ionivac IM 210, which is available for pressures in rangefrom 10−1 Pa down to 10−7 Pa and are used for the measurement ofthe pressure in the deposition chamber and they also support theautomatic operation of the whole vacuum system.

The winding system is adapted for the roll to roll transport ofPET at speed up to 7 m/min. The system comprises of six rollerswith different diameters. The three main rollers are motor driven,the additional rollers are powerless. The web speed is controlled bythe rotational speed of the coating roller (water cooled) and both

winding directions are possible. The outer roller diameter of thewindings rollers is typically 150 mm. Their web width is 280 mmand the length of deposited film can reach 1000 m [32].

Hardness (H) and reduced elastic modulus (E) of the coatingswere assessed by means of a nanoindentation system (Hysitron

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riboLab system). A pyramidal diamond Berkovich tip with a totalncluded angle of 142.3◦ and a radius of curvature of approxi-

ately 100 nm has been used for both indentation and scratcheasurements. A trapezoidal loading–unloading ramp was used.

he maximum load is reached after 5 s, kept constant for 5 s andecreased to zero during the same time duration. The H andhe E were determined from the experimental load–displacementurve using the Oliver and Pharr model [33] and the valuesrom more than ten indents were averaged for each sample.cratch measurements were conducted using a two-dimensionalorce–displacement transducer. This transducer allows the mea-uring of both lateral force and normal displacement. The totalcratch length is fixed at 10 �m. Scratch measurements have beenealized at a normal force of 10,000 �N and a constant normal load-ng rate of 33.33 �N s−1. The nanoindentation system employed inhis study is equipped with a Scanning Probe Microscope (SPM), inhich a sharp probe tip is moved in a raster scan pattern acrosssample surface using a three-axis piezo positioner. Tests were

erformed in a clean-air environment with a relative humidity ofpproximately 25%, while the temperature was constantly kept at0 ◦C (ambient temperature). The samples to be measured werelaced into the cabinet for several hours prior to testing, in ordero reach equilibrium with the thermal mass inside the cabinet.

The reported thickness of the SiOx films was measured by an X-ay fluorescence meter which had been calibrated by transmissionlectron microscopy. The surface morphology of the evaporatediOx (x ∼ 1.8) film was examined using the techniques of scanninglectron microscopy (SEM) and SPM. The atomic ratio x, which ishe ratio of Si to O, was determined by X-ray Photoelectron Spec-roscopic analysis (XPS) [32]. For the thickness of the SiO1.8 filmsxamined in this work the atomic ratio (1/1.8) is uniform.

.2. Glass-ceramic materials

Glass-ceramic materials have been produced by the thermalreatment of two vitreous products. IS1 product contains 50 wt% ofrich in PbO and Fe2O3 incinerated sludge, 35 wt% SiO2 and 15 wt%a2O. IS2 product contains 60 wt% of the same incinerated sludge,5 wt% SiO2 and 15 wt% Na2O. The resulting glass-ceramic mate-ials have different final morphologies. IS1 glass-ceramic is madep of a vitreous matrix with homogeneously dispersed rectangu-

ar shaped Pb8Fe2O11 crystallites. IS2 glass-ceramic is made up ofvitreous matrix with inhomogeneously dispersed oblong Fe2O3

rystallites of varying widths [34,35].The modes of fracture were determined by observing the paths

f crack propagation on the surface of the specimens with respect tohe crystallite/amorphous matrix interfaces by optical microscopy.ll indentation induced cracks were produced on flat and polishedurfaces of the thermally treated products with a Vickers dia-ond indenter. This was necessary since an indentation on a rough

urface cannot be observed clearly by high magnification opticalicroscopy, due to the particularly narrow depth of focus. Polishingas achieved by wet mechanical grinding with SiC grinding papers

nd alumina (Al2O3) pastes. An Anton-Paar MHT-10 microhard-ess tester was utilized [36] attached on a Zeiss Axiolab-A opticalicroscope. The indentation parameters were set at: load = 1–3 N,

uration = 10 s and loading rate = 0.2 N/s. The duration of 10 s is aypical time, that is, long enough for the onset of plasticity, whilehe load range was selected in order to produce indentations withlearly defined radial cracks.

. Results

.1. SiO1.8/PET system

Fig. 1 shows three representative indentation curves for max-mum normal loads approximately equal to 18, 90 and 450 �N.

ence 256 (2010) 7583–7590 7585

The load hysteresis and residual penetration depths indicate thatelastic deformation occurred at the low load of 18 �N whileelastic/plastic deformation occurred with loads 90 and 450 �N.Unloading produces a highly nonlinear load response characteristicof a viscoelastic material, for the normal load of 450 �N.

Both H and E are increased in the surface/near surface region(surface region: 0.47 and 5.70 GPa, SiO1.8 film 50 nm thick). Forhigher contact depths, i.e. >350 nm, the H(E) values tend to decreaseapproaching the bulk PET H and E values (0.3 and 3.0 GPa, respec-tively). The bulk H and E values of PET are in good agreement withmeasurements in the literature [31,37–39].

Fig. 2 shows the variation in hardness and modulus with pen-etration depth (displacement) calculated from 10 indentations ineach displacement, covering the loading range 18–1000 �N, intoSiO1.8/PET system. Fig. 2 shows a pronounced rise in hardness andmodulus as the indentation depth decreases. Both H and E areincreased in the surface region (0.47 and 5.70 GPa) and the nearsurface region (i.e. 0.39, 6.65 GPa), while for higher penetrationdepths, corresponding in the bulk of the PET membranes, the H(E)values tend to decrease approaching the bulk H and E values (0.3and 3.0 GPa, respectively).

In all the conventional indentation data reported in this paper,the loading data have been fitted to the power-law function P = ˛hn,where P is the load, ˛, a material parameter and n, the index of thedeformation (indentation index), to determine this depth offset.

The fit of the loading curve data to the function P = ˛hn is verygood. Fig. 3 shows a typical power-law fit to 18, 90 and 1000 �Nindentation into the SiO1.8/PET system; the fitted curve superim-poses virtually exactly over the loading data. When a Berkovichindenter is used, the value of n is often very close to 2 [40–42].However, deviations from this are known to occur, particularly forpolymers (e.g. due to strain-rate effects), and are discussed in moredetail in the following paragraphs.

In Fig. 4 indentation index vs. Displacement of the indenter inSiO1.8/PET system is presented. Fig. 4 reveals the existence of: (i) asurface region of about 50 nm (thickness of the SiO1.8 film) wherethe indentation index is∼0.96, and (ii) a near surface region of about250 nm in thickness, where the indentation index rises from 1.04 to1.60. For penetration depths above ∼350 nm, the indentation indexis nearly stable (n = 1.60). These results are in agreement with aprevious study [31], with the indentation index values somewhatlower than it was measured in [31] due to the fact that in our casethe PET membranes are stretched uniaxially.

The increase in hardness in surface/near to the surface is anindentation size effect and represents a change in the mechani-cal properties of the surface and near surface region due to SiO1.8film and enhanced crystallinity of the near surface PET region.

Mechanical stretching resulted in the elongation of the poly-meric membranes and in the reduction of their thickness, andcaused a preferable orientation of the macromolecules close to thesurface. Such treatments cannot be fully controlled to achieve thedesired degree of ordering and the thickness of the formed orderedoverlayer (so called crystalline layer) [42].

Several factors can affect whether the normal relationship[19,20] between load and depth during an indentation (P = ˛hn,n = 2) holds. Deviations from this (i.e. n < 2) can occur when (i) thereis strain-rate hardening, (ii) there are viscoelastic effects, (iii) themechanical properties (hardness and modulus) vary with depthdue to a changing crystallinity profile. It is clear that viscoelasticeffects are important for PET since n is always <2 in our tests. Vis-coelastic effects can be also inferred from the observed variation in

the indentation exponent, n, with loading rate.

It has previously been suggested that PET films are more crys-talline in the near surface region than in the bulk [42], possibly dueto orientation-induced crystallization. In view of this, the conclu-sion that industrial treatment processes applied on PET before its

7586 C.A. Charitidis / Applied Surface Science 256 (2010) 7583–7590

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ig. 1. (a) Evolution of the applied load vs. displacement upon nanoindentation canoindentations for load 450 �N.

ommercial disposal [stretching (drawing) the films] have a greaterffect of topmost layers, seems reasonable, although clearly, careust be taken in interpreting the near surface indentation response

f viscoelastic–plastic materials.Scanning electron and atomic force microscopes (SEM, AFM)

ere used to investigate the structural integrity of the evapo-ated SiOx films. Extensive SEM examination of the barrier coatingsevealed no obvious pinholes or macro-scale defects in the struc-ures. SEM and AFM investigations showed that films were uniformnd completely covered the PET substrate.

Adhesion characteristics of the SiO1.8/PET were examined usinghe nanoindenter under the applied loads. Arai et al. [43] exam-ned the failure modes of thin films following their indentation andrgued that their adhesion characteristics could be categorized intove general groups depending on the presence (or not) and appear-

Fig. 2. Hardness and elastic modulus vs. displac

or loads 18, 90 and 450 �N, (b) Scanning Probe Microscopy (SPM) imaging of the

ance of the radial or lateral cracks around the indentation. Crackingand flaking are among the possible failure modes of the thin films.For loads used in this work, the failure mode observed (by SPM)was characterized by plastic deformation only (Fig. 1b). The lack ofcracking or flaking in this failure mode indicates that the SiO1.8 filmhas excellent adhesion to the PET substrate.

Friction is a complex phenomenon, which involves asperityinteractions involving adhesion and deformation (plowing). Fig. 5shows the coefficient of friction as a function of normal loadobtained from nanoscratch tests performed with a maximum nor-

mal load of 10,000 �N for SiO1.8 50 nm thick onto PET membranesand uncoated PET. The coefficient of friction increased with increas-ing normal load to a maximum value (∼0.2) corresponding to thetransition load (∼3000 �N). The behavior of coefficient of fric-tion of the uncoated PET, in the same load range, was found to

ement (solid lines are guide for the eye).

C.A. Charitidis / Applied Surface Science 256 (2010) 7583–7590 7587

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ig. 3. Typical loading behavior of SiO1.8/PET system during nanoindentation in the l= ahn , n: indentation index at different displacements of the indenter in SiO1.8/PET

ecrease (from a maximum value ∼0.6) to a value of ∼0.3. Theohesive strength of the SiO1.8/PET system is enhanced comparingo that of uncoated PET, and as a result, the friction force needed tovercome the adhesion of the system will increase, leading to thencrease of coefficient of friction. The high-load range for SiO1.8/PET

eveals higher coefficient of friction (� ∼ 0.2 to 0.3). Taking intoccount that the deformation of SiO1.8/PET was essentially elasticnd elastic–plastic for low normal loads (Fig. 1) the low coefficientf friction in this regime is considered to be a manifestation of sur-

ig. 4. Indentation index (n) vs. displacement of the indenter into SiO1.8/PET system.

nge 18–1000 �N. The solid lines represent fitting of the load data using the equationm.

face adhesion forces. The scratches produced on SiO1.8/PET at loads>3000 �N indicated that plastic deformation was dominant andthat the higher coefficient of friction in this regime was mainly dueto the plowing process occurring at the nanometer and micrometerscales.

3.2. Glass-ceramic materials

Fig. 6 depicts an optical micrograph of the path of anindentation-induced radial crack. The indentation was made on theIS1 glass-ceramic product. This micrograph depicts a mixed crackpropagation pattern. In places indicated by the arrows numberedby 1, the crack is deflected by the crystallites. More specifically,the crack propagates following the boundaries of the separatedcrystallites causing interfacial debonding, i.e., by separating theamorphous from the ceramic phase. In places indicated by thearrows numbered by 2, the crack propagates in straight line as itcuts through the microcrystallites, without being deflected by thecrystallite/amorphous matrix interfaces. Consequently, the frac-ture mode is both transgranular and intergranular for IS1 product.

In IS2 product, the morphology is different, with randomly ori-ented elongated crystallites of various widths dispersed in theremaining glass matrix. Whether these crystallites will act as abarrier to crack propagation depends on their width and relativeorientation with respect to the direction of the crack propagation.

In Fig. 7 all radial cracks, except for the lower one, follow mixedmode propagation. The cracks follow for a short distance the crys-talline/glass matrix interface (intergranular mode) and then theycut through it (transgranular mode). This interchange between thetwo propagation modes is most clearly observed in the upper crack,

7588 C.A. Charitidis / Applied Surface Science 256 (2010) 7583–7590

Fig. 5. Coefficient of friction profile, as a function of the normal load applied for (a) SiO1.8/PET and (b) uncoated PET.

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here the crack path has a step-like appearance. The lower cracknds a thicker elongated crystallite and is been annihilated. Thepper and left side cracks were eventually annihilated from thickerlongated crystallites.

ig. 7. When the microcrack propagates perpendicularly with respect to the needle-ike crystallites the cracking occurs in a transgranular manner. When the microcrackropagates at 45◦ with respect to the needle-like crystallites the cracking occurs inmixed manner (both trans- and inter-granularly). The crack itself has the shape ofbroken line.

aterial. The microcrack propagation is both inter- and transgranular.

4. Discussion

4.1. SiO1.8/PET system

Gas and water vapour transport in gas barrier films is comprisedof contributions from three components: un-hindered transportthrough ‘macro-defects’ (>1 nm) in the oxide layer, hinderedtransport through ‘nano-defects’ (<1 nm), and hindered transportthrough the amorphous lattice of the oxide (interstice <0.3 nm).The relative contribution of each component depends on the sizeof the permeant molecule (or atom), and the number and sizeof each class of defect. The term ‘macro-defects’ is based on thedefect size being three or four times as large as a typical permeantatom/molecule (0.2–0.3 nm). In this case, the defect will providelittle resistance to gas flux. Boundaries between grain-like struc-tures observed in SiOx coatings would correspond to macro-defectsor nano-defects depending on whether they were larger or smallerthan 1 nm [44]. The role of nano-defects is pronounced for relativelylarge molecules (e.g. O2 and H2O) since their lattice permeabilitiesat room temperature are virtually negligible.

Results from the nanomechanical and nanotribological char-acterization of the SiO1.8/PET revealed significant differences inthe nanomechanical properties of the surface/near surface region.SiO1.8/PET system exhibited higher hardness, elastic modulus andelastic recovery than uncoated PET membranes under the samenormal load, implying that SiO1.8 films sustain higher scratchinduced stresses and are effective as protective coating material. In

addition, the lack of cracking or flaking in the SiO1.8 film, indicatedits excellent adhesion to the PET substrate. Finally, the thickness ofthe near surface region of PET (∼150 nm) is more crystalline thanin the bulk [42] which is possibly due to the orientation-inducedcrystallization.

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The SiO1.8/PET system revealed a transition from elastic to plas-ic deformation accompanied by a change of the dominant friction

echanism from adhesion to plowing, observed with increasingormal load. The results of the present work illustrate that signifi-antly different friction properties can occur at different ranges ofoads.

.2. Glass-ceramic materials

For both IS1 and IS2 glass ceramics, the mode of fracture isixed i.e. both transgranular and intergranular. A change in the

oughness of the crystallite/amorphous matrix interface cannot bepplied for the change between the two modes of fracture. This isecause the composition of the amorphous matrix and the sepa-ated crystallites is fixed for all IS1 and IS2 samples. The two aboveases, although they present similar behaviors, should be treatedeparately in order to reach a plausible interpretation.

In the case of IS1 glass ceramic, the reason for the mixed modef fracture is possible to stem on the size of the crystallites. It haseen observed in the same glass ceramic a predominant transgran-lar mode of fracture for crystallites of sizes lower than one micron45]. For sizes higher than several microns, a predominant inter-ranular mode of fracture has been observed. The interpretation forhis behavior has been attributed to the relative sizes of a distortedrystalline zone that acts as a boundary between the crystallite andhe amorphous matrix and the size of the crystallites. In the case ofS1 glass ceramic the crystallites have sizes between 2 and 5 �m,.e. in a region where no predominant mode of fracture is expected.uch a strength degradation with coarsening of the grain structureas been reported previously for ceramic [46] and glass-ceramic47] systems. The interpretation for this common behavior thoughs different in the present work.

In IS2 glass ceramic, the crack propagation mode depends onhe orientation of the crack with respect to the longer dimensionf the crystallites and the elongated crystallites’ width. When therack propagates parallel to the longer dimension, it follows therystallite/amorphous matrix interface, i.e. the propagation is inter-ranular. When the crack propagates at right angles it ruptureshe crystallites, i.e. the propagation is transgranular. This is mostrobably due to their relatively small width, namely below 1 �m48].

In Fig. 2, the orientation of the crack is approximately at 45◦

ngle, therefore it lies at the centre between the above two extremeases. In this case, a mixed crack propagation mode is observed.n order to interpret the above observation we followed the con-lusions of the analysis of a semi-infinite wedge loaded crack thatpproaches an interface at an oblique angle by Ming-Yuan He andutcinson [49]. The analysis deals with the problem of the compe-

ition between penetration through the interface or deflection intot. By taking energetic considerations, i.e. ratio of energy releaseate of deflected crack to maximum energy release rate of pene-rating crack He and Hutchinson came into conclusion of what isntuitively expected: the competition between deflection and pen-tration becomes more favorable to deflection the more obliques the crack impinging the interface, i.e. what is observed in theresent work.

Intergranular mode has been proved to enhance the toughnessf a material [50], since easier interfacial decohesion provides andditional energy dissipative mechanism [51]. Although the inter-ranular cracks become longer, they are easily deflected by thenterfaces and as a result they are spatially located with respect

o a straight transgranular crack.

On the other hand, transgranular mode is favourable when theesulting products are designed to have higher toughness. Dueo the mixed fracture mode both IS1 and IS2 glass ceramics arexpected to show mechanical performance that is between a per-

ence 256 (2010) 7583–7590 7589

formance of a predominantly strong and predominantly toughmaterial.

The above discussion is based on the presumption that theunder study cracks belong to the “long crack regime”, i.e. they areclearly distinguished radial that have emanated from the cornersof Vickers indentation prints and have created and propagated out-side the contact area and have occurred in a brittle manner. In allthe indentation tests made there was not any case that a brittleto quazi-plastic transition was observed [52,53]. As a result, thetoughness mentioned above implies the long crack toughness. Thisis the reason for the interconnection between long straight (trans-granular) cracking and low toughness.

5. Conclusions

The deviations in nanomechanical properties of the surface/nearsurface region are clearly defined through nanoindentation andnanoscratching measurements. The thickness of the near surfaceregion of PET (∼150 nm) exhibits greater crystallinity than this ofbulk PET, possibly due to the appearance of orientation-inducedcrystallization. Under the same normal load, SiO1.8/PET systemrevealed higher values of hardness, elastic modulus and elasticrecovery than uncoated PET membranes, showing that SiO1.8 filmssustain higher scratch induced stresses and are effective as protec-tive coating material. The lack of cracking or flaking led to excellentadhesion of the SiO1.8 film to the PET substrate. The transitionof the SiO1.8/PET system from elastic to plastic deformation wasaccompanied by a change of the dominant friction mechanism fromadhesion to plowing (increase of normal load).

For both glass-ceramics studied in this work, the mode of frac-ture is mixed i.e. both transgranular and intergranular. The understudy cracks are clearly distinguished radial (“long crack regime”).Those cracks have emanated from the corners of Vickers indenta-tion prints, propagated outside the contact area and have occurredin a brittle manner. For all indentation tests performed, no case ofa brittle to quazi-plastic transition was observed. Thus, the tough-ness mentioned above implies the long crack toughness, being thedominant factor for interconnection between long straight (trans-granular) cracking and low toughness.

Acknowledgements

The author thanks Dr. Klaus Noller, Head of Department Mate-rials Development, Fraunhofer Institute Process Engineering andPackaging, Freising, for providing the silicon oxide thin films ontopoly (ethylene terephthalate) membranes.

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