Comparison of test methods estimating the stiffness of ultrathincoatings

Marcus Vinıcius Tavares da Costa , Cristian Neagu, Pierre Fayet, Urban Wiklund, Hu Li, Klaus Leifer,

E. Kristofer Gamstedt

� The Author(s) 2018

Abstract A key engineering parameter of thin coatingsis their stiffness. Stiffness characterization of ultrathincoatings with a nanometer scale thickness is experimen-tally challenging. In this work, three feasible methodshave been used to estimate the Young’s modulus ofmetal coatings on polymer films. The methods are: (1)nanoindentation, (2) strain-induced elastic buckling and(3) peak-force measurements integrated in atomic forcemicroscopy. The samples were prepared by atomic layerdeposition of TiO2 (6 and 20 nm thick) and mixed oxidesof TiO2 and Al2O3 (4 and 20 nm thick). The differencesin estimated Young’s modulus are interpreted in termsof the underlying assumptions and test conditions. Theirspecific advantages and drawbacks are also comparedand discussed. In particular, the nanoindentation neces-sitates a sufficiently sharp indenter tip to make localized

measurements dominated by the coating. The strain-induced elastic buckling method is simple in practice,but showed a large scatter due to variation in localcoating thickness and irregular deformation patterns.The stiffness characterization using atomic force micro-scopy gave the most consistent results, due to a sharp tipwith a radius comparable to the thinnest coatingthickness. All methods gave a higher Young’s modulusfor the TiO2 coating than for the mixed oxide coating,with a variation within one order of magnitude betweenthe methods.

Keywords Nano coating, Young’s modulus, Polymersubstrate

Introduction

Carton food packages are indeed an important inven-tion. Every day, billions of liters of water, milk, juiceand other liquid foods are consumed around the world.Carton packaging is mostly used to contain and protectbeverages so that they can be transported to theconsumers with increased lifetime. The barrier func-tion is today mainly assured by aluminum foils. Thismaterial protects food from the outside (intrusion ofoxygen, humidity, etc.), as well as prevents leakage bykeeping the nutrients inside.1 One of the potentialcandidates which can replace aluminum foil is a thinmetal oxide brittle coating deposited on a polymer filmsubstrate.2 A number of material properties are typi-cally of interest in coatings development. One of theseproperties is the stiffness, which affects the strain andstress state in coating structures. It is needed in theestimation of interfacial fracture toughness in coatingssubject to mechanical loading.

There are some well-known methods to characterizecoatings on polymers such as the fragmentation test.3

This test can be used to calculate the interfacial shear

This paper was presented at the 13th Coatings Science Inter-national Conference (COSI) on June 26–30, 2017, in Noordwijk,The Netherlands

M. V. Tavares da Costa, E. K. Gamstedt (&)Division of Applied Mechanics, Department of EngineeringSciences, Uppsala University, Box 534, 751 21 Uppsala,Swedene-mail: kristofer.gamstedt@angstrom.uu.se

M. V. Tavares da Costae-mail: marcus.tavares@angstrom.uu.se

C. NeaguTetra Pak AB, DSO Packaging Materials, Ruben RausingsGata, 22186 Lund, Sweden

P. FayetAdhemon Sarl, Thin Technology, Avenue Edouard-Dapples20, 1006 Lausanne, Switzerland

U. Wiklund, H. Li, K. LeiferDivision of Applied Materials Science, Department ofEngineering Sciences, Uppsala University, Box 534,751 21 Uppsala, Sweden

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https://doi.org/10.1007/s11998-018-0085-0

743

strength between coating and polymer substrate, basedon the observation of crack density of the coating, asan increasing tensile strain is being applied. Analyticalmodels have been developed to determine adhesiveproperties of the interface from the crack density atsaturation.4 Finite element simulation has also beenused for the coating/substrate assemblies to determinethe cohesive fracture toughness of the interface.5,6 Touse these analytical or numerical approaches, it is firstnecessary to determine the elastic properties of thecoating. Most coatings are assumed to be isotropic, andthe foremost parameter is then the Young’s modulus,E, since the Poisson ratio, m, does not show as muchvariation and has less influence on adhesive behavior.The determination of coating Young’s modulus is ofparticular interest for thin coatings, where it is knownto depend on the coating thickness.7

The focus of this work is therefore to investigate theYoung’s modulus of thin metal oxide coating usingindependent methods. Three experimental methodshave been used to obtain values of the Young’s modulus,namely nanoindentation, mechanical buckling measure-ment and Peak-Force Quantitative NanomechanicalProperty Mapping technique (PF-QNM). Since theultrathin coatings (thickness in the order of 10 nm) donot allow for stiffness characterization by in-planetensile testing, all three methods are based on localout-of-plane deformations. Using three complementarycharacterization methods for the same material alsogives the opportunity to compare the methods inpractice and identify the different advantages anddrawbacks of the methods for the present class ofcoating materials. Furthermore, this work also seeks todetermine the mechanical properties of coating alterna-tives to aluminum and their thickness dependency.

Materials and methods

Samples

The materials in this study were selected based on theirbarrier function for carton food packages. All coatingswere deposited by atomic layer deposition (ALD) onpolymer film substrate. The materials of these coatingsare TiO2 (titanium oxide) and MOX (mixed oxide,composed of even contributions of TiO2 and Al2O3).The thickness of these coatings was 20 nm for thethicker coatings of both materials, 4 nm for the thinMOX coating and 6 nm for the thin TiO2. Thethickness was determined by the number of depositioncycles since only one atomic layer is produced for eachcycle. The substrate, biaxially oriented polyethyleneterephthalate (BoPET) film 120-lm-thick Teijin Meli-nex ST504 from Dupont, was the same for all barriercoatings. The thickness of the coating is negligiblecompared with the thickness of the substrate, and thein-plane stiffness of the coated film is therefore entirelydominated by the stiffness of the substrate.

Tensile testing

The Young’s modulus of the substrate was obtained bytensile test from the slope of the initial linear regime ofthe stress–strain diagram. Three tests were carried outfor each barrier coating. The influence of the ultrathincoating on the film stiffness was substantially smallerthan the measurement scatter. The samples were cut insmall pieces with dimensions 4 mm 9 20 mm usingscalpel and ruler. The samples were stretched up to40% of nominal strain, the speed was 0.1 mm/min, andthe gauge length was > 10 mm. All data were mea-sured by the Deben ‘‘Microtest’’ tensile stage, and thescatter was taken as the standard deviation of the testresults.

Nanoindentation

The main principle of this technique is the nanoscaleindentation of a rigid indenter into the surface of adeformable material. From the load–deflection curveof the unloading phase, the elastic properties close tothe surface can be obtained.8 The CSM Ultra NanoHardness Tester with a diamond cube corner indenterof 40 nm tip radius was used for all ALD samples. Atthe time of testing, this was the sharpest tip availablefor this particular nanoindentation device. The inden-tation depth was up to 30 nm. The loading rate waschosen to be 100 lN/min, as a compromise to be slowenough to obtain stable data and fast enough to avoidviscous effects. A holding time of 10 s before unloadingwas chosen to let the material set, and thus avoidviscous effects in the unloading curve.9 The sample sizewas around 5 mm 9 5 mm, and 10 indentations weremade for each sample from which the average Young’smodulus of the coating was calculated. Each specimenwas attached using double-sided adhesive tape on topof the horizontal sample holder. The influence of theadhesive tape was neglected since the indentationdepth was extremely shallow. The contact stiffness wasexperimentally measured from the initial portion ofunloading curve10,11

S ¼ dP

dh

� �max

¼ mPmaxðhmax � hpÞ�1 ¼ b2ffiffiffip

p Er

ffiffiffiffiA

p

ð1Þ

where Er is the reduced modulus, A is the projectedarea, hmax � hp is the total penetration depth, Pmax isthe maximum force, and m and b are constants whichdepend on the geometry of the indenter. The value of bwas set to 1.0 for a hemispherical indenter,11 and theprojected area A and constant m were experimentallycalibrated by the nanoindentation device. Thecalibration of A accounts for the assumed sphericaltip. Once the reduced modulus Er is known, theYoung’s modulus of the coating, Ec, is obtainedthrough9–12

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744

Ec ¼1 � mc

1Er� 1�mi

Ei

ð2Þ

where Ei is the Young’s modulus of the indenter, 1145GPa. The Poisson ratios of the indenter and coating,vi and vc are 0.070 and 0.30, respectively, for thediamond indenter8 and titanium and aluminum oxidecoatings.13

Strain-induced elastic buckling instabilityfor mechanical measurement

The phenomenon of elastic buckling instability of asuperficial coating subjected to in-plane compressivestrains can be instrumental in estimating the stiffness ofthe coating. This kind of test is usually abbreviatedSIEBIMM in the literature, denoting strain-inducedelastic buckling instability for mechanical measure-ment. The principle is to measure the buckling wave-length caused by an applied or residual compressivestrain on coating/substrate–film. The analytical model-ing typically requires that (1) the stiffness of thecoating is much higher than that of the substrate,Ec � Es, (2) the substrate is thicker than the coating,hs � hc, and (3) the buckling takes place in the smallstrain regime. Figure 1 shows schematically how thebuckles appear as well as the important dimensions inthe cross section. In our case, a tensile strain wasapplied and, due to the Poisson effect (quantified by mc

and ms ), the elastic buckling forms by an effectivelateral contraction force F.

A practical approach has been developed by Volyn-skii et al.,14 where they reported the mechanism ofbuckling formation in the gold layer of 10 nm thick onPET substrate and derived the relation for the wave-length

k ¼ 2phc

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffið1 � m2

s ÞEc

3 1 � m2c

� �Es

3

sð3Þ

as a function of the thickness of the coating, hc, theYoung’s modulus and Poisson’s ratios of the coatingand the substrate using Euler derivative equation.

Later, Stafford et al.15 used this equation to deter-mine the Young’s modulus of the coating using adesigned tensile stage with laser diffraction to measurethe wavelength of the buckled coating. Solving for Ec

gives SIEBIMM main equation

Ec ¼3

8p3Es

1 � m2c

1 � m2s

khc

� �3

: ð4Þ

A state of plane strain is assumed. In our case, thePoisson ratios mc and ms are 0.3013 and 0.44,1 respec-tively. Equation (4) was also used, e.g., by Cranstonet al.16 to determine the Young’s modulus of thin films(� 70 nm) of nanofibrillated cellulose multilayersusing a tensile microtest stage together with scanningelectron microscopy (SEM). Such studies show thatSIEBIMM may be effective for elastic characterizationof thin layers.

It should be noted that equation (4) is applicableonly for small deformations. Jiang et al.17 proposed acorrection of equation (3) for the case of finite defor-mation in buckling due to substrate prestraining. Byuse of their approach, a corrected wavelength k0 can becalculated

k0 ¼ kð1 þ eTÞð1 þ nÞ1=3 ð5Þ

where

n ¼ 5eT 1 þ eTð Þ=32:

The compressive transverse strain, eT ¼ �msek, isdetermined from the applied longitudinal tensile strainat buckling onset ek. The buckling strain ek and thewavelength k are obtained experimentally.

In this study, four SIEBIMM tests were carried outin situ in a Deben ‘‘Microtest’’ tensile stage installed ina Hitachi TM-1000 table-top SEM operating at 15 kVwith a backscattered electron detector. For highmagnification, the specimen was subsequently broughtto a Zeiss Merlin field emission gun SEM used at anacceleration voltage of 2 kV and the high-efficiencysecondary electron (HE-SE2) detector suited fordetailed topographic analysis. The considered coatingswere made from nonconductive materials. Conse-quently, unwanted charging effects were observed inthe SEM during the first attempts. In order to reducethe charging effects, all these samples were coated ontop of barrier layers with Au–Pd conductive coating byQuorum SC7640 polaron sputter coater.

The effects of the Au–Pd coating on the mechanicalexperiments are presented in Appendix A. In thesetests, the coating properties were examined in thelateral direction, measuring the buckling size beforeand after the conductive layer, and longitudinal direc-

Crack

ε

ε

FF

E v hC , C , C

E v hS , S , S

λ

Fig. 1: Buckling formation due to an effective compressiveforce (Poisson effect) in a coated film caused by an appliedtensile strain

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tion, observing the cracking progress regarding tensilestrain in different Au–Pd thicknesses. Essentially, itcould not be demonstrated that the conductive coatinghad any influence on the estimation of coating stiffness.

Figure 2 illustrates how the SIEBMM test wasperformed. A removable rig was designed and manu-factured to fit the Deben tensile stage in the vacuumchamber of the table-top SEM, as shown in Fig. 2a.When the buckling was clearly observed in situ duringtensile testing in the SEM, the loading was stopped andthe applied strain in the sample was fastened bytightening two fixation screws in the removable rig, asshown in Fig. 2b. The compact rig was then detachedand taken to the high-resolution SEM. The removablerig can thus swiftly be taken out from the microtensilestage and moved to a high-resolution SEM whilemaintaining the tensile strain at buckling onset ek, asshown in Fig. 2c.

AFM peak-force measurements

The Young’s modulus of the coatings was alsoestimated using a Multimode 8 Atomic Force Micro-scope, Bruker, in the Peak-Force QuantitativeNanomechanical Mapping (PF-QNM) mode.12,18 Asilicon probe was utilized with the tip radius of 8 nm.PF-QNM mode is a technique to quantitatively mea-sure the surface mechanical properties of the sampleby recording the force–separation feedback loop in theAFM as illustrated in Fig. 3. The Young’s modulus, theadhesion force, energy dissipation and the sampledeformation can be extracted from this loop, where theYoung’s modulus can be estimated by fitting theretraction part of the force–separation loop to theDerjaguin-Muller-Toporov (DMT) model19:

F � Fadh ¼ 4

3EAFM

r

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiRðd� d0Þ3

qð6Þ

where F � Fadh is the difference between the force onthe cantilever and the adhesion force, R is the tipradius, EAFM

r is the reduced modulus, and d� d0 is thesample deformation. Just as for nanoindentation, theYoung’s modulus of the coating is calculated from thereduced modulus by equation (2).

Results and discussion

Tensile testing

The in-plane tensile test confirmed that there was noinfluence of the coating stiffness on the film stiffness,which was effectively dominated by the substrate. Theobtained value of the Young’s modulus of the substratewas comparable for all samples, and no tangibledifference was found compared with the same sub-strate material without any barrier coating.20 That also

Fig. 2: Steps in the elastic buckling instability analysis: (a) The removable rig attached in tensile stage. (b) The fixation rigwhich is dismounted at the point of buckling formation. (c) The rig is brought to a high-resolution SEM for measurement ofthe buckle dimensions

0

Fadh

DeformationForc

e

Tip-sample separation

DTM fit for elastic modulus

Fig. 3: Force–separation curve generated in AFM PF-QNMmode

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shows that the ALD manufacturing process did notchange the elastic property of the polymer substrate.Figure 4 shows the results from tensile test.

Nanoindentation

The contact stiffness was determined by least squaresfitting from unloading curves starting at 98% down to40% of the maximum measured force. With a holdingtime of 10 s before unloading, obvious viscous effectscould be avoided. Figure 5 shows the results from the

nanoindentation tests, from which the coating Young’smoduli were determined using equations (1) and (2) asan average from 10 measurements. These values aresummarized in Table 1 together with those from theother independent test methods. For some indenta-tions, in particular for the thinner coatings, viscouseffects due to the polymer substrate were apparent.This resulted occasionally in nearly vertical unloadingcurves in the force–indentation diagrams.

SIEBIMM

The buckling formation was observed in situ by SEM.The strain for onset of buckling could then bedetermined. It was noticed that the strains for initialbuckling showed a significant variation, since theelastic instability was controlled by the local thicknessand elastic properties. A high tensile strain at bucklingonset was observed for the thinner TiO2 and MOXcoatings of 0.31 and 0.42, respectively. For 20-nmcoatings, the buckling was perceptible at tensile strainof 0.1. Occasionally, cracks were found along the ridgesof the larger buckles (see Fig. 6). The wavelength ofthe cracked buckles was, however, found to be thesame as with the uncracked ones. If a band of adjacentbuckles could not be found, the local wavelength oflocalized buckles was measured.

Figure 6 shows the buckles on top of a TiO2 coatingof 6 nm strained at 0.31 in the longitudinal direction.This image was taken from the center of the sample,where the load was relatively uniform. This region had

0

1

2

3

4

0.5

1.5

2.5

3.5

TiO2 6nm TiO2 20nm MOX 4 nm MOX 20 nm

Youn

g’s

mod

ulus

of c

oate

d po

lym

er (G

Pa)

Fig. 4: Young’s modulus of the coated substratedetermined from tensile testing

00

5

25

20

20 4030

15

10

100

0

5

25

20

20 4030

15

10

100

0

5

25

20

20 4030

15

10

100

0

5

25

20

20 4030

15

10

10Indentation depth (nm) Indentation depth (nm) Indentation depth (nm) Indentation depth (nm)

Forc

e (μ

N)

(a) (b) (c) (d)

Fig. 5: Force versus indentation depth in nanoindentation of the coatings. (a) TiO2 6 nm, (b) TiO2 20 nm, (c) MOX 4 nm,(d) MOX 20 nm

Table 1: Summary of estimated Young’s moduli for the characterization methods

Barrier coating hc (nm) Tensile test Nanoindentation AFM SIEBIMM Bulk19

Es (GPa) Ec (GPa) Ec (GPa) Ec (GPa) Ec (GPa)

TiO2 6 3.5 ± 0.4 44 ± 19 45 ± 2 420 ± 140 115TiO2 20 3.2 ± 0.4 35 ± 17 76 ± 2 64 ± 16 115MOX 4 3.7 ± 0.2 31 ± 21 29 ± 2 240 ± 110 160–180MOX 20 3.0 ± 0.3 23 ± 09 43 ± 3 52 ± 13 160–180

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a high and relatively uniform concentration of buckles.The wavelength pattern was different that thosereported by Stafford et al.,15 Volynskii et al.14 andCranston et al.16 The buckling process is influenced bythe stiffness ratio Ec/Es, according to equation (3). Thereported values ranged in the order 103–105, whereas inthis study, the ratio is in the range of 7–25.

All coatings had similar buckling shapes, but showeddifferent dimensions. The size of all buckles wasmeasured in SEM, and then averaged to calculate thecoating Young’s modulus Ec using equation (4).Thevalues are reported in Table 1.

AFM peak-force measurement

The indentation depth in AFM was controlled within afew nm by a sharp AFM tip (radius of 8 nm), resultingin a limited influence from the substrate, which issuitable for ultrathin coating measurement. Figure 7shows the effective Young’s modulus maps of an areaapproximately 1 lm2 for each barrier coating. Eachpixel in the AFM images corresponds to a loopdescribed previously in Fig. 3. Each image is madeup of more than 250,000 measurements. The measuredforce deformations for all pixels were used to deter-mine the average Young’s modulus of the coating bymeans of equations (2) and (5).

The average values of each coating are reported inTable 1.

Comparison of results

The results from the various characterization tech-niques are summarized in Table 1, where the scatter isthe standard deviation of these values. Although thevalues of Young’s moduli of the coatings are all in thesame order of magnitude, there are considerabledifferences. The TiO2 coating of 6 nm and MOXcoating of 4 nm had a relatively good agreementbetween the nanoindentation and AFM values, butvery high values for the SIEBIMM values. The 20-nmcoatings showed a large variation between the differentcharacterization methods. For comparison, bulk mod-ulus from literature13 is presented in the last column.The presented methods showed higher stiffness for theTiO2 coatings than for the MOX coatings, which wasnot the case for the literature values.13 The bucklingmethod, SIEBIMM, showed higher Young’s moduli forboth materials, compared with AFM and nanoinden-tation. The sensitivity and accuracy of this method isadversely affected by the relatively low coating–substrate stiffness ratio and inelastic deformations athigh strains. As for the thickness dependency, only theAFM test method showed a higher stiffness for thethicker coatings.

Concluding remarks

Four test methods were experimentally assessed for thedetermination of the Young’s moduli of ultrathincoatings deposited on polymer films. Tensile testing isgenerally unsuitable to characterize the coating stiff-ness for thin coated polymer films, since the measuredstiffness is dominated by the relatively thick polymerfilm.

Nanoindentation gave crude measures of the coatingYoung’s modulus, since the radius of the indenter tipwas higher than the coating thickness. Viscous effectswere observed, in particular for the thinly coatedsamples. The viscoelastic polymer substrate thenaffects the estimated coating stiffness.21 Shallow inden-tations with sharp tips on nonviscous materials wouldlead to improvements in this method. We tried to useapproaches that account the substrate effect: Song/

Fig. 6: SEM image of a TiO2 coating of 6 nm stretched at0.31 strain

Fig. 7: AFM mapping of the effective Young’s modulus of the coatings. (a) TiO2 6 nm, (b) TiO2 20 nm, (c) MOX 4 nm,(d) MOX 20 nm

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Pharr,22 Hay/Crawford23 and King21 models. Thesemodels were developed in a range of hundred nanome-ters of coating thickness and depend of normalizedcontact area over coating thickness. It turned out thatthe values of the coating modulus did not improve.

The buckling method, SIEBIMM, has the advantageof being straightforward, not requiring an advancedapparatus to measure local loads and displacements. Inour case, the buckles appeared at relatively highstrains, where the polymer substrate was subject tosome inelastic deformation. The moduli estimatedusing this method were much larger than those fromthe other methods. With a fully elastic and relativelycompliant substrate, such as an elastomer, the clarity ofthe buckles and the calculated stiffness values wouldmost likely be improved.

The AFM method showed the most consistentresults compared with the other methods. In this case,the tip radius was much smaller than that used innanoindentation, which provides better local deforma-tion. For nanoindentation, it is likely pushing thecoating down rather than being measured from a pointload at a pyramidal tip (Oliver Pharr method). Addi-tionally, the indentation depth in AFM is small enoughto have only a limited influence of the viscoelasticpolymer substrate. A considerable amount of stiffnessdata is generated, since the AFM equipment automat-ically maps a given area. This leads to average values ofhigher statistical confidence compared with few point-wise measurements in nanoindentation.

Although the variation was large in absolute num-bers, all methods showed the same trend with a higherYoung’s modulus for the TiO2 coating than for theMOX coating (TiO2 and Al2O3 mixture). This isconsistent with Cherneva et al.,24 who measured aslightly higher modulus for TiO2 than for Al2O3.

The influence of the coating thickness on Young’smodulus of the coating was not conclusive. The AFMmethod showed a higher stiffness for the thickercoating, whereas the other methods showed the reverseorder, which was also found by Chen et al.7

To estimate Young’s modulus of coatings with athickness of less than 0.1 lm is indeed challenging. Thepresent paper presents a number of candidate meth-ods, each with their specific advantages and drawbacks.Both AFM and nanoindentation are based on pushingin and retracting a rigid tip. AFM allows for moreshallow indentation, which makes it more suitable forvery thin coatings. The advantage of SIEBIMM is itssimplicity since only the wavelength of coating bucklesneeds to be measured, but it relies on elastic reversibledeformation at relatively high strains which is notalways feasible. Overall, all presented methods stillshow potential to rank the stiffness properties of thincoatings. Quantitative measures are, however, needed,e.g., in models predicting the fragmentation of strainedbrittle coatings.

Acknowledgments Financial support and supply ofthe test samples for this study by Tetra Pak PackagingSolutions AB, Development and Service Operations(DSO), are much appreciated. MVTdC is grateful tothe European Commission within the framework of theErasmus Mundus Programme, Action 2-STRAND 1,Lot 9, Brazil for PhD studies. The authors would liketo thank Mr. Luimar Correa for the technicalassistance in the SEM experiments and Dr. JessicaBolinsson for valuable scientific discussions.

Open Access This article is distributed under theterms of the Creative Commons Attribution 4.0International License (http://creativecommons.org/licenses/by/4.0/), which permits unrestricted use, distribu-tion, and reproduction in any medium, provided yougive appropriate credit to the original author(s) andthe source, provide a link to the Creative Commonslicense, and indicate if changes were made.

Appendix A: Effects of the conductive layerin the barrier coating properties

Observation of buckling caused by elastic instability inthin coatings typically requires a conductive coating inSEM analysis since the features of interest are in thesub-micron scale. In order to reduce charging effects,all the samples were coated on top of the barrier layerwith a Au–Pd conductive coating by Quorum SC7640polaron sputter coater. Au–Pd is one of thinnestcontinuous coating materials for SEM purposes.25 Forthe deposition process the following parameters wereused: an accelerate voltage of 2 kV, a plasma dischargecurrent of 20 mA, a pressure between 5 and 8 Pa and adeposition time of 1 min. Using these parametersprovides a layer of 8.5 nm thick,26 which was theminimum thickness allowing observation of surfacebuckling and fracture without excessive chargingeffects. The issue to address here is whether or notthis additional layer would have any influence on thebuckling formation and size.

To investigate the influence of the Au–Pd coating,two different test samples were produced with the TiO2

20-nm coating. In the first one, the sample wasstretched up to a strain level of 0.12, where the buckleswere clearly observed and then Au–Pd was sputteredon top of coating. In the second one, the Au–Pd wassputtered before the stretching at the same strain. Bothtest samples were observed in the SEM using anacceleration voltage of 3 kV and the high-efficiencysecondary electron HE-SE2 detector. The micrographscan be seen in Fig. 8, where (a) the precoated and (b)the post-coated samples show the same deformation

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behavior. The similarities are also found at highermagnification for (c) the precoated and (d) post-coatedsamples. Not only qualitatively, but also quantitativelydid the dimensions of the buckles not show anysignificant difference if the conductive coating wasapplied before or after the buckling formation. It istherefore assumed that the Au–Pd coating did notinfluence the mechanical behavior in the SIEBIMMtest of the TiO2 coated film.

In addition, the conductive coating did not havesignificant influence on the fragmentation process ofthe brittle barrier coating in tensile loading. Figures 9and 10 show that the crack accumulation was notinfluenced by the thickness of the conductive coatingfor the MOX and TiO2 films, respectively.

The Au–Pd is not a continuous monolithic material,but shows a granular structure at high magnification as

shown in Fig. 11, obtained in 246 k9 of magnificationwith an in-lens secondary electron detector. The Au–Pd layer has a granular structure on the polymersurface in the crack, as well as on the TiO2 barrier. Ifthese granules are separated from each other intension, the conductive layer cannot expect to carrymuch stress, which has also been pointed out byRochat et al.27

Based on the negligible influence of the Au–Pdcoating on the buckling formation and tensile crack-ing of the investigated metal oxide coatings, it isassumed that the conductive coating does not haveany influence in the mechanical analysis of theSIEBIMM test. The observed microstructure withseparable granules in the conductive coating can be areason for the very limited mechanical effect of thematerial.

Fig. 8: Stretched samples coated with 20 nm TiO2: (a) and (c) have been deposited with a conductive Au–Pd coating beforestretching, and (b) and (d) after stretching

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1.5 nm 8.5 nm4 nm

ε = 0.05

ε = 0.10

ε = 0.20

Fig. 9: SEM images showing a comparable accumulation of crack with increasing tensile strain for various thicknesses ofconductive coatings on for 20 nm MOX films

1.5 nm 8.5 nm4 nm

ε = 0.05

ε = 0.10

ε = 0.20

Fig. 10: SEM images showing a comparable accumulation of crack with increasing tensile strain for various thicknesses ofconductive coatings on for 20 nm TiO2 films

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Fig. 11: High-resolution SEM image of a crack of TiO2

20 nm showing the granular nanoscale structure of the Au–Pd coating both on the polymer substrate in the crack andon top of TiO2 layer

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