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Adhesion and degradation of hard coatings on poly (methyl methacrylate) substrates
Ani Kamer a, Kjersta Larson-Smith b, Liam S.C. Pingree b, Reinhold H. Dauskardt a,⁎a Materials Science and Engineering Dept., Stanford University, CA, USAb Boeing Research & Technology, Chemical Technologies Division, WA, USA
The adhesion and time dependent crack growth behavior of polysiloxane based hard coatings on poly(methyl methacrylate) substrates were investigated. The adhesive fracture energies for different coatingswere quantitatively characterized and varied between 1.4 J/m2 and 22 J/m2. Significant time dependent crackgrowth in various moist environments was observed and was consistent with a viscoelastic crack growthmodel. The effect of selected weathering treatments was also examined and resulted in a significant drop incoating adhesion. The coatings were analyzed using surface sensitive techniques; structural changes in thecoatings resulting from various exposure doses were studied and mechanisms responsible for the observeddegradation in adhesion were discussed.
Organic polymers are an attractive alternative to silica basedglasses in optical and structural applications such as eyeglass lensesand airplane passenger windows due to their low density and highresistance to fracture. Poly (methyl methacrylate) (PMMA) is athermoplastic widely used in these applications owing to additionaldesirable properties such as visible light transparency and low glasstransition temperature. However, PMMA's soft surface is prone toscratching and absorption of small molecules such as oxygen andwater vapor. In most applications PMMA needs a protective coatingthat is hard, water repellent and that adheres well to the organicpolymer surface. Hybrid organic/inorganic silicon-based coatings arean excellent candidate for such an application because the inorganicsilica part provides hardness and the organic part accounts forflexibility and hydrophobicity.
While many studies have investigated the influence of compositionand curing conditions on the hardness [1] and barrier properties [2] ofhybrid sol–gel coatings, less research has focused on coating–substrateadhesion. Attempts to measure adhesion have been made by cross-cuttape tests [3,4], andoptical observation of debonding [5]. Thesemethodsprovide results that are at best qualitative and add little to theunderstanding of interface damage mechanisms. More quantitativemethods have been employed by loading a coated substrate in uniaxialtension andmeasuring the crack density observed at the coating surface[6]. In this approach, the fracture energywas calculated using the classicperfectly plastic stress transfer theory. While this method is more
reliable compared to tape tests, it requires the knowledge ofmechanicalproperties of the film and the intricate nature of plastic deformation ofthe substrate surface.
In this work, we employ a modified double cantilever beam (DCB)method [7,8] for the quantitative assessment of hard coating–polymersubstrate adhesion. When a weak interface is present between acompliant film and an elastically stiff substrate, DCB specimen testsresult in crack growth at the weak interface. However, in the case of acompliant substrate and a stiff coating the crack does not necessarilystabilize at the interface and can deflect into the substrate. Anasymmetric DCB specimen was developed to reverse this intrinsictendency of the crack to deflect from the weak interface into thesubstrate. The asymmetric DCB specimen yielded reproducible fractureenergy results with small scatter andwas shown to be a convenient andreliablemethod for characterizing theadhesionof hard coating–polymerinterfaces. This method was used to investigate the role of interfacechemistry anddegradationon thedurability of a coating operatingunderoutdoor environmental conditions. The subcritical debonding undervarious environmental conditions was studied. The hard coatings wereexposed to prolonged simulated solar radiation andwater sprayingwithtemperature cycling and the effect of these treatments on adhesion wasinvestigated. Lastly, possible mechanisms for the degradation ofadhesion under simulated outdoor conditions are discussed.
2. Experimental procedures
2.1. Materials and coatings
Military grade (MIL P 25690) PMMA plates with a thickness of5.53 mm or 12.8 mm were used. Hard coatings were prepared using
1908 A. Kamer et al. / Thin Solid Films 519 (2011) 1907–1913
proprietary sol–gel formulations starting from alkoxysilane basedreactants. Tetraethoxysilane and methyltriethoxysilane were firsthydrolyzed in an alcoholic solvent and were subsequently con-densated to form a coating precursor. The precursor was applied onthe PMMA sheets by dip-coating. The solvent in the films wasallowed to evaporate and the dried coatings were thermally cured inorder to realize their intended hardness. Commercially availablecoatings labeled A to E, made with different precursors, were appliedon planar PMMA sheets. The coating labeled as “window” wasapplied on passenger aircraft hardware (a PMMA sheet curved andmachined to fit in an airplane hull); the used window was in servicefor 5 years. All coatings were 4 μm thick, except for coating C, whichwas 10 μm thick.
2.2. Weathering treatment
The effects of operating conditions on hard coated airplanepassenger windows were simulated using weathering treatments inaccordance with the SAE J1960 automotive exterior standard.Outdoor exposure was simulated by a Xenon Arc weatheringchamber [Atlas Ci4000 Xenon WeatherOmeter, Atlas MaterialTesting Technology, Chicago, IL] with Quartz and Type S Borosilicatefilters. The simulation consisted of 3 hour cycles of 40 min light,89 °C, 50% relative humidity (RH); 20 min light, front spray, 89 °C,50% RH; 60 min light, 38 °C, 95% RH; 60 min dark, front spray, 38 °C,95% RH. The samples were treated until they reached the desireddose in units of energy per area at 340 nm wavelength. Tests werealso performed on new and used airplane windows, and onweathered samples.
2.3. Adhesion tests
Symmetric (h1=h2) and asymmetric (h1bh2) DCB specimenswere prepared by gluing a blank (uncoated) substrate onto a coatedsubstrate, (Fig. 1). The thickness of the glue layer was ~2 μm. Whenspecimens were prepared from aircraft hardware samples, the beamswere cut in the direction of the smallest curvature. The curvature wassmall enough to be neglected in beam-bending compliance calcula-tions. Prior to bonding, the coated surface was cleaned withisopropanol and was treated with UV light and ozone in a UV-Ozone cleaner [Jelight 144AX, Jelight Company Inc., Irvine, CA] for3 min to improve the adhesion of the hydrophobic hard coating to theepoxy glue. Aluminum loading tabs were attached on both beams. Thespecimens were tested at least one week after preparation to ensure
proper curing of the glue at room temperature. Curing at highertemperatures for shorter periods resulted in cracking of the coating.The specimens had in plane dimensions of 7 mm×80 mm, and thetotal thickness was 10 mm (symmetric specimens) or 8.5 mm(asymmetric specimens). The fracture test was started by first makinga 1–2 mm cut with a slow-turning diamond saw near the interfacebetween the PMMA substrate and the coating, and then inserting arazor blade into the cut and applying a gentle opening movement topropagate the crack a few millimeters (Fig. 1). The crack length waseither registered by visual inspection or was calculated using anextension of the compliance of a beam on an elastic foundation due toKanninen [9]:
C =4E′B
ah1
+ 0:64� �3
+ah2
+ 0:64� �3� �
ð1Þ
where C is the compliance of the double cantilever beam specimen,defined as the slope of displacement versus load, E′ is the plane strainelastic modulus, a is the crack length, h1 and h2 are the beamthicknesses and B is the beam width. Note that since the thickness ofthe coating was less than one thousandth of the thickness of thesubstrates the sandwich specimen was analyzed like a homogenousspecimen [10].
The adhesion tests were conducted on a micromechanical testsystem [DTS Mechanical Delaminator Test System, DTS Company,Menlo Park, CA], in displacement control mode. The beam ends weredisplaced with a known displacement rate, and the load wasmeasured simultaneously. Both the displacement and the load weremeasured in far-field and the compliance of the loading system wasaccounted for. The pre-cracked specimen was loaded in tension untilcrack growth occurred, letting the crack grow and then unloading it.Data of crack growth over tens of millimeters was obtained from eachspecimen. The strain energy release rate or driving force, G, wascalculated using Eq. (2) [11]:
G =6P2
E′B2
ah1
+ 0:64� �2 1
h1+
ah2
+ 0:64� �2 1
h2
� �ð2Þ
where P is the load applied to each arm of the DCB. The fractureenergy in units of energy per area, Gc, was calculated using the load Pcwhere the load–displacement curve started to deviate from linearity,Fig. 2.This analysis was used for both symmetric and asymmetric DCBspecimens by inserting the appropriate beam thickness.
2.4. Phase angle and crack path
There are two parameters used to define the loading of a crack: theglobal driving force for crack growth, G, and the stress intensity factor,
Fig. 2. Load–displacement data obtained from an asymmetric DCB specimen. Pc is thecritical loadwhere the load–displacement curve deviates from linearity; 1/C is the slopeof the linear load–displacement region, defined as the inverse compliance.
1909A. Kamer et al. / Thin Solid Films 519 (2011) 1907–1913
K, defining the stress state near a crack tip. For linear elastic materialsthese two parameters are uniquely related as:
G =K2I
E′+
K2II
E′ð3Þ
where KI is the stress intensity factor resulting from opening (mode I)loading and KII results from in-plane-shear (mode II or sliding). E′ isYoung's modulus, in plane-stress or plane-strain, depending on thethickness of the specimen.
The phase angle Ψ is defined in Eq. (4), where KI and KII arerespectively the opening and sliding stress intensity factors resultingfrom a particular loading condition applied remotely or locally at thecrack tip:
Ψ = tan−1 KII
KIð4Þ
In the reference system shown in Fig. 1 for a crack at the interfacebetween the PMMA substrate (bottom) and the hard coating (top), apositive phase angle drives the crack towards the substrate and anegative phase angle drives the crack towards the hard coating. In asymmetric DCB specimen, applying equal loads P on each arm resultsin pure mode I loading at the crack tip (Ψ=0) and the crack stays atthe interface. In an asymmetric DCB specimen, where the beams havesimilar elastic properties but their thicknesses are different, the crackis driven towards the thinner beamwhen equal loads P are applied oneach arm [12] (Fig. 3). This implies that when the top beam is thinnerthan the bottom beam the phase angle is negative, and when it isthicker the phase angle is positive.
Remote loading is only one of the factors that affect the crack pathat an interface. Elastic properties, fracture toughness of the substratesand the layers adjacent to the fracture interface and plasticdeformation characteristics of these determine whether a crack willstabilize at the interface of interest. In the case of the PMMA–hardcoating system, the elastic mismatch of the substrate and the thinlayer can be characterized by a Dundurs' parameter α~−0.12, whichis a small value. The effect of α on the phase angle is less than 1° at thisinterface, as presented in Fig. 8 in [13]. It can be concluded that theeffect of the elastic mismatch on the phase angle is negligible.
In the present work, asymmetric loading was applied to opposethe intrinsic tendency of the crack to deflect into the bottomsubstrate. This was achieved by choosing a combination (modemixity) of KI and KII stress intensity factors that would force the crackin the opposite direction, i.e. towards the hard coating. The desiredmode mixity was obtained by reducing the thickness of the topsubstrate. The top to bottom beam thickness ratio of 0.7 resulted in a
y hh
30o
20o
10o
-10o
-20o
-30o
-40o
0o
-1.0 -0.8 -0.6 -0.4 -0.2 0.2 0.4 0.6 0.8 1.0
y/h0
Ψ =
tan-1
(KII/
KI)
40o
Fig. 3. The mode mixity Ψ=tan−1(KII/KI), plotted against the offset y/h for a doublecantilever beam with a crack off the mid-plane, shown in the inset. Figure adaptedfrom [12].
mode mixity phase angle Ψ=−17°. This small mode mixity wasenough to stabilize the crack at the interface of interest and thefracture energy values of several different hard coatings on PMMAsubstrates were measured.
2.5. Subcritical crack growth tests
Subcritical (i.e. propagating with a driving force below Gc) crackgrowth can be observed in materials that undergo some time-dependent process at the crack tip, such as a reaction, diffusion orviscoelastic relaxation. A v–G curve (plot of crack growth rate v in m/sversus crack driving force G in J/m2) can exhibit several differentregions depending on the nature of the time-dependent mechanismat the crack tip. At driving forces close to Gc the crack growth rate isvery fast and independent of the environment. At intermediate high G,the crack growth rate is limited by the diffusion of the reactive speciesto the crack tip and is independent of the driving force; an example ofsuch behavior can be seen in strained Si–O–Si bonds reacting withwater vapor at the crack tip of silica-based glass [14]. Just above thethreshold driving force for crack growth, Gth, the crack growth ratedepends on G and the activity of the environmental species takingpart in the reaction. The mechanism can be simply modeled as achemical reaction with activation energy barrier Ea which is loweredas bonds at the crack tip are strained by G and are more susceptible toreact with the environmental species.
Time dependent crack growth in different environments wasquantified by load relaxation tests conducted in a humidity andtemperature controlled environmental chamber [Associated Environ-mental Systems, Ayer, MA]. These tests were performed on samplesmachined from flight hardware at constant temperature of 23 °C andat several relative humidity conditions. The system was allowed toequilibrate for 4–6 h after the specimen was inserted. In the case ofliquid environment, only the specimen was immersed in the solutionof interest. The pre-cracked specimen was then loaded until crackgrowth initiated and at that point displacement was fixed. Crackgrowth continued for several days with monotonically decreasingapplied G. As the crack grew, the load decreased and the complianceincreased. Crack growth rate was calculated from the change incompliance with time, by plotting a(t) calculated using Eq. (1) andtaking the numerical derivative with respect to time. The applied G ateach t was calculated from load and crack length data. Whenapplicable, calculated crack length was calibrated by optical mea-surement of the crack length at the end of the test.
2.6. Channel cracking
Time dependent crack growth in the coatings was also studied bychannel cracking [15]. A beam machined from coated aircrafthardware was loaded in 4-point-bending in a mechanical testmachine [Mini Bionix, Materials Testing Systems Corporation, EdenPrairie, MN]. Prior to loading, scratches perpendicular to the bendingstress were introduced on the surface of the coating by a microindenter. The scratcheswere 100 μm long and 100 μmapart from eachother. When the beam was loaded cracks started to grow from eachend of the scratches, perpendicular to the bending stress direction.Crack growth rate was recorded by a long distance microscope[Questar Step Zoom M100, Questar Corporation, New Hope, PA] as afunction of crack driving force G, which was calculated using Eq. (5)from [15]:
G =Zσ2
f hfEf
ð5Þ
Where Z is a dimensionless parameter close to 4 for channelcracking with substrate damage [12], hf is the thickness of the film, Ēf
Fracture surface of coating side
Fracture surface of PMMA side
crack
100 µm
100 µm
Fig. 4. Fracture surfaces of a symmetric DCB specimen, Coat A, Gc~8 J/m2.
Fig. 5. Fracture energy of various coatings on PMMA substrates. Coat A to E are coatingson PMMA plates, “New window” is a pristine hard coated PMMA airplane window,“Used window” was in service for 5 years.
1910 A. Kamer et al. / Thin Solid Films 519 (2011) 1907–1913
is the biaxial modulus of the film and σf is the stress in the film. Theappropriate Z value was chosen from [12] taking into account theelastic mismatch between the substrate and the coating. Z alsodepends on the exact depth of the channel crack in the substrate;since this value was not experimentally measured a reasonableassumption of ~1 μmwasmade. The films studied in channel crackingwere 4 μm thick and were assumed to have zero residual stress beforeloading; Ēf was 5 GPa (V.S. Sundaram, personal communication). Thefilm stress was calculated using Eq. (6) from [15]:
σf =3PLh2SB
E′fE′s
!ð6Þ
where P is the load, L is the distance between the inner and outerloading pins, Es′ is the substrate plane strain modulus 3.94 GPa, hs isthe substrate thickness and B is the beam width.
2.7. Surface characterization
Fourier Transform Infrared Spectroscopy (FTIR) [Bruker OpticsVertex 70, Bruker Optics Inc., Billerica, MA] with an Attenuated TotalReflectance (ATR) accessory [MIRacle™ Single Reflection, PikeTechnologies, Madison, WI] and X-Ray Photoelectron Spectroscopy(XPS) [Monochromatized X-ray Photoelectron Spectrometer with AlKα radiation, Surface Science Instruments, Mountain View, CA]analyses were used to detect whether failure was interfacial. Optical[Polyvar MET optical microscope, Reichert-Jung, Wien, Austria] andscanning electron microscopy [FEI XL30 Sirion SEM, FEI Company,Hillsboro, OR] observations were used to characterize fracturesurface topography.
3. Results and discussion
3.1. Fracture energy
3.1.1. Symmetric double cantilever beam specimensThe fracture energy of Coat A measured using a symmetric DCB
specimen was around 8 J/m2 and the fracture surfaces resulting fromthis test are shown in Fig. 4. The crack deflected from the interface intothe PMMA substrate and then back to the interface generating a wavyfracture surface. Load–displacement curves exhibited sharp drops inload indicating unstable “stick–slip” cracking behavior. Stable inter-facial crack growth could not be obtained with a symmetric DCBconfiguration although the DCB test is an established technique forthe assessment of the fracture energy of thin films [16]. Thisobservationwas surprising because it is well known that an interfacialcrack will stabilize at a weak interface when loaded in pure openingmode. Note that the majority of thin film adhesion results reported inthe literature involve compliant films (such as low-k layers andadhesive bonds) on stiff crystalline substrates (such as single crystalsilicon and metal).The stiff substrate will constrain the crack either atthe weak interface or in the coating. In contrast, the material systeminvestigated in this work consisted of a compliant amorphoussubstrate and a much harder coating. The soft surface of the PMMAsubstrate could not prevent the crack from deflecting out of theinterface. Thus, we observed a tendency of the crack to propagate intothe substrate under the coating.
Similar behavior concerning interfaces between incompatible bulkpolymers has been reported in a study investigating PMMA andpolystyrene joints [17]. The interfacial fracture between theseimmiscible polymer blocks yielded unexpectedly high fracturetoughness values when loaded in pure mode I. These values did notreflect the fracture toughness of the interface because the crackdeflected into the polymer with lower crazing resistance. Crazing ledto the overestimation of the interfacial fracture toughness by almost
an order of magnitude. It was shown that an appropriate asymmetricloading confined the crack at the interface and eliminated the effect ofcrazing on the interfacial fracture toughness.
3.1.2. Asymmetric double cantilever beam specimensThe adhesive fracture energy values measured for five different
coatings on stretched PMMA plates, a pristine coating on a PMMA
1911A. Kamer et al. / Thin Solid Films 519 (2011) 1907–1913
window and a window that was in service for 5 years are shown inFig. 5. These values were measured using an asymmetric DCBspecimen with a remote loading mode mixity of −17°, as describedin the experimental section. The tests resulted in clean interfacialdelamination of all coatings, which was confirmed by opticalmicroscopy, XPS and FTIR analyses. The fracture energy values ofdifferent coatings varied between 1.3 and 22 J/m2, with very smallscatter and high reproducibility. It should be noted that the standardcross-cut tape test, which all of these coatings had passed, wasinsufficient to distinguish between the adhesion quality of coatingswith an order of magnitude difference in fracture energy. Whencompared to the fracture toughness values for polysiloxane basedlow-k films on single crystal Si reported in the literature [18–20], thevalues measured in our study are significantly higher. This is anexpected result, since the hard coatings investigated in the presentwork contained a higher organic fraction, and were on plasticallydeformable substrates. This resulted in increased energy dissipationthrough material pullout and plastic deformation in the coating andin the substrate.
3.2. Subcritical crack growth
Subcritical crack growth rate plotted as a function of applied G for anew window with a hard coating is presented in Fig. 6. In theinvestigated G range a threshold for crack growth between the PMMAand the hard coating was not observed. At high driving forces (higherthan 3 J/m2) the subcritical crack growth behavior was not influencedby the change in relative humidity (RH) from 95% to 15%. However, atlower driving forces, delamination at 15% RH happened faster than at45% and 95%. This behavior was surprising since previous studies onsilica-like films indicated that subcritical debonding at interfacescontaining Si–O–Si bonds was faster at higher humidity [21]. Apossible explanation for this behavior could be the hydration of thePMMA substrate; PMMA can absorb more than 2 wt% of water in ahumid environment [22,23]. At room temperature and average RH of~50% (the storage condition for our samples), PMMA reachesequilibrium at ~0.5 wt.% water content [23]. Water absorption anddesorption happening in the substrate at the crack tip may thereforemask the effect of environmental humidity on crack growth rates atlow driving forces.
When the driving force was higher than 3 J/m2 the crack growthrate was not affected by the relative humidity and the slope of the v–Gcurve remained close to 8. This behavior can be explained using a K-controlled fracture analysis for linear viscoelastic materials [11]. Thefundamental assumption in this analysis is that the bulk of thespecimen is linearly elastic while only a small region near the crack tip
Fig. 6. Crack growth rate as a function of applied strain energy release rate at 23 °C. Fifteenpercent relative humidity (15% RH) is marked with open and filled circles; 45% RH withopen andfilled triangles, 95% RHwith open,filled anddotted squares.More than one curveat each humidity are presented to demonstrate the reproducibility of the method.
deforms viscoelastically. The crack growth rate da/dt is expected todepend on G as:
dadt
∝G12n ð7Þ
where n is the exponent of the polymer relaxation modulus in Eq. (8)
E = Eot−n ð8Þ
In the case of a crack between a hard polysiloxane coating andPMMA, the time dependent behavior at the crack tip will bedominated by the viscoelastic deformation in the polymer. Indeed,the subcritical crack growth that was observed in this system is inexcellent agreement with that reported for bulk PMMA. It was foundthat subcritical crack growth rate in bulk PMMA depends on mode Istress intensity factor KI with an exponent of 1/0.06 [24].
Using the relationship between G and K in Eq. (3) it can be shownthat this is equivalent to:
dadt
∼G8:33 ð9Þ
This behavior was closely reproduced by our experiments forseveral orders of magnitude of crack growth rate. The modeldeveloped in [24] assumed a crack tip opening displacement of1.8 μm for bulk PMMA. In this work, the crack tip openingdisplacement was assumed to be close to half of the bulk (0.7 μm),to account for the lack of viscoelastic deformation in the hard coatingon one side of the crack, Eq. (10), where Eo is ~3 GPa, u is the crack tipopening displacement 0.7 μm, εy is the yield strain 0.06 [24]. Themodel with this modification of the viscoelastic zone size agreed verywell with the experimental data, shown in Fig. 7.
G = Eou�yπ�yu
� �2n dadt
� �2nð10Þ
At low humidity conditions, diffusion of water out of the PMMAwould de-plasticize the polymer (i.e. increase its glass transitiontemperature) and n would decrease [25] (note that n is related to theloss factor δ as tanδ= tan(πn/2) [26]). In this case, when the crackgrowth rate was slow enough for this diffusion process to reachequilibrium an increase in the v–G curve slope would be expected.However, a decrease in the slope was observed, which indicates thatunder conditions of low humidity (below 15%RH) and low strainenergy release rate (below 3 J/m2) the mechanism of crack growthswitches to a faster mechanism independent of viscoelastic deforma-tion at the crack tip. Further investigation at low humidity is needed
Fig. 7. Viscoelastic model of subcritical crack growth described by Eq. (10) (solid line)and experimental data obtained at 23 °C and 45% RH (squares).
Fig. 9. Change in fracture energy with increasing simulated outdoor treatment.
1912 A. Kamer et al. / Thin Solid Films 519 (2011) 1907–1913
to understand the precise mechanism of crack growth at low strainenergy release rates in this material system.
3.3. Channel cracking
The v–G curve obtained from channel cracking was in goodagreement with data from asymmetric DCB tests under similarconditions (Fig. 8). This agreement is most probably due to the channelcracks penetrating into the substrate and employing the previouslydiscussed subcritical crack growth mechanism with a viscoelastic zone.A channel crackingmechanismwith lagging viscoelastic deformation inthe substrate could explain this behavior: the channel propagatedrapidly in the coating (cohesive channel crack)while propagation in thesubstrate lagged behind the cohesive crack. Extension of the crack couldbe detected only after it penetrated into the substrate and created alarger opening of the crack in the film (this phenomenon is illustrated inFig. 49 in [12]). The observed channel crack growth rate and thesubcritical crack growth rate at the interface of the substrate and thehard coating are in good agreement because both are limited by theviscoelastic deformation in the substrate close to the interface.
Substrate damage in channel-cracking specimens could not berecorded optically due to insufficient contrast between the film andthe substrate. Nevertheless, weathered windows showed very similarchannel cracking patterns and substrate damage was observed underthe coating (after delamination of the coating).
Calculating the driving force for channel cracking requires theknowledge of the elastic properties of the film, the film thickness andthe residual stress present in the film resulting from processing priorto loading. Because the mentioned parameters were known for thecoatings in this study it was possible to compare channel crackingwith the asymmetric DCB method. A drawback of the channelcracking method is that it is very difficult to obtain sufficient contrastin transparent substrate–coating systems for crack growth ratedetermination. The asymmetric DCB is a much more versatile methodfor studying subcritical crack growth in transparent coatings withunknown elastic properties.
3.4. Effects of weathering
Fracture energy results using asymmetric DCB specimen for windowsamples exposed to simulated outdoor treatment can be seen in Fig. 9.The samples were treated until they reached doses of 625, 1250 or2500 kJ/m2 (a 2500 kJ/m2 dose is equivalent to 1 year solar exposure inFlorida). Even the lowest dose caused a significant drop in fractureenergy.While the fracture energymaintained this lower value for dosesup to 2500 kJ/m2, it subsequently increased dramaticallywith a 5000 kJ/m2 dose to 16.7 J/m2. The surfaces of pristine samples and samples
Fig. 8. Crack growth rate as a function of applied strain energy release rate: opensquares were measured by asymmetric DCB, filled squares were measured by channelcracking.
exposed to lowdoses appeared smooth and featurelessunder theopticalmicroscope. However, the surface of a sample exposed to the highestdose was cracked, with damage extending deeper into the PMMAsubstrate, shown in Fig. 10. The cracks cut the hard coating intomillimeter-sized islands still attached to the substrate and the lighterregions around the cracks are evidence of edge delamination.
Fig. 11 shows the ATR-FTIR spectra for a hard coating exposed todifferent weathering doses. The curves have been normalized withrespect to the Si–CH3 peak at 760 cm−1 [27] for comparison of the Si–O–Si peaks at ~1000 cm−1 [28]. It can be seen that the intensity of theSi–O–Si peak increased with increasing exposure. The same effect ofUV light on siloxane-based films has been reported in the literature[29]. The ATR-FTIR results show that network Si–O–Si bonds formedand organic end groups were lost as a result of prolonged solar-likeexposure. The loss of organic groupsmight be diminishing the amountof favorable interactions between the coating and the organicsubstrate. This could account for a small drop in fracture energy withsimulated outdoor treatment. Furthermore, cross-linking leads to thedensification of the hard coating and must be creating tensile stressesin the film. However, the asymmetric DCB method that we employedfor adhesionmeasurement eliminates the effect of the film stresses onthe crack driving force by maintaining the constraint on the coatingbefore and after the measurement. This is why the shrinkage of thecoating by itself is unlikely to explain the decrease in fracture energy.
UV exposure has a cross-linking effect on the polysiloxane basedhard coating and can induce chain scission in the PMMA substrate[30]. A UV photon can cleave a PMMA chain, forming two radical ends,which can propagate this defect by reacting with nearby polymerchains. In the presence of oxygen radical ends can stabilize and defectscan grow to form micron-size cracks. The number and size of thesemicrocracks would be limited by diffusion of oxygen from the surfaceinto the polymer and by the UV absorbance of the coating and the
Fig. 10. Optical micrograph showing the surface of a coating exposed to a 5000 kJ/m2
dose of simulated outdoor treatment.
Fig. 11. FTIR data for coatings exposed to simulated outdoor treatment. Solid line: noexposure, dashes: 2500 kJ/m2 dose, dots: 5000 kJ/m2 dose. The peak at ~1000 cm−1 isattributed to network Si–O–Si; the peak at ~760 cm−1 to Si–CH3.
1913A. Kamer et al. / Thin Solid Films 519 (2011) 1907–1913
polymer. Thus, the most favorable place for microcracks to formwould be the interface region between the hard coating and thePMMA substrate. The mechanism of microcrack formation with UVexposure is likely to explain the significant drop we see in the fractureenergy of weathered samples.
While lower exposure doses led to a decrease in the fractureenergy of the hard coating, the 5000 kJ/m2 dose increased themeasured value. Optical characterization of the fracture surfacesrevealed that only the sample treated with the highest dose hadextensive substrate damage. The channel cracks on the filmpenetrated into the substrate and created additional stress concen-tration loci deeper in the substrate. This increased plastic deformationduring fracture tests and the measured fracture energy was higherwhen compared to an undamaged film and substrate. In summary, theincrease in the resistance to fracture of a coated window exposed to5000 kJ/m2 dose was a direct result of substrate damage, not evidenceof interface healing.
4. Conclusion
We demonstrated that the mixed mode asymmetric DCB is areliable, quantitative method for measuring adhesion of hard coatingson polymers. It was found that introducing a small amount of mode IIloading helped to keep the crack at the interface of a hard coating anda PMMA substrate. Time dependent crack growth was observed inasymmetric DCB and channel cracking tests and both methodsproduced similar results. The observed time dependent crack growthat subcritical crack driving forces was mostly due to viscoelasticdeformation of the PMMA substrate near the crack tip. Crack growthat high driving forces was insensitive to changes in relative humidity;
additional investigation is needed to explain the behavior at lowdriving forces. Simulated outdoor exposure at low doses significantlydecreased the fracture energy without macroscopic damage, whilethe highest dose resulted in cracking of the coating and extendeddamage in the substrate. FTIR analysis of weathered coatings showedthat polysiloxane based coatings cross-linked and lost organic groupsas a result of exposure to UV light.
Acknowledgements
The authors would like to thank the Boeing Company for providingmaterials and funding for this study. We are grateful to Mr. Vasan S.Sundaram (Boeing Company) for helping initiate this work and for hissupport.
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