Interfacial and Annealing Effects on Primary α-Relaxation ofUltrathin Polymer Films Investigated at NanoscaleHung K. Nguyen,*,† Massimiliano Labardi,‡ Simone Capaccioli,†,‡ Mauro Lucchesi,†,‡ Pierangelo Rolla,†,‡

and Daniele Prevosto‡

†Dipartimento di Fisica “Enrico Fermi”, Universita di Pisa, Largo Pontecorvo 3, 56127 Pisa, Italy‡CNR-IPCF, Consiglio Nazionale delle Ricerche, Istituto per i Processi Chimico-Fisici, c/o Dip. Fisica Largo Pontecorvo 3,56127 Pisa, Italy

ABSTRACT: The influence of interfacial interactions and annealing time on dynamics of the α-relaxation in ultrathin poly(vinylacetate) films deposited on different substrates has been studied using local dielectric spectroscopy at ambient pressure andcontrolled humidity. After annealing at 323 K for about 3 days, for polymer films supported on gold and aluminum substrates,an increase of the relaxation rate with decreasing film thickness below 30−35 nm was observed, whereas for films deposited onsilicon substrates a thickness-independent dynamics was found for films as thin as 12 nm. The difference in size effect ondynamics of the films could reasonably be related to the difference in interfacial energy between polymer films and substrates,even though a criterion simply based on interfacial energy cannot be used to explain all the results. In fact, further annealing at ahigher temperature evidenced an annealing-dependent dynamics in films prepared on aluminum substrates consistent with thepresence of long-living metastable states at the polymer/substrate interface. The lifetime of such metastable states seems relatedto the nature of the substrate as well as to the molecular weight of the polymer.

1. INTRODUCTIONEffects of interfacial interactions on relaxation dynamics ofpolymeric materials under confinement are considered as a keyfactor for understanding the deviation of relaxation propertiesand glass transition temperature, Tg, from the bulk behavior.1−3

It is suggested that at the free surface of a polymer film a layerof few nanometers of thickness exists,1,2,4 in which chainmotions are faster than in the bulk system. Moreover, depend-ing on the degree of the polymer/substrate interaction, themobility of polymer chains at the interface with the substratecan be increased, decreased, or remain the same as in the bulk.3,5

The most used parameter to estimate the degree of interfacialinteractions is the interfacial energy between the substrateand polymer, γsp. It is generally accepted that with increasinginterfacial energy the segmental mobility of polymer chainsin the regions close to the substrate slows down, leading toincrease the Tg in such regions compared to that in the bulk.6−9

Recently, such a scenario has been strongly criticized by aseries of works showing that polymer relaxation dynamics, andconsequently the dynamic glass transition, is independent ofthe film thickness, even in the nanometer range, as well as on

polymer/substrate interactions.10−12 Variations ascribed todifferent sample preparation procedures were supposed to bethe main factor leading to observed deviations of dynamicsof ultrathin films from the bulk.12 For example, thickness-independent dynamics has been found in poly(vinyl acetate)(PVAc) films as thin as 13 nm when subjected to an annealingprocedure lasting for several tens of hours at a high temperature(366 K, i.e., ∼Tg(bulk) + 50 K).11 Surprisingly, such annealingtime is much longer than any conventional time scale usedto determine physical processes of polymers, for example, theprimary α-relaxation or reptation relaxation time at the anneal-ing temperature. On the contrary, polymer films annealed atlower temperatures and/or for shorter time usually confirmthe presence of the size-dependent dynamics, as reportedby broadband dielectric spectroscopy (BDS) on capped anduncapped films9,13−15 and by other techniques on supportedfilms.3,16,17

Received: December 21, 2011Revised: January 27, 2012Published: February 13, 2012

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Besides sample preparation issues, the need of annealingover a very long time has led some authors to hypothesize thepresence of long-living metastable states in supported polymerfilms. In this respect, the existence of an interfacial polymerlayer in contact with the supporting substrate was discussedseveral years ago18 and recently measured in silicon-supportedpolystyrene (PS) films as a function of the annealing time.19

Moreover, in a recent report Napolitano et al.20 have evidencedthe correlation between the annealing-dependent relaxation dy-namics of ultrathin PS films deposited on aluminum substratesand the annealing-controlled growth of an irreversibly adsorbedpolymer layer at the substrate surface. A dimensionless numbergiven by the ratio between the time scale of the adsorption andthe annealing time was introduced to describe the growth of theadsorbed layer and consequently the annealing-dependentdynamics observed in ultrathin polymer films.20 The proposedmechanism seems to be suited to reconcile the results obtainedon several Al-supported polymer films under different anneal-ing conditions11 and to explain on a physical basis the conclu-sions drawn by Erber et al.12 about the importance of samplepreparation procedures.In recent years, a new method named local dielectric

spectroscopy (LDS)21 has been successfully applied to measuredielectric relaxation of supported polymer films. The LDSmethod was initially implemented to measure primary α-relax-ation on thick PVAc films (thickness ∼1 μm) with nanometerscale resolution under ultrahigh-vacuum conditions.21 Soonafterward, the method has been developed to be applied onultrathin films under ambient conditions and with controlledhumidity.17,22,23 A benefit of LDS method is that it can beused to measure dielectric relaxation on uncapped films, thusavoiding problems related to the evaporation of the cappingelectrode, such as shorts and metal particles diffusion inside thefilm. Moreover, by such technique dielectric properties of poly-mer films are measured with a lateral positioning resolution oftens of nanometers, releasing the need of continuous and uni-form ultrathin polymer films over extended areas, also allowingLDS to be applied to detect inhomogeneities of dynamicbehavior of the sample on the nanometer scale.22,24

In the current work, we apply LDS developed by using animproved frequency-modulated electrostatic force microscopy(FM-EFM) method to get access to a frequency interval up tofour decades, to study with greater accuracy confinement effectson dynamics of uncapped ultrathin PVAc films deposited ondifferent substrates, namely gold, aluminum, and silicon. By asystematic study of samples annealed in the same conditions,the annealing effects on the primary α-relaxation of the filmsare also addressed. In addition, interfacial energies betweenPVAc and supporting substrates are determined and correlatedto the effects of nanoconfinement on the dynamics.

2. EXPERIMENTAL METHODS2.1. Sample Preparation. Poly(vinyl acetate) (PVAc), with

molecular weight Mw = 350 kg/mol, polydispersity index PDI = 2.80,and gyration radius Rg ∼ 17 nm,25 was purchased from ScientificPolymer Products, Inc. Ultrathin PVAc films were prepared by spin-coating solutions of PVAc in toluene onto gold, aluminum, and siliconsubstrates. Gold layers of 30 nm thickness were obtained by thermalevaporation on glass disks previously evaporated with a ∼5 nm adhe-sion layer of chromium, whereas 50 nm thick aluminum layers weredirectly evaporated on glass disks. Silicon substrates were obtainedfrom standard monocrystalline, doped Si(100) wafers for micro-electronics use. The film thickness was controlled in the range from

8 to 233 nm by changing the concentration of polymer solutions aswell as spinning speed.

Prior to dielectric measurements, all polymer films were annealedwith a similar procedure at 323 K under vacuum for about 3 days. Wewould like to clarify that it is not our intent to propose a new criterionof annealing. Our choice is merely based on a previous experimentalprocedure,17 used on PVAc films with Mw = 167 kg/mol, that wasfound sufficient to completely remove the residual solvent and mois-ture from our thickest film. After such common procedure some filmsused for studying annealing effects have been annealed at highertemperature, which will be detailed in section 3.

Before measurements, a scratch was made on each film by a steelcutter in order to uncover part of the conductive substrate beneath.Calibration spectra were taken on the conductive substrate and used asreference to obtain local dielectric spectra of the polymer. Moreover,AFM profiling of the scratch was used to determine the film thickness.

2.2. Interfacial Energy Calculation. In order to determine theinterfacial energies, γsp, between PVAc and solid substrates, contactangles of three liquids, namely water, glycerol, and diiodomethane, oneach substrate were measured using a CAM 200 optical contact anglemeter (KSV) at room temperature. The values of the contact angles,φ, were obtained from averaging five measurements for each sample(Table 1).

The dispersive component, γLW, the electron acceptor, γ+, and theelectron donor, γ−, contributions to the surface energy of a liquid and asolid are related to the contact angle of a droplet of the liquid on thesurface of the solid, through the Young−Dupre equation26

+ ϕ γ = γ γ + γ γ + γ γ+ − − +(1 cos ) 2( )L sLW

LLW

s L s L (1)

where “s” and “L” indicate the substrate and liquid, respectively. Usingmeasured contact angles of three liquids and their previously deter-mined surface tension parameters,27 surface energies of the solids werecalculated by eq 1.

Interfacial energies of PVAc film with the three substrates wereestimated using the well-accepted Good−Girifalco−Fowkers combin-ing rule28

γ = γ − γ + γ γ + γ γ − γ γ

− γ γ

+ − + − + −

− +

( ) 2(

)

sp sLW

pLW 2

s s p p s p

s p (2)

where “p” denotes the polymer. The three surface tension parametersof PVAc were determined from the literature.29

2.3. Local Dielectric Spectroscopy Technique. In this work, aVeeco Multimode atomic force microscope (AFM) (Nanoscope IIIawith ADC5 extension) was adapted to a frequency-modulated electro-static force microscope (FM-EFM) and operated in lift mode as pre-viously described.22 FM-EFM with improved bandwidth was hereimplemented through a RHK Technology PLLProII phase-locked-loop (PLL) frequency detector, having a nominal response bandwidthof 4 kHz that could be extended to about 10 kHz by appropriatesetting of internal filters and DAC sampling rates. In essence, suchbandwidth indicates the limit frequency at which the PLL detector isable to follow a rapid change of the resonant frequency of the AFMcantilever. The latter frequency is made to shift by the effect of the

Table 1. Measured Contact Angles of Water, Glycerol, andDiiodomethane on Solid Substrates and CalculatedInterfacial Energies with PVAc

contact angle φ [deg]

substrate water glycerol diiodomethane

interfacialenergy with

PVAc [mJ/m2]

gold 89.2 ± 0.9 74.6 ± 1.5 33.4 ± 1.8 1.2 ± 0.2aluminum 27.9 ± 1.3 39.2 ± 3.1 28.7 ± 0.8 1.6 ± 0.3silicon 41.9 ± 2.4 32.2 ± 1.9 57.3 ± 1.3 3.4 ± 0.3

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electric force gradient induced by biasing of the tip. Such gradientis proportional to the term d2C/dz2, where C is the tip/samplecapacitance.17

As demonstrated by the spectra herein reported, we were able toacquire spectra with electric excitation frequency as high as 5 kHz,thereby extending our previous bandwidth17 by 2 decades and there-fore up to 4 decades in total. This was possible while not giving upthe high sensitivity and spatial resolution obtainable by frequency-modulation EFM30 in comparison to different methods based ondirect force measurement.23

For electric excitation and acquisition of dielectric spectra, anexternal dual-phase lock-in amplifier (SRS SR830DSP) was used,controlled through a GPIB interface by a homemade LabView routine.Doped silicon, Pt-coated AFM cantilevers (Nanosensors PPP-NCLPt)were used, with spring constant k ∼ 40 N/m, resonant frequency f res ∼165 kHz, and nominal tip radius of around 25−30 nm. Oscillation am-plitude was calibrated as customary by AFM tapping mode amplitude/distance curves. By this way, the distance between the AFM tip and thesample surface during the electric measurements was calculated bysumming the average tip−sample distance on tapping mode with theset-lift height. For all dielectric measurements, the tip/sample distancewas estimated around 20 nm.Our microscope was operated in controlled humidity within a

homemade enclosure and during measurements the relative humiditywas about 3−4%. The sample temperature was controlled using athermal application controller (TAC, Veeco). The actual temperatureon the sample surface was calibrated using a PT100 sensor.

3. RESULTS AND DISCUSSION3.1. Dielectric Measurements by LDS. In LDS, dielectric

relaxation of polymer films is represented in terms of an electricphase shift angle, δv, the tangent of which is related to the tip/substrate capacitance by

δ = ∂ ″ ∂∂ ′ ∂

C zC z

tan/

/v

2 2

2 2 (3)

where z is the distance between the tip and the sample surfaceand C′ and C″ are the real and imaginary parts of the tip/substrate capacitance. As described in previous reports,22,23,31,32

in the distance range of the order of the tip radius, R, the tip/substrate capacitance can be conveniently expressed by approxi-mated relations. We adopt the one proposed by Fumagalli et al.31

ω = πε + − θ+ ε ω

⎧⎨⎩⎫⎬⎭C z R

Rz h

( , ) 2 ln 1(1 sin )

/ ( )0(4)

where h and ε(ω) are the thickness and the relative dielectricpermittivity of the polymer film, respectively, θ is the aperturehalf-angle of the tip shaft, assumed of conical shape, and ε0is the vacuum permittivity. Noticing that the capacitance (eq 4)is not linearly dependent on ε, thus the quantity tan δv is de-pendent on the material properties as well as on the geometryof the capacitor. In other words, tan δv is a different quantityfrom the tangent of the loss angle usually considered in con-ventional dielectric spectroscopy, even if their temperature andfrequency dependencies are qualitatively very similar.In Figure 1, the tangent of δv, in the following also called

loss-tangent, obtained in a frequency range from 0.5 Hz to5 kHz and at different temperatures for a 233 nm thick filmdeposited on aluminum substrate is presented. The spectrashow the presence of a relaxation peak that moves towardhigher frequencies with increasing temperature in a similar wayof dielectric loss spectra obtained by BDS.14 A measurementaround the Tg value (∼310 K) is also presented to show thebackground of the instrument. At Tg the secondary relaxation

peak is at much higher frequency and the α-structural relaxationpeak is at lower frequency than the measured interval, so thatonly its high-frequency tail is observable. By α-relaxation orstructural relaxation we refer to the process associated with theglass transition and cooperative segmental dynamics. The solidcurves represent a fitting performed by using analytical modelsdescribing the substrate/tip capacitance (eq 4), the phase shift ofthe tip resonant frequency modulation (eq 3), and the Havriliak−Negami relation, often used in the phenomenological descriptionof the dielectric function of a polymer:

ε ω = ε + Δε+ ωτ

∞ −α βi( )

[1 ( ) ]HN1 HN HN (5)

In our fitting procedure we fixed ε∞ as constant as that ob-tained for bulk samples (ε∞ = 3.1),33 while the other para-meters in eq 5 were adjusted to get the best fits that wereevaluated by checking the residues between the experimentaldata and fitting curves. As shown in Figure 1, fitting curvesinterpolate well the measured data at all temperatures. We pointout that the parameters obtained from the analysis describe thedielectric function ε(ω) of the material and consequently arecomparable to the parameters obtained from conventional di-electric measurements.The shape parameters of the dielectric function of eq 5

obtained by fitting are almost temperature-independent withinthe fitting uncertainties (data not shown), while the peak fre-quency of the α-structural relaxation process increases withincreasing temperature as expected due to the enhancement ofsegmental mobility at higher temperatures. The data measuredon the 233 nm thick film will be assumed as the bulk value.Indeed, in a previous study17 we have found that the relaxationrate measured on PVAc film (with lower molecular weight thanin the present paper) of 137 nm thickness by LDS is similarto that of the bulk PVAc sample with a similar molecularweight measured by BDS. Moreover, by using a similar tech-nique, Schwartz et al.23 have observed a thickness-independentdynamics (once corrected for the geometrical contribution tothe spectra) on films with thicknesses from 250 to 1000 nm,confirming that the thickness-independent regime is alreadyreached at 250 nm.

Figure 1. Loss-tangent spectra on a 233 nm thick PVAc film,deposited on aluminum, measured at different temperatures. Linesrepresent fitting with eqs 3, 4, and 5.

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The logarithm of maximum frequency of the dielectric loss,ε″, as a function of temperature, has been fitted by the Vogel−Fulcher−Tammann (VFT) equation (Figure 2):

ν = ν −−

DTT T

log logmax 00

0 (6)

where νmax is the frequency of the maximum dielectric loss, ν0 isthe relaxation frequency of the α-process at infinite tempera-ture, D is a constant anticorrelated to the fragility parameter,and T0 is the so-called Vogel temperature.34 The VFT para-meters resulting from fitting of such data, as well as the glasstransition temperature, Tg, according to the definition νmax(Tg) =1/(200π) Hz, are listed in Table 2. The good agreement of

such parameters with those obtained for the bulk samples ofPVAc35,36 using BDS supports the fact that the loss-tangentspectra obtained by LDS represent the α-structural relaxation ofthe polymer film, and the used models are working well at leastin the present experimental conditions.Finally, we mention that LDS technique has been proposed

as a surface technique, but the estimation of the thickness ofthe surface layer to which it is sensitive is matter of debate.21,37

According to our understanding based on previous investi-gation from us and others17,23 in the thickness range here inves-tigated, the data report information averaged over the entirethickness.3.2. Effects of Film Thickness and Supporting

Substrate. The size effect on the α-structural relaxation ofAl-supported PVAc films is presented in Figure 2 where thelogarithm of the frequency of the maximum of ε″ measured onfilms with thickness from 233 to 13 nm is reported as afunction of reciprocal temperature. The relaxation rate is almostthickness-independent from 233 to 57 nm, but it increases on

the film of 23 nm, and even more on the thinnest film of13 nm. This result is quantitatively consistent with the findingsreported in our previous study about the confinement effect onthe α-structural relaxation of gold-supported PVAc films withlower molecular weight.17

As shown in a variety of reports, interactions at surfaces andinterfaces play an important role in ruling the deviations ofdynamics of confined polymers from the bulk behavior.6,8 Bychanging the supporting substrates, the interplay betweenpolymer/substrate interaction and size effects on α-structuralrelaxation of ultrathin films can be observed. In this study, threedifferent substrates, namely gold, aluminum, and silicon, whichare very commonly employed in the study of supported ultrathinpolymer films, were used. Loss-tangent spectra obtained at thetemperature of 327.4 K for PVAc films on gold and aluminumsubstrates and 328.2 K on silicon are reported in Figures 3a−c,

respectively. As clearly visible, a shift of the relaxation peak tohigher values of frequency is observed on gold- and aluminum-supported PVAc films (Figures 3a,b). The shift becomes morepronounced on thinner films, and spectra seem also to becomebroader. In contrast, there is no significant change in the relaxa-tion peak of all films deposited on silicon substrates with thick-ness from 85 nm down to 12 nm (Figure 3c).38

Figure 2. Logarithm of maximum frequency of the α-structuralrelaxation in PVAc films deposited on aluminum, as a function ofreciprocal temperature. Solid line represents VFT fitting for the film of233 nm thickness.

Table 2. VFT Fitting Parameters for the 233 nm Thick FilmSupported on Aluminum in Comparison with Bulk Data onPVAc with Mw = 170 kg/mol

sample log ν0 [Hz] D T0 [K]Tg (νmax = 1/200π Hz)

[K]

233 nm 11.8 (fixed) 6.3 ± 0.1 261.1 ± 0.9 310.0 ± 1.1bulk35 11.8 ± 0.1 6.3 ± 0.1 261.4 ± 1.0 310.4

Figure 3. Isothermal loss-tangent spectra of PVAc films with differentthicknesses deposited on (a) gold (measured at 327.4 K), (b) aluminum(measured at 327.4 K), and (c) silicon (measured at 328.2 K).

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The effects of film thickness and supporting substrate on theshape parameters of the relaxation peak were also analyzed,as shown in Figure 4. The values of αHN and βHN both slightly

increase with decreasing thickness of films supported on goldand aluminum, meaning that the α-structural relaxation on suchsubstrates becomes more symmetric and broader on decreasingthe film thickness. However, for films deposited on silicon, onlyαHN seems to increase with decreasing film thickness, while βHNremains independent of the sample thickness. In principle, theobserved broadening of the relaxation peak can be ascribed to amore heterogeneous dynamics, which is consistent with thepresence of layers with different mobility. The observed in-crease of the symmetric broadening of the spectra in samplesprepared on all substrates should be contrasting with the hypoth-esis of a selective suppression of slow modes.14

In Figure 5, the difference in the value of Tαdefined asthe peak temperature of the α-structural relaxation process

measured at 1 Hzcompared to the 233 nm thick film isplotted as a function of the film thickness. For films depositedon gold and aluminum substrates, there was almost no changein the Tα when the film thickness was decreased down to about35 nm, while a systematic reduction of Tα was observed withfurther decreasing film thickness, and the Tα-reduction mea-sured on films as thin as 10 nm was about 2−3 K. Even though

the reduction of Tα obtained here is small, somehow compa-rable to the measurement uncertainties, such reduction occursin a systematic way on both gold and aluminum substrates.Furthermore, a small change in dynamics of PVAc with de-creasing film thickness was found in previous reports, also usingdifferent techniques, on samples with lower molecular weightand different procedures of sample preparation.14,22,39 Forexample, using the microbubble inflation method to determinethe creep compliance of freely standing PVAc films, O’Connellet al. found a reduction of the Tg from the bulk value of lessthan 1.5 K on the thinnest film of 23.7 nm.39

Films prepared on Si substrates did not show change ofdynamics for film thicknesses down to 12 nm. This observationagrees with BDS results obtained with the air-gap capacitorgeometry by Serghei et al.10 Indeed, after annealing the sampleswell above the Tg in a pure nitrogen atmosphere for severalhours until the reproducible measurements were reached, noshift of the relaxation rate respect to the bulk was found onseveral polymers studied in the uncapped geometry, amongwhich ultrathin PVAc films supported on silicon. Interestingly,these results were apparently contrasting with the size effectson relaxation dynamics of capped PVAc films between twoaluminum electrodes previously reported by the same authorsusing BDS.14 Indeed, they found a reduction of the Tα, corres-ponding in this case to the frequency of 38 Hz, of about 2−3 Kfrom the bulk value when the film thickness decreases downto about 10 nm, in agreement with our present observation onAu- and Al-supported films. Such results suggest that the discre-pancy in refs 10 and 14 could not be due to the evaporatedaluminum electrode or to a noneffective removal of solvent inthe Al-supported films, but more probably to the difference inpolymer/substrate interactions. Moreover, we can confirm thatthe confinement effects on PVAc relaxation dynamics are muchweaker than those on other polymers such as PS40 or poly-(methyl methacrylate) (PMMA).41 The difference might be theresult of diverse macromolecular structures or arrangements ofpolymer chains close to the free surface or to the interface. Arecent study by Kim et al.42 addressed the importance of watercontent in PVAc when measuring in ambient conditions to thestrength of the Tg reduction with decreasing film thickness.However, in our measurements at relative humidity of 3−4%this water content is very small, less than 0.5% as found byMiyagi et al.43 for PVAc bulk. Therefore, we can deduce thatthe humidity-induced change of Tg or Tα herein is smaller thanour experimental errors.In order to understand how interfacial interactions impact

the Tα reduction, the interfacial energy, γsp, between PVAc filmsand the three solid substrates was also obtained (Table 1).The results show a similarity in interfacial interactions of PVAcfilms with gold and aluminum surfaces, which are less thanthe critical value of about 2 mJ/m2 estimated in Fryer’s report6

separating the energy region where a negative deviation ofTg in ultrathin films was observed (γsp < 2 mJ/m2) from thatof positive deviation (γsp > 2 mJ/m2). Noteworthy, for Au-and Al-supported films the values of γsp are very similar toeach other and they correspond to similar ΔTα, whereas forSi-supported films we have larger values of γsp and smaller ΔTα

compared with the two other substrates (gold and aluminum)(Figure 5).The value of γsp measured on the silicon substrate is much

higher than the critical value, indicating a strong interactionbetween PVAc and silicon surface. Fryer et al.6 have found anincrease of the Tg on both PS and PMMA films with decreasing

Figure 4. Shape parameters of the α-structural relaxation (eq 5) as afunction of film thickness.

Figure 5. Thickness dependence of the difference in the Tα betweenultrathin films and the 233 nm thick film. The dashed and solid linesare guides to the eye.

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film thickness compared to the bulk value when the polymer−substrate interfacial energy was higher than 2 mJ/m2. However,in our case, instead of an increase of the Tα we observed nosignificant Tα shift for all films deposited on silicon. Therefore,we confirm that samples made of the same material, and an-nealed in the same way on different substrates, exhibit confine-ment effects that can be qualitatively related to the interfacialenergy. Nevertheless, the comparison of our results with thosefrom the literature evidence that γsp alone is not sufficient toexplain the variation of Tg from the bulk with decreasing filmthickness, as also mentioned elsewhere.44 This result is not sur-prising as differences in molecular structure and chain stiffnessare factors affecting polymer dynamics in the bulk and likelyalso the confinement effects on dynamics. This is evidencedfor example by Priestley et al.,8 who found that the change ofTg at the free surface and substrate interface of polymer filmsdepends substantially on the structure of polymers. However,the mechanism behind the interaction between polymer andsupporting substrate still awaits to be clarified in detail.3.3. Annealing Effects. The dependence on annealing of

the α-structural relaxation time and glass transition was evi-denced on Al-capped PVAc films,11 and it was inferred thatsuch dependence might be due to the growth of an absorbedpolymer layer at the polymer/substrate interface as for the caseof the Al/PS system.20,45 We considered interesting to inves-tigate such aspect on PVAc films deposited on substrates withquite different interfacial energies. As a consequence, samplesannealed at 323 K for 3 days (first annealing) were furtherannealed at 366 K for different times ranging from 20 h to 8days. For the 226 nm thick PVAc film deposited on aluminum,the α-structural relaxation is found independent of further an-nealing; i.e., the loss-tangent spectra measured before and afterannealing up to 6 days at 366 K are unchanged (Figure 6a).In contrast, the relaxation rate of the structural process on the23 nm thick Al-supported PVAc film decreases slightly after3 days of annealing and eventually decrease to that of the thickfilm after 6 days of annealing. These results demonstrate thatfurther annealing seems to affect the dynamics of the ultrathinfilm only. As a simple hypothesis, it is suggested that the regionof polymer located at the layers close to the free surface orpolymer/substrate interface evolves through a series of meta-stable states characterized by different time scales. The latterpossibility is in agreement with, and in case extends, the resultsobtained by Napolitano et al.,20 evidencing that the influenceof annealing on the dynamics of ultrathin polymer films iscorrelated to the growth of an irreversibly adsorbed layer at thesubstrate interface and to packing of macromolecules in suchlayer.We have also observed that the time scale of the evolution

of such metastable states in PVAc films during annealing isrelated to the molecular weight, as already found for PS.20 Thedisappearance of dynamics enhancement in our sample ofMw =350 kg/mol and thickness 23 nm was obtained with more than3 days of annealing at 366 K (plus the first annealing at 323 Kfor 3 days), as seen in the inset of Figure 6a showing the timeevolution of the relative maximum frequency of the α-processmeasured on 226 and 23 nm thick films at 327.4 K, where n =νmax(t, h)/νmax(0 s, 226 nm). However, only 32 h was necessaryfor the 13 nm thick PVAc film with Mw of 157 kg/mol11 torecover bulklike dynamics. The Mw dependence of the evolu-tion time of metastable states is found more than linear, butit appears weaker than that observed on PS20 in the samerange of molecular weights. In fact, Napolitano et al.20 found

that increasing the molecular weight from 97 to 160 kg/mol(both larger than the critical value for entanglement) thecharacteristic time increase from less than 3 h to more than6 weeks (more than a factor 100). In our case for an increase ofa factor 2 of Mw the increase of the characteristic time is a bitlarger than a factor 3.In order to clarify whether the interfacial energy between the

polymer and the supporting substrate has actually a role in theannealing effects observed above, a similar annealing procedurewas performed on the 26 nm thick film deposited on the siliconsubstrate. As shown in Figure 6b, no difference in the peakposition among loss-tangent spectra obtained at different an-nealing times could be observed, suggesting that annealing at323 K for 3 days is sufficient to equilibrate polymer chains evenat interfacial layers. This result also suggests that the adsorptiontime of PVAc onto substrates can depend on the polymer/substrate interactions as well as on substrate roughness. Furtherinvestigation is in progress to elucidate this point.

4. CONCLUSIONS

Confinement effects on the α-structural relaxation of supportedultrathin PVAc films on a frequency range of up to 4 decadeswere studied by means of an improved local dielectric spectro-scopy approach. Relaxation dynamics was measured for PVAcfilms deposited on different substrates and subjected to differ-ent annealing conditions. After the first annealing step (at 323 Kfor 3 days), faster relaxation was observed on films supportedon gold and aluminum substrates when film thickness decreasedto below 30−35 nm. The maximum observed Tα reduction fromthe bulk value was 2−3 K for film with thickness of about10 nm. In contrast, no thickness-dependent relaxation wasfound for all films deposited on silicon substrates with thickness

Figure 6. Loss-tangent spectra measured after different annealingtimes at 366 K of PVAc: (a) on Al measured at 327.4 K, the insetshows the evolution of the relative maximum frequency, n, uponannealing time on both thicknesses; (b) on Si measured at 328.2 K.The data have been normalized to the maximum and rescaled forclarity.

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from 85 nm down to 12 nm. Such difference can be reasonablyrelated to the difference in interfacial energy between polymerfilms and substrates, which is larger for silicon than for gold andaluminum. However, a criterion simply based on interfacialenergy cannot be used to explain all the results. In fact, byannealing at higher temperature (366 K) up to 8 days, no shiftof loss-tangent spectra was observed for the 226 nm thick filmon aluminum as well as on the 26 nm thick film on silicon.However, an effect was observed on the 23 nm thick film onaluminum; in particular, the bulk dynamics was re-covered after 6 days of annealing. Our results are consistentwith the presence of long-living metastable states of the poly-mer at the interfacial layer close to the substrate. We find thatthe characteristic time of such states may depend on the mole-cular weight of the polymer as well as on the nature of thesubstrate. These findings help to reconcile literature dataobtained on supported and capped thin PVAc films10,14 thatappeared conflicting.

■ AUTHOR INFORMATION

Corresponding Author*E-mail: nguyen@df.unipi.it.

NotesThe authors declare no competing financial interest.

■ ACKNOWLEDGMENTS

We thank Dr. P. Pingue and M. Cecchini (Scuola NormaleSuperiore, Pisa) for assistance in sample preparation andcontact angle measurements and M. Bianucci (University ofPisa) for assistance in developing our LDS setup. We also thankProf. K. Fukao (Ritsumeikan University) for information aboutRg of PVAc. H.K.N. thanks the Galileo Galilei School ofGraduate Studies at Pisa University for financial support.

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Macromolecules Article

dx.doi.org/10.1021/ma202757q | Macromolecules 2012, 45, 2138−21442144