Polymer Physics (From Suspensions to Nanocomposites and Beyond) || Dynamics of Materials at the Nanoscale: Small-Molecule Liquids and Polymer Films

5DYNAMICS OF MATERIALS AT THENANOSCALE: SMALL-MOLECULELIQUIDS AND POLYMER FILMS

Gregory B. McKennaDepartment of Chemical Engineering, Texas Tech University, Lubbock, TX, USA; EcoleSuperieure de Physique et de Chimie Industrielles de la Ville de Paris, Laboratoire dePhysicochimie de Polymeres et des Milieux Disperses, Paris, France

5.1 Background and introduction

5.2 Dynamics and the glass transition of small molecules at the nanoscale5.2.1 Early studies5.2.2 Dynamics of small molecules at the nanoscale: ups and downs

5.3 Ultrathin polymer films

5.4 Aging in confined systems5.4.1 Physical aging basics5.4.2 Physical aging and structural recovery at the nanoscale

5.5 Summary and perspectives

5.1 BACKGROUND AND INTRODUCTION

The behavior of materials at the nanoscale has become a subject of intense interestfor several reasons. First, it was a bit of a laboratory curiosity when my co-workerCatheryn Jackson and I [Jackson and McKenna, 1991a, 1996] and, nearly simultane-ously, Jiri Jonas and his co-workers [Liu et al., 1991] discovered that glass-forming

Polymer Physics: From Suspensions to Nanocomposites and Beyond, Edited by Leszek A. Utracki andAlexander M. JamiesonCopyright © 2010 John Wiley & Sons, Inc.

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192 DYNAMICS OF MATERIALS AT THE NANOSCALE

organic liquids, upon confinement to extremely small (now called nanoscale) pores,exhibit reduced glass transition temperatures. In those works the reduction of the Tg

was less than 18 K even when the pore diameter was on the order of 4 nm. Severalyears later, Keddie et al. [1994a,b] found similar behavior for ultrathin polystyrene(PS) films supported on silicon substrates. As indicated, these works were interest-ing because they were not readily explained by conventional theories of the glasstransition, but they were not obviously of great practical importance. This situationchanged dramatically when the group at the University of Guelph in Ontario, Canada,under the direction of John R. Dutcher [Forrest et al., 1996, 1997, 1998], developeda method to make measurements on freely standing (both sides exposed to air) poly-mer films that indicated a reduction of the glass transition temperature in polystyreneof over 70 K. Because polymers are used widely in microelectronics applications,and this field was pushing the nanoscale in feature sizes, such large changes in theglass transition implied potential difficulties in actual applications where mechanicalstability or dielectric loss properties might become important.†Although well over1000 papers have been published on the behavior of glass-forming materials at thenanoscale, it is still fairly true to say that we do not fully understand this behavior.Furthermore, there is interest not only in the behavior of glass-forming materials at thenanoscale [Alcoutlabi and McKenna, 2005] but also in phase transitions at this sizerange [Alba-Simionesco et al., 2006]. Here we limit ourselves to the behavior of glassyand glass-forming systems. In the following sections I provide a look at the behaviorsof confined materials observed and how we perceive them. To provide a referenceframe, I begin with a section on the dynamics of glass-forming liquids in confine-ment by discussing the developments relevant to small-molecule glass-formers. Inthis section I provide a view of what is known and what is controversial, as well as aperspective of what can be or needs to be done to resolve the issues. Next, I describethe behavior of ultrathin polymer films. This second part is important because poly-mers can actually be formed into freely standing films and therefore experiments canbe performed with polymers that are not readily performed with unentangled liquidswhich are unstable once heated above the glass transition temperature, due to sur-face tension effects. The long-chain nature of polymers stabilizes the freely standingstructures.

The penultimate section of the chapter deals with the aging of materials at thenanometer size scale and in confinement. We see that this is an area that is justbeginning to be explored and that, perhaps, there is as much difficulty in findingcomplete answers as there is in determination of the behavior of the Tg in confinement.Finally, I provide a summary of the chapter along with a perspective of where we areand where the field should be going.

†It is important to remark that the mechanical properties of polymers are very sensitive to temperature.Hence, a material such as PS with a Tg values of approximately 100◦C, at ambient temperature is at nearly80% of the Tg value on an absolute temperature scale. Changing the Tg by 65 K has a large effect on thisratio, and the result is a potentially large reduction in the polymer mechanical properties, such as yieldstrength, modulus, and creep resistance. Dielectric response is similarly affected by proximity to the glasstemperature.

DYNAMICS AND THE GLASS TRANSITION OF SMALL MOLECULES 193

5.2 DYNAMICS AND THE GLASS TRANSITION OF SMALLMOLECULES AT THE NANOSCALE

5.2.1 Early Studies

Organic Liquids in Controlled Pore Glasses As indicated above, the first work thatidentified nanometer-length scales as being important to the glassy response of smallmolecules came from the present author’s labs. In the specific work of Jackson andMcKenna [1991a, 1996] it was found that the imbibing of two organic liquids [ortho-terphenyl (o-TP) and benzyl alcohol] into the pores of a controlled pore glass (CPG)lead to the observation of a reduction in the glass transition temperature. The Jackson–McKenna results are shown in Figure 5.1. Here we see that there is a reduction forthe o-TP of approximately 18 K for the smallest pore size of 4 nm, while the benzylalcohol shows a change of only about 7 K. Similar results were reported by Liu et al.[1991] for a larger series of organic liquids.

The results shown in Figure 5.1 were, perhaps, most interesting because it isdifficult to explain them in the context of simple views of either free volume orconfigurational entropy theories [Williams et al., 1955; DiMarzio and Gibbs, 1958;Gibbs and DiMarzio, 1958; Adam and Gibbs, 1965; Simha and Somcynsky, 1969;Somcynsky and Simha, 1971; Ferry, 1980; McKenna, 1989; Angell et al., 2000]. Inthe case of the former, an 18 K decrease in Tg would require an increase in specificvolume due to the physical constraint of the pore walls that is not reasonable.† Anotherway of looking at this issue is that if Tg had increased rather than decreasing in theearly studies, there may have been much less interest in the problem.

Tg in Microemulsions Before moving forward in time to look at how the problemof the Tg reduction in confined geometries has been studied since the early 1990s,we take a look back at earlier evidence of the effects of confinement on the glasstransition event. There are two sorts of results that I am aware of and that are relevantto the problem. First, in a set of extremely clever experiments in the early 1980s,C.A. Angell and co-workers [MacFarlane and Angell, 1982; Angell et al., 1984;Dubochet et al., 1984] used a method of making microemulsions (now we wouldcall the smallest of these, nanoemulsions) for creating volumes of simple molecularliquids that were small enough that crystallization could be avoided and vitrifica-tion observed in otherwise extremely difficult to vitrify liquids such as benzene andtoluene. Since the purpose of the work was to find the Tg of the materials investigatedand compare it with the Tg obtained by, for example, extrapolating in concentration

†Here we assume that the dTg /dP = 0.26 K/MPa for o-TP. This would imply a negative hydrostatic pressureof nearly 70 MPa for a ∆Tg of 18 K. This corresponds to a volume strain of the o-TP of 2.7%, which seemshighly unlikely, as it would imply a sticking of the o-TP to the pore walls with no further volume change ata temperature approximately 37 K above the observed and reduced Tg . While packing effects could affectthis, these are beyond the present discussion of simple views of these theories. Similarly, the entropy ofthe confined liquid would be expected to decrease due for example, to orientation at the walls of the pores,and this would also lead to increases in the value of Tg . (See Jackson and McKenna [1991a], Atake andAngell [1979], Simon et al. [2002], and Angell and Qing [1989] for parameters used in these estimates.)


FIGURE 5.1 Tg versus reciprocal pore diameter for o-terphenyl and benzyl alcohol. (AfterJackson and McKenna [1991a].)

from binary mixtures of glass-forming liquids [Angell et al., 1978], there is not astrong conclusion to be extracted from the work about the effect of size on the Tg .However, it is important to note that the work gave little evidence of a strong decreasein the value of Tg in the microemulsions as the droplet size decreased. This may beimportant because the nature of the constraint of a microemulsion is different thanthat imposed by a controlled pore glass. The former is relatively soft and the latteris hard. This will be discussed in detail later when we examine macroscopic con-straints in the context of the work of R. Richert’s group [He et al., 2005], where theyhave taken the microemulsion technique and used it to study more fully the effectsof size and rigidity of the constraint on the dynamics of nanoscale glass-formingliquids.

Tg in Block Copolymers There is another set of experimental data in which an appar-ent change of the Tg was observed with decreasing size [Kraus and Rollman, 1976;Toporowski and Roovers, 1976; Gaur and Wunderlich, 1980; Krause et al., 1982a, b;Lu and Krause, 1982; Lu et al., 1982; Granger et al. 1986; Wang and Krause, 1987;Tyagi et al., 1988]. In this instance, block copolymers can be synthesized in whichmicro- or nanophase separation occurs between the two blocks. In some instances it isobserved that the glass transition can decrease, and in others it will increase. Althoughthere are multiple events that occur in such systems, such as partial miscibility of thetwo blocks, broad interfaces, and defects in the individual blocks due to differencesin density of the two phases, at least some of the changes in the Tg were attributed tohydrostatic tension and pressure effects that occur due to the mismatch between thecoefficient of thermal expansion differences between the two phases (one glassy andone rubbery) over an appropriate temperature range. When the continuous phase isrubbery, it is found that the Tg can increase due to a hydrostatic compression, while in


the opposite case a hydrostatic tension can cause a decrease in the Tg value. Data fromLee et al. [1998] for a polypropylene/polystyrene/styrene–ethylene–propylene blockcopolymer blends is in semiquantitative agreement with the postulated hydrostatictension effect. For a constrained rubber phase, the Tg can decrease by as much as10 K. Since in this case the constraint occurs at 373 K, which is the Tg of polystyrene,the ∆T of the rubbery phase with a Tg of approximately 273 K is 100 K, and thistranslates into a strain of 4% and a negative hydrostatic pressure of approximately60 MPa (assuming that the glassy phase is rigid). With a dTg /dP of approximately0.25 K/MPa, this would lead to a reduction of Tg of 15 K. Hence, to account for theobserved decrease of only 10 K, one could invoke the fact that the glassy polystyreneis not rigid (unlike the controlled pore glass), as its bulk modulus is only about threetimes the bulk modulus of the rubbery phase, so the constraints would not be as greatas just calculated. In addition, it is possible that the breadth of the glass transition alsomakes the difference between a 10 K observation and a 15 K estimate unimportant.Also, one anticipates that in the block copolymers, the constraints imposed on theblock can be different than in the case of the small polymers confined to pores due tochain connectivity effects.

One other aspect of the effect of size and constraint at the nanoscale on materialdynamics was first considered in block copolymers in the work of Donth [1984].Because the length scale of the block domains can be controlled by the molecularweight of the individual blocks, one can use this control to develop questions concern-ing the length scales that control the glass transition. Donth used data from Krause et al.for estimating the size of regions of cooperative dynamics in the polystyrene blocks.Although the initial Donth paper could not come to firm conclusions about the sizes ofthe “cooperatively rearranging regions,” it suggested an approach that has continuedto be of interest in the study of size and confinement effects in glass-forming systems.

Amorphous Phase in Semicrystalline Polymers Perhaps one of the reasons for thesurprise of observing the decrease in the glass transition value in small moleculesin controlled pore glasses in the Jackson and McKenna [1991a,1996] and Liu et al.[1991] publications is the prior knowledge one had that in semicrystalline polymers,it was virtually always observed that the amorphous phase either had the same glasstransition as when no crystals were present or that the glass transition broadenstoward higher temperatures, often dramatically [Struik, 1978; Menczel and Wun-derlich, 1981; Aharoni, 1998]. The portion of the amorphous chains constrained bythe rigid crystallites has also been interpreted as a new “rigid amorphous” state hav-ing a higher glass temperature [Menczel and Wunderlich, 1981]. The observation ofincreasing Tg is broadly accepted as being due to surface-induced constraints on themolecular motions in such systems and helps partially to explain the ups and downs ofthe glass transition in confinement discussed below. As mentioned above, an increas-ing Tg is understood and often expected. When Tg decreases, further investigationsseem to be required for full understanding.

Polystyrene Nanospheres One last early experiment is discussed here. Gaur andWunderlich [1980] investigated the behavior of polystyrene nanospheres using


FIGURE 5.2 Surface area effect on the broadening of the glass transition temperature ofpolystyrene nanospheres and lamellar phases in block copolymers. Open squares are forlamellar phases, and filled triangles are for nanospheres. (Modified figure based on originalfrom Gaur and Wunderlich [1980].)

differential scanning calorimetry. They found that the Tg was broadened toward lowertemperatures and interpreted the resulting increased mobility as being due to the freesurface in the polymer nanoparticles. As shown in Figure 5.2, the broadening wasfound to depend strongly on the surface area of the particles. Why particles and thelamellar blocks, similar to the works mentioned above, showed similar effects wasnot resolved.

5.2.2 Dynamics of Small Molecules at the Nanoscale: Ups and Downs

The results described above are already suggestive of some of the issues that arise indetermination of the influence of size and confinement on the behavior of materials atthe nanometer size scale. Nearly immediately subsequent to the work of Jackson andMcKenna [1991a] and Liu et al. [1991] there were several publications using polarliquids in controlled pore glasses that seemed to show the opposite result[Mel’nichenko et al., 1995; Schuller et al., 1995; Schonhals and Stauga, 1998a, b].That is, rather than small size or confinement decreasing the glass transition temper-ature, the evidence was that the dielectric response was retarded (i.e., equivalent toincreasing Tg ). Hence, as in the case of the polymer in the semicrystalline confinement,the confining surfaces seem to be able to increase the Tg by constraining the motion


FIGURE 5.3 Dynamic response measured dielectrically for propylene glycol (PG) confinedin 10-nm pores, showing bulk material (filled circles), a central core of material (open circles)and a new surface layer (open triangles). (From Mel’nichenko et al. [1995], with permission.Copyright © 1995 American Institute of Physics.)

of the molecules at the surface of the confining medium, in this case the pore walls.As an example, the results from Mel’nichenko et al. [1995] are shown in Figure 5.3,where we see both a slower central core of material and a new surface-constrainedlayer of material in the relaxational dynamics of the hydrogen-bonding propyleneglycol confined to a controlled pore glass. It is interesting that this work shows twodistinct dynamics. Subsequent work using differential scanning calorimetry (DSC)on the o-TP (nonpolar) system by Park and McKenna [2000] showed two distinctTg values for the liquid confined to CPGs of different size. However, as shown inFigure 5.4, for the non-hydrogen-bonding o-TP, a heat capacity jump at a tempera-ture below the bulk glass transition was observed along with a second, smaller stepat a temperature above the bulk Tg . It seems clear that the surface can slow down thedynamics. It is less clear what is happening in the core of the pore to reduce the Tg .Recently, Zheng and Simon [2007] investigated the effects of polar molecules andsurface treatments on Tg using calorimetry. In that work they found small decreasesin the glass temperature for glycerol and propylene glycol, being approximately 5◦Cfor the former and 3◦C for the latter, although for the propylene glycol they also foundtwo Tg values, with the higher Tg being approximately 30◦C higher than that of theunconfined fluid. They also provided a critical review of related works.

In 2000, the present author published an arrow chart in which he summarizedthe behavior observed for materials in nanoscale confinement [McKenna, 2000]. Anupdated and abbreviated version of this chart was used by R. Richert [2006] at a recent


FIGURE 5.4 DSC thermograms for o-terphenyl confined to nanometer-sized pores, as indi-cated. Bulk material is shown for comparison. It is clear that there are two glasslike transitionsin the confined o-TP, with one being lower than that of the bulk and one being higher. (AfterPark and McKenna [2000].)

symposium on dynamics in confinement [Koza et al., 2007] and is shown in Figure 5.5.The reason to show such a chart is that it is still perceived as relevant today becausethe Tg or dynamics continue to be studied extensively and the results are frequentlyapparently at odds with one another. In fact, the question that is frequently posed is:Does size matter? [Bansal et al., 2005; Mayes, 2005; McKenna, 2007]. Then, perhaps,another way to view the problem is to ask what we do know. There is evidence thatthere is both a finite size effect on the dynamics and a surface effect. An example ofsuch competing factors is given in Figure 5.6, where the Tg of toluene is first seen todecrease as size decreases (1/d) and then to increase as the surface interactions beginto dominate [Alba-Simionesco et al., 2003]. We know that surface interactions havean effect on the Tg or dynamics [Teichroeb and Forrest, 2003; Sharp et al., 2004,2007]. Other factors [Hutcheson and McKenna, 2005, 2007; McKenna, 2007] mayplay important roles, although this view is not shared universally, and further workneeds to be done to establish the relative importance of the various effects.

Two other issues that are important concerning the effect of confinement on theglass transition temperature and its associated dynamics arise from the interactionsof the confinement itself with the fluid or (polymer) of interest. The first effect issimply that of macroscopic confinement effects. Hence, if the confining mediumcauses a mechanical stress to be applied to the sample, there could be, for example, ahydrostatic tension applied to the fluid, and this would result in an apparent change in


FIGURE 5.5 Arrow chart used by R. Richert in presentation at the Third International Con-ference on Dynamics in Confinement, Grenoble, France, April 2006. (Originally reproducedby McKenna [2007].)

0.0

100 10 5d/nm

1/d/nm-1

3 2

100

120

140

160

180

0.1 0.2

finite size effects

surface effects

T/K

0.3 0.4 0.5

FIGURE 5.6 Glass transition temperature versus reciprocal pore diameter for toluene show-ing competition between finite-sized and surface effects. (From Alba-Simionesco et al. [2003],with permission. Copyright © 2003 European Physical Journal.)


the glass transition temperature. This possibility was already considered by Jacksonand McKenna [1991a,1996] in their early work and has also been discussed above inthe context of the block copolymers [Krause et al., 1982a, b; Lu and Krause, 1982;Lu et al., 1982; Granger et al., 1986; Wang and Krause, 1987] and, more recently,by the present author [McKenna, 2000]. It is clear that such macroscopic effectscannot fully account for the observed behaviors, although they may contribute. Thesecond important factor is the exact nature of the confinement as “hard” or “soft”,and this is an aspect of the confinement properties that has been emphasized byRichert and co-workers [Wang et al., 2004; He et al., 2005, 2007]. The importantpoint to be taken away from the work of Richert is that there does seem to be aninfluence of the actual “hardness” of the confinement on the dynamics (or the Tg )of the confined fluid of interest and how strongly the slowing due to this interactionis transmitted into the interior of the liquid. This is contrasted to the “soft” micro-or nanoemulsions [Wang et al., 2004; He et al., 2005] that were also mentionedabove [MacFarlane and Angell, 1982; Angell et al., 1984] and where the dynamicsof the confined fluid have been interpreted to take on the rapid dynamics of the softwall.

Needless to say, these results continue to be widely discussed. In this author’sview, it seems that the results we are dealing with are, for the vast majority of cases,properly obtained and not per se artifactual. Rather, we are missing a widely acceptedtheory or model that can account for the body of existing results and this remainsa challenge for the community that deals with theories of confined and nanoscaleliquids and glass-forming liquids.

5.3 ULTRATHIN POLYMER FILMS

As indicated in the introductory section, polymeric materials provide the possibilityof studying materials that have nanometer dimensions in the form of ultrathin filmsbecause these materials do not break up into droplets, as expected from surface–volume considerations. Rather, the samples are at least metastable as films becauseof the very high viscosities of polymers (due to the long-chain nature of the polymermolecule and consequent entanglements). Hence, above the glass transition, thesematerials remain intact and can be investigated. The first such studies came fromKeddie et al. [1994a,b] for films supported on substrates. These films were studied inwhat we refer to as a pseudothermodynamic mode in that the thickness was measuredas a function of temperature and the “break” in the curve was interpreted as theglass transition temperature, just as if a macroscopic sample had been measured in adilatometer. Figure 5.7 shows a compilation of results for similar experiments fromthe literature [Forrest and Dalnoki-Veress, 2001].

One can see from Figure 5.7 that depressions of the Tg of approximately 35 Khave been reported for the supported films. It is important to note, however, that inother systems, such as poly(methyl methacrylate) (PMMA), when the polymer andsubstrate have strong interactions, the Tg value increases. Figure 5.8 shows results forthe Tg of PMMA and PS films supported on substrates having different interaction

ULTRATHIN POLYMER FILMS 201

FIGURE 5.7 Summary of the Tg dependence on film thickness for films supported on oneside by a substrate. (From Forrest and Dalnoki-Veress [2001], with permission. Copyright ©

2001 Elsevier.)

strength. As shown, the interfacial energy affects the Tg dramatically, especially asthe films get thinner [Fryer et al., 2001]. Hence, it appears that the strength of thesurface interaction itself is important in the discussion of confinement effects on theglass transition temperature [Fryer et al., 2000, 2001; Grohens et al., 2000; Tsui et al.,2001; Sharp and Forrest, 2003].

A major advancement in the study of nanometer-thick films came when Dutcherand co-workers [Forrest et al., 1996, 1997, 1998] developed a capability to per-form ellipsometry or Brillouin scattering measurements on freely standing films. One

FIGURE 5.8 Change in Tg relative to the bulk value for thin films of PMMA and PS adheredto substrates having different surface energies. (From Fryer et al. [2000], with permission.Copyright © 2000 American Chemical Society.)


0

20

40

60

80

100

50 100h (nm)

Tg

(ºC

)

150 200

Mw = 767 k, ellipsometry

Mw = 2.24 M, ellipsometry

Mw = 767 k, BLS

Mw = 2.24 M, BLS

FIGURE 5.9 Glass temperature Tg versus film thickness for free-standing PS films with twodifferent molecular weights. The experiments were carried out using Brillouin light scatteringand ellipsometry. (From Dalnoki-Veress et al. [2001], with permission. Copyright © 2001American Physical Society).

reason that this turned out to be an important result was that they discovered thatfreely standing or unsupported PS films could exhibit much greater decreases inthe glass temperature than had been observed in the supported films to that time, andunlike the supported films, the freely standing films exhibited a very strong molecular-weight dependence of the glass temperature in addition to a very strong thicknessdependence. Figure 5.9 shows some of their results. The important point of thatwork was the finding that the Tg could be reduced by over 70 K, and that the reduc-tions, above some critical but still high molecular weight, began to exhibit a strongmolecular-weight dependence, which is not seen in the supported films. Hence, wesee that freely standing and one-side-supported films exhibit reductions in the Tg ,but we also see that the magnitude of the reductions is much greater for the freelystanding films than it is for the supported films (or films with one free surface).

In a way, the findings of such large changes in the Tg for the freely standingpolystyrene films are somewhat serendipitous. This is because there is not a broadlyaccepted theory of the effects of confinement or size on the glass transition, andthe polystyrene films were selected for study only because PS is a generally widelystudied polymer, not because anything special about the material’s behavior wasanticipated. Also, the PS freely standing films exhibit very large changes in theglass temperature (see Figure 5.9), and such changes are not seen in all polymers[O’Connell and McKenna, 2005; Roth and Dutcher, 2005a], and the nano-confinedsmall-molecule glassformers (see our earlier discussion) also do not show such largeeffects. O’Connell and McKenna [2005] used a novel bubble inflation method toexamine the creep and temperature response of poly(vinyl acetate) films and foundno effect of thickness on the glassy response down to thicknesses less than 27 nm.


FIGURE 5.10 Glass temperature versus film thickness for freely standing films of PMMA(circles) of M = 767 kg/mol and PS (triangles) of M = 790 kg/mol. (From Roth and Dutcher[2005a], with permission. Copyright © 2005 Elsevier.)

Subsequently O’Connell and McKenna [2006] also found that PS in similar exper-iments (i.e., creep and temperature dependence) showed reductions in Tg similar tothose found using the pseudothermodynamic methods of Dutcher’s group. Perhaps amore dramatic demonstration of how the Tg reduction in thin, freely standing filmsis nonuniversal (depends on the chemical structure of the polymer) comes from addi-tional work in Dutcher’s group, where the Tg reductions for PMMA are compared tothose found for a PS of similar molecular mass shown in Figure 5.10. Subsequently,Torkelson’s group [Torkelson, 2006] has shown that the Tg depression of supportedfilms depends on molecular structure and found for polycarbonate and polysulfonethat the Tg depressions can be even greater than found for PS, and depend on molec-ular structure in a series of n-methacrylates and in plasticized systems [Ellison et al.,2004; Priestley et al., 2007a].

Because the ultrathin films seem to show an extremely large reduction in the Tg

value, there has been considerable effort to understand the cause of the reductions.Many workers have suggested that the effect has to do with the free surface, just asone finds melting point depressions in small crystals [Jackson and McKenna, 1990,1991; Alba-Simionesco et al., 2006], and there is some evidence that this may be so.For example, Forrest and co-workers [Sharp et al., 2004] have considered the caseof capping polystyrene, hence having no free surface, and have sometimes foundno apparent reduction in the glass temperature. On the other hand, work by Kohet al. [2006] suggests that simply removing the free surface by stacking films isinsufficient to restore the macroscopic Tg . The Torkelson group’s work using fluo-rescent dyes to probe the structure of the films shows that the Tg exhibits a largegradient from the free surface to the supporting substrate [Ellison and Torkelson,2003]. Interestingly, that work suggests a 12 to 14-nm-thick layer at the surface hav-ing a reduced Tg . The thickness of the free surface by a nanosphere embedment


experiment by Teichroeb and Forrest [2003], on the other hand, suggests a thicknessof less than 3 to 4 nm. This result is in better agreement with positron annihilationlifetime experiments, where the density profile at the surface is different from thebulk only to a depth of approximately 2 to 5 nm [Algers et al., 2004a,b]. The longlength scales seen by Torkelson seem to require further exploration to fully understandthe result.

Although there is considerable evidence for reduction of the surface Tg or increasedsurface mobility [Agra et al., 2000; Schwab and Dhinojwala, 2003; Teichroeb andForrest, 2003; Erichson et al., 2004, 2007; Sharp et al., 2004; Fakhraai and Forrest,2008], some of the experiments have been challenged, while in others differenttypes of measurement suggest that the surface exhibits very close to bulk dynamics[Liu et al., 1997; Kerle et al., 2001; Schwab et al., 2003; Hutcheson and McKenna,2005, 2007]. It is clear that observations of a surface with mobility corresponding toa 70◦C reduction in the glass temperature have not yet been made directly. Rather,there is evidence of mobility increases that are more akin to a 5 to 10◦C drop in thesurface Tg . (The reader should keep in mind that a 10◦C drop in Tg would correspondto an increase in surface mobility of approximately three to four orders of magni-tude, taking the approximation that mobility changes by approximately one order ofmagnitude for each 3◦C change in Tg [Ferry, 1980].)

One other aspect of the apparent reductions reported for ultrathin films of polymersthat remains to be understood is why some experiments exhibit a strong reductionin the Tg , as just shown, whereas other, apparently similar experiments do not. In away, this is the equivalent to the arrow chart of Figure 5.5. A good example comesfrom a comparison of the author’s own work and the work of Bodiguel and Fretigny[2006a,b, 2007] both on polystyrene films of high molecular weight. In the author’swork [O’Connell and McKenna, 2005, 2006, 2007, 2008] we were able to place ultra-thin films on a template that was pierced with through-holes and then to inflate themembranes in a miniaturization of the classic bubble inflation experiment [Wineman,1976; Treloar, 1975]. In this case, pressure is placed across the membrane and theinflation is tracked as a function of time using an atomic force microscope (AFM) toimage the bubble changes. Figure 5.11a shows an array of bubbles, and Figure 5.11bshows the time-dependent bubble profile. This can be analyzed to obtain the creepcompliance as well as temperature shift factors from time–temperature superposi-tion. Without going into the analysis, this is an apparently classic experiment, andFigure 5.12 shows the estimated glass transition range for polystyrene films of varyingthicknesses. These results are consistent with those discussed above from the Dutchergroup (i.e., there are dramatic reductions in the Tg as film thickness decreases). How-ever, in their rather elegant set of experiments, Bodiguel and Fretigny [2006a,b, 2007]took freely standing films of polystyrene and floated them onto the liquid glycerol.Upon heating above the Tg the films shrink symmetrically, and from a considerationof the surface tension and film shrinkage, the creep compliance and viscosity can bemeasured. The results of these workers are not consistent with the observation of adramatically reduced Tg in the ultrathin films. In addition to these findings, whichare inconsistent with the reduction of Tg in thin films, there is a considerable amountof work in the investigation of hole growth in thin films [Xavier et al., 2004, 2005;


FIGURE 5.11 (a) AFM image of an array of inflated bubbles of a 65-nm PS film at 28 kPapressure and 100◦C. (b) Plot that depicts the bubble shape at different creep times. This is a70.1-nm-thick PS film, the pressure is 41 kPa, and the temperature is 80◦C. (After O’Connelland McKenna [2006].)

Roth and Dutcher, 2005b, 2006] which is inconsistent with a reduced glass transitiontemperature in ultrathin films. Holes are found to grow at measurable rates, generallyonly near and above the macroscopic glass temperature rather than the greatly reducedTg appropriate to the film thickness.

FIGURE 5.12 Change in the glass temperature as a function of film thickness for polystyreneand poly(vinyl acetate). The PS results are consistent, with there being a large reduction of Tg

for the thinnest films. For the PVAc, the results are consistent with there being no reduction inTg . Tg estimates come from the time–temperature superposition behavior of bubble inflationexperiments. (After O’Connell and McKenna [2008].)


Finally, it is worthwhile to here comment on one set of ideas concerning the reduc-tions of the glass transition temperature in ultrathin films as put forth by Sergheiet al. [2005]. In that work, dielectric spectroscopy, ac calorimetry, and capacitive acdilatometry are used to study the dynamics of confined films, and the authors considersamples degraded in air at temperatures from 140 to 200◦C. Infrared spectroscopy isused to look for chain scission. Most of the results are obtained at 180 or 200◦C. Theauthors claim that observation of a reduced glass transition temperature of approxi-mately 15◦C in the polystyrene studied is consistent with reductions in the Tg observedin freely standing films being due to chain scission and consequent reduction of themolecular weight and, therefore, the glass transition temperature. Although this resultmay be consistent with the 15◦C reduction observed in their studies, in fact, it is notconsistent with the studies on thin films. There are two reasons for this. First, forthe thinnest films, the Tg of polystyrene falls below 50◦C (323 K). Examination ofFigure 2.11 in the textbook by Cowie [1991] shows that the molecular weight ofPS required for a Tg of 323 K is approximately 30 mainchain carbon atoms, whichcorresponds to a PS molecular weight of approximately 1500 g/mol. Hence, to obtaina glass temperature as low as 50◦C, the sample would be virtually unmanageablebecause of the extremely low molecular weight. In addition, very few of the experi-ments run on ultrathin films for the purpose of investigating glass transition behavioras a function of film thickness ever see a temperature of 140◦C, let alone 180 or200◦C. Furthermore, if as was done in the work of the present author, one annealsthe samples at temperatures relative to the Tg of the thin film (i.e., annealing of a filmhaving a Tg of 50◦C would be at 60 or 70◦C), the temperatures are much too low togive thermal degradation of the magnitude required to decrease the molecular weightfrom nearly 106 g/mol to 1500 g/mol. Finally, the observation that there is a greaterreduction in the Tg of freely standing PS films for the highest molecular weightswould also be inconsistent with the proposed mechanism of chemical degradationsince, for a fixed rate of degradation, one would expect a smaller reduction in Tg fora higher-molecular-weight polymer.

5.4 AGING IN CONFINED SYSTEMS

5.4.1 Physical Aging Basics

Physical aging of glassy materials is the result of the fact that glasses are nonequilib-rium materials and, as such, their properties evolve spontaneously toward equilibriumin a process known as structural recovery. Structural recovery is considered the evolu-tion of thermodynamic state variables such as volume or enthalpy toward equilibriumsubsequent to an arbitrary thermal history. Physical aging is the response of otherproperties, such as the viscoelastic or dielectric responses, to the changing state param-eters. Figure 5.13 shows the important aspects of physical aging. Figure 5.13a showsa volume–temperature schematic for a down-jump history and illustrates that theglass so-formed is in a nonequilibrium state. The arrow indicates that the structure(volume) will evolve toward the equilibrium value. Figure 5.13c shows the typical

AGING IN CONFINED SYSTEMS 207

enth

alp

y o

r vo

lum

eδ

x 10

00

t-ti (h)

Ten

sile

cre

ep c

om

plia

nce

, 10-1

0 m2 /

N

temperatureCreep time t (sec)

(a) (c)

(b) (d)

Supercooled Liquid

Aging time te, days

Structural recovery

6

5

4

3

5.0

T0 = 40ºC

19.8ºC

22.424.9

27.5

30

3532.5

2.00119.5 ºC124.2 ºC126.2 ºC127.3 ºC129.7 ºC132.5 ºC135.0 ºC

1.50

1.75

1.25

1.00

0.75log

a te

0.50

0.25

0.00

-0.253 5 64

log (ageing time [s])

4.54.03.5

3.02.52.0

1.5

1.00.50.0

10-3 10-2 10-1 100 101 102

102 104 106

2%Glassy

Equilibrium line

αg or Cpg

Tg

Tf

Ta

αI or CpI 0.03 -0.1 -0.3 -1 -3 -10 -30 -100 -300 -1000

FIGURE 5.13 (a) Schematic of temperature down-jump in volume–temperature space, alsoshowing a definition of fictive temperature Tf . (b) Creep compliance for a PVC quenched from90◦C to 40◦C and aged for 1000 days. The arrow indicates the direction of shifting and time-aging time superposition at 1000 days. (From Struik [1978], with permission.) (c) Volumerecovery after the temperature jumps from 40◦C to the temperatures indicated. δ is the volumedeparture from equilibrium [δ = (vt − v0)/v0] and v0 and vt are the volumes at equilibriumand at time t, respectively]. (Plots in (a) and (c) from Zheng and McKenna [2003]; data forglucose in plot (c) from Kovacs [1963].) (d) Aging time shift factors versus aging time for apolycarbonate at different temperatures, as shown. (After O’Connell and McKenna [1999].)See the text for a discussion.

volume recovery response for a series of temperature down-jumps and that the timesfor equilibrium to be attained depend on the depth of the temperature jump below theglass transition temperature, in this case approximately 40◦C. The classical physicalaging picture from the Struik 1978 landmark work is shown in Figure 5.13b for apoly(vinyl chloride) (PVC), in which the system was quenched from above to some40◦C below the glass temperature. As can be seen, the creep response at increas-ing aging times after the temperature jump stiffens by shifting the response towardlonger times. This response is well characterized by time–aging time superposition.The aging time shift factors are defined as ate = τte /τte,ref . τte is the creep retarda-tion time at the current aging time te and τte,ref is the creep retardation time at the


reference aging time. Then the aging time (double logarithmic) shift rate µ is definedas µ = (d log ate )/(d log te ). These are important parameters in the investigation ofthe physical aging process, and it is generally found that 0 < µ < 1. Figure 5.13dshows typical behavior of the shift factors at different aging times from data similarto those of Figure 5.13b for a polycarbonate at different temperatures. We see therethat at short aging times (or at temperatures far below the Tg ) the aging follows apower-law behavior, whereas at temperatures near Tg the aging slows as the sampleapproaches equilibrium, as does the volume in Figure 5.13c. We also note that atT � Tg no aging occurs. Furthermore, if experiments could be performed at muchshorter aging times, the power law at short times would transition to a regime ofvery low slope; hence, the entire aging regime is sigmoidal as one goes from veryshort aging times toward the long times where equilibrium is achieved [McKenna,2003]. When thermal histories are more complicated, the aging behavior can lookmore complex [Struik, 1977, 1978, 1988; McKenna, 2003; Zheng and McKenna,2003; Zheng et al., 2004]. A good example is shown in Figure 5.14, where the vol-ume recovery in a two-step or memory experiment is shown to be nonmonotonic(Figure 5.14a) and the shift factors (Figure 5.14b) are also shown to be nonmono-tonic. However, to a very good approximation, the creep retardation times (or shiftfactors to a known reference) depend on the volume alone (although there is evidencethat there is more to the picture than simply volume [McKenna et al., 1995; Simon etal., 1997; McKenna, 2003; Simon and Bernazzani, 2006]. In this case, for memory andother experiments where the structural recovery behavior is different from that in the

specific volume

(a) (b)

7

6

5

4

3

2

1

76

5

4

3

2

1

1

0.967

0.9690

-1

-2

-31 10 100 1000

0.968

10time, te, elapsed at 85ºC, min time, te, elapsed at 85ºC, min

1000100

shift 30log acm3/g

FIGURE 5.14 (a) Volume as a function of elapsed time after the second step of two-steptemperature histories, (b) aging time shift factors corresponding to the changing specific volumeof part (a). (After Struik [1978], with permission.)


shift, 10log a-3

-2

-1

0

10.96700 0.96800 0.96900

v, cm3/g

FIGURE 5.15 Shift factors versus specific volume for a polymer glass in different tempera-ture histories as indicated. (After Struik [1978], with permission).

simple down-jump instance, the responses still depend on the state parameters (i.e.,volume or enthalpy). This is shown in Figure 5.15, where the shift factors for creepexperiments at different aging times and for different thermal histories (indicated inthe plot) are shown to depend on volume alone. Hence, the down-jump temperaturehistory is an excellent surrogate for the full structural recovery behavior, at least to afirst approximation, and will be the history for results discussed subsequently.

The discussion above is a brief overview of structural recovery and physical aging,with emphasis on volume as the measure of structure and the mechanical response tothe structure being well explained through free-volume concepts. However, althoughmuch of the richness of the structural recovery and physical aging responses is cap-tured with these sorts of measurements, the reader needs to be aware that other methodsof measurement have also been used to investigate structural recovery and physi-cal aging. Calorimetry, rather than volume, is often used as the means of structurecharacterization (enthalpy), and dielectric spectroscopy has also been used as a toolfor the investigation of aging phenomena in glassy materials. We do not providedetails of these measurements here simply for space reasons. Next we describe whatis known about aging at the nanoscale.

5.4.2 Physical Aging and Structural Recovery at the Nanoscale

There are only a handful of investigations of aging of materials at the nanome-ter size scale. The earliest was done by the author in collaboration with workersat Eastman Kodak in a structural recovery investigation of o-TP confined in con-trolled pore glasses having a variety of pore sizes [McKenna et al., 1992]. Thiswork is of interest because of the errors of interpretation made in that study. The


Temperature (K)

Unaged

3 min

10 min

30 min

100 min

Unaged

3 min

10 min

30 min

100 min

Temperature (K)

Norm

ali

zed

Hea

t C

ap

aci

ty, C

pN

Norm

ali

zed

Hea

t C

ap

aci

ty, C

pN

210 220 230 240

(a) (b)

250 260 270 280 210 220 230 240 250 260 270 280

FIGURE 5.16 Enthalpy recovery results for o-terphenyl aged at Tg − 11◦C in (a) bulk stateand (b) confined in an 11.6-nm pore diameter controlled pore glass material showing muchsmaller buildup of enthalpy overshoot upon aging of the confined material. Originally, similardata were interpreted [McKenna et al. [1992] to imply reduced aging in confined systems.(Data from Simon et al. [2002].)

work treated the structural recovery by examining the enthalpy overshoot responsein DSC experiments after different aging times and within the framework of the Tool[1946a,b]–Narayanaswamy 1971–Moynihan et al. [1976] (TNM) model of structuralrecovery. Figure 5.16 shows the results of the recovery of o-TP compared to the bulkbehavior from a later study from NIST and Texas Tech [Simon et al., 2002] that showsthe same features.

The important feature observed in Figure 5.16 is that the enthalpy overshoot inmaterial confined to the nanoscale is greatly reduced. In the first work [McKennaet al., 1992], the TNM model fits to the data were interpreted to mean that the reducedenthalpy overshoot and its buildup with aging time corresponded to increased relax-ation times; that is, despite a reduced Tg at the nanoscale, the structural recovery ata constant temperature below the reduced Tg seemed to imply that the material wasless mobile than at the macroscale. Although possible, this did not seem very satis-fying, and that work appeared only in abstract form. In the later work, Simon et al.[2002] recognized an additional feature of the enthalpy recovery of the o-TP at thenanometer size scale. This is shown in Figure 5.17, where the difference between thefictive temperature Tf (a measure of the structure of the glass)†and the aging temper-ature Ta is shown as a function of aging time. As shown in the figure, the differenceTf − Ta does not go to zero but remains finite. Simon et al. had the insight to see thatthe fact that Tf − Ta does not go to zero implies that the confined material is in a

†The fictive temperature gives a measure of the structure of the glass that is frozen-in at the glass tem-perature. It is defined by taking a point in the nonequilibrium glassy state and drawing a line parallel tothe glassy line (this can be volume or enthalpy); the point of intersection with the extrapolated liquid-state(equilibrium) line is called the fictive temperature, Tf . See Figure 5.13a for the graphic depiction of thefictive temperature. It is related to the departure from equilibrium, which is a slightly different measure ofthe structure of the glassy state.


10

8

6

4

2

0101 102 103

Time (s)

Bulk

47.9 nm

25.5 nm

d = 11.6 nmT f

- T a

(K)

Ta = Tg - 8ºC

104 105

FIGURE 5.17 Enthalpy recovery as Tf − Ta for o-terphenyl in bulk and confined in nanoporesas indicated. Test temperatures are at Tg − 8◦C. Importantly, Tf − Ta does not go to zero forthe confined materials. See the text. (After Simon et al. [2002].)

different state than that of the bulk material. The kinetics of the structural recovery, inthis case, could be modeled successfully using an extension of the TNM and KAHR[Kovacs et al., 1979] models to a case in which the material in the nanopores is agedisochorically (constant volume) and with a Tg that is reduced from that of the bulkmaterial due to the nanoconfinement. Importantly, the modeling gave a quantitativeexplanation for the reduced enthalpy overshoots seen in Figure 5.16a and also pro-vided a clear view that the behavior in the confined pores seemed to age in a way thatwas related only to the reduced glass transition and the increased glassy enthalpy dueto the pore confinement to an isochoric state. The interpretation of a retarded processmade originally was incorrect, and from Figure 5.17 one could even argue that thesample equilibrates somewhat faster in the pores simply because the material evolvesto a state other than the bulk state. In the confined system, it is in fact a state of higherenthalpy and higher free volume. Of course, if this were the final answer, we couldimagine that the behavior of aging at the nanoscale is resolved. Unfortunately, forultrathin polymer films, and even just for very thin polymer films, this does not seemto be the case.

The first aging experiments in ultrathin polymer films were performed by Kawanaand Jones [2003] on polystyrene on a substrate. Thickness measurements by ellip-sometry were used as an elongational dilatometer, and volume overshoots akin to theenthalpy overshoots in Figure 5.16 were observed. Similar to the results from enthalpymeasurements, the overshoots were observed to decrease with decreasing film thick-ness, although the results were interpreted to be due to a gradient of properties frombulk to the liquid surface layer rather than as being due to an isochoric transition, aswas the case for o-TP confined to nanopores. Also, the work by Kawana and Jones[2003] did not follow the kinetics of aging, as these authors were more interested in


the film thickness effect and, in fact, observed no aging or glass transition in 10-nm-thick films. Of greater interest, perhaps, is a recent series of papers from Torkelson’sgroup in which fluorescent probes have been used not only to probe glass transitiongradients in ultrathin films supported on substrates, but also to investigate the phys-ical aging response in these thin films [Ellison and Torkelson, 2003; Priestley et al.,2005a,b; Mundra et al., 2006].

The work has also been complemented with some dielectric spectroscopy [Priestleyet al., 2007b, c]. The first thing to comment on is that in the bulk, the fluorescent probeintensity is a function of the density of the host matrix at a constant temperature.Therefore, making measurements of intensity versus time is nominally equivalentto making measurements of specific volume or density versus time (i.e., these arestructural recovery measurements). One of the major findings of the work is thatconfinement of the ultrathin film on a substrate leads to reduced structural recoveryin the ultrathin films, and the amount of recovery decreases as the film thicknessdecreases. Furthermore, the work from the Torkelson group suggests that the agingdepends on where in the thin film one places the probes [Priestley et al., 2005a,b], andthis is interpreted to be consistent with a gradient in Tg values reported previouslyby the group [Ellison and Torkelson, 2003; Mundra et al., 2006]. Typical resultsare shown in Figure 5.18, where we see the effects of confinement and positionon evolution of the normalized fluorescence intensity for poly(methyl methacrylate)layers at different locations relative to the supporting substrate.

The other aging work relevant to the present discussion comes from work initi-ated in the studies of Pfromm and Koros [1995] on polymer membranes in whichthe film dimensions are not as small as those in the aging studies above. This workhas been greatly enlarged upon in continuing work both from Paul and co-workers

FIGURE 5.18 Structural recovery at different temperatures as measured by fluorescenceintensity of probes in PMMA layers located at different distances from the substrate or freesurface, showing that the aging depends on position: (a) middle layer exhibiting bulk-likebehavior: (b) surface layer showing reduced Tg value as evidenced by no aging at a temperaturebelow the bulk Tg but still exhibiting reduced aging at the lower temperatures; (c) substratelayer showing the impact of constraining substrate to reduce aging at all temperatures. (circles),T = 32◦C; (triangles), T = 75◦C; (squares), T = 115◦C. Tg,bulk = 120◦C. (From Priestley et al.[2005a], with permission. Copyright © 2005 American Association for the Advancement ofScience.)


PPO

++

+

++

++

++

25.2 µmℓ

4.98 µm2.01 µm0.98 µm0.73 µm0.40 µm

“Bulk”

N2

Per

mea

bilit

y (b

arre

r)

2

3

4

5

6

0.1 11

10010 1000 10000

Aging Time (hr)

+

FIGURE 5.19 Evolution of the nitrogen permeability with aging time for a poly(phenyleneoxide) film at 35◦C as a function of film thickness. The Tg of the material is 210◦C. (FromHuang and Paul [2007b], with permission. Copyright © 2007 American Chemical Society.)

[McCaig et al., 2000; Huang and Paul, 2005, 2007a,b] and from Pfromm’s group[Dorkenoo and Pfromm, 1999, 2000]. Figure 5.19 shows the permeability of nitrogenin poly(phenylene oxide) films between 400 nm and 25 �m thick. It can be seen that thethinnest films show more rapid aging than the thickest films and that all films exhibitdecreases in permeability below that for the bulk at longer aging times. Importantly,unlike the experimental results on confined films reported above, aging is acceleratedrather than retarded. Such results are not totally incongruous, but it is surprising thatthe free surfaces in Torkelson group’s experiments (Figure 5.18b) do not show theaccelerated aging observed in the permeability of the thin membranes. In fact, theresults from Huang and Paul [2005,2007a,b] are reminiscent of results from Swallenet al. [2007] in which vapor-deposited glasses are found to exhibit unusually high den-sity and stability that has been ascribed to enhanced mobility of the surfaces as thedeposition takes place. Interestingly, in earlier work McCaig et al. [2000] had been ableto describe the aging of their microscopic films nearer the Tg using a combined free-volume recovery model and a free-volume diffusion model. Although such modelingis promising, there is also evidence from dilatometry of micrometer-sized particlesfrom Braun and Kovacs [1963] which suggests that free volume does not diffuse, andthis has been an issue over the years regarding free-volume descriptions of the glasstransition. In any event, the results of Figure 5.19 were not so-modeled, undoubtedlybecause the aging here takes place far below the glass transition temperature, and


the types of models that form the basis of our understanding of the aging phenom-ena breakdown under such conditions. As suggested by these authors, further workis clearly required. Another item of interest here is that the observed reduction inpermeability to below that of the bulk is also reminiscent of the implosion or densifi-cation event observed in mechanical straining of glassy polymers far below their Tg

values [Colucci et al., 1997]. The similarities may suggest that there are unaccountedfor residual stresses in the thin films or that the diffusion–permeability process itselfinduces stresses sufficient to cause implosion. The changed aging of glasses confinedat the nanoscale is also seen in polymer nanocomposites [Priestley et al., 2007c]. Infact, thin films have been taken to be good models for the confined polymer in thenanocomposites, where the distances between particles can be nanometric, but areaverages and because of the particle shapes, often are curved. Furthermore, poly-mers in the semicrystalline state also have nanoconfined amorphous regions that canbe either rigid amorphous [Menczel and Wunderlich, 1981] or simply confined orconstrained [Struik, 1987a,b, 1989a,b; Wunderlich, 1994; Aharoni, 1998], dependingon the view that one takes. Aging of such confined systems has been little studied,although some thoughtful results were found early with measurements of aging ofthe rigid amorphous phase [Menczel and Wunderlich, 1981; Huo and Cebe, 1992;Wunderlich, 1994; Krishnaswamy et al., 2003] in semicrystalline polymers and sim-ply the aging of semicrystalline polymers in the broadened glass transition regime ofthe constrained amorphous phase [Struik, 1987a,b, 1989a,b; Beckmann et al., 1997].More recent work on nanocomposites has suggested that aging can be suppressed bymaking the nanoparticles attractive [Priestley et al., 2007c], which can also increasethe glass transition temperature. Yet a full understanding of the physical aging processat the nanometer size scale eludes us. One possible reason for this is that the behaviorof both the glass transition and of aging have been studied primarily using confinedsystems, such as supported films, nanoparticle-reinforced resins, or fluids imbibedinto rigid supports. To the best of the author’s knowledge, there are no measurementsof the viscoelastic properties during aging of ultrathin, unsupported polymer films ofnanometer thickness.

5.5 SUMMARY AND PERSPECTIVES

It is clear that material behavior at the nanometer size scale is not completely under-stood. In the case of the glass transition behavior, there are significant debates in thecommunity, in which it is argued whether or not reduction in the glass temperaturewhen a material is confined to extremely small dimensions (e.g., in pores, as films)is due entirely to surface effects or if there is a combination of finite size effects andsurface effects. Furthermore, it seems that macroscopic effects may be important ifthe confining medium (e.g., substrate, pore wall) can lead to mechanical stress orpressure being applied to the fluid of interest. Furthermore, there is evidence thatthe mechanical nature of the confinement (hard versus soft) may also play a role inhow the dynamics of the glass-forming liquids change at the nanoscale. There is littledoubt that there is a finite size effect that is complicated by the presence of large

SUMMARY AND PERSPECTIVES 215

surface contributions to the behaviors observed. In this chapter I have dealt primarilywith the experimental situation because I view this as the area not only where themost work has been done but also because currently, the field seems to be driven byexperiment. In my view the experimental situation is unsettled because of the subtletyof the issues at hand and not, as some might have us believe, that the experimentsthemselves are full of artifacts. As I noted in a recent paper [McKenna, 2007], it isimportant that the community recognize the range of results and not only the resultsfor which the individual researcher or research team is responsible. It is important torecognize that most of the results, although often apparently conflicting, are obtainedfrom well-executed experiments. Hence, the challenge of determining the molecularor microstructural bases of the observations may well depend on the developmentof a widely accepted theory of the glass transition that goes beyond those classicaltheories, such as the Simha–Somcynsky [1969] free volume–based cell model, theoften used Doolittle [Doolittle, 1951; Doolittle and Doolittle, 1957]–Ferry [Williamset al., 1955; Ferry, 1980] free-volume model, and the Gibbs–Dimarzio [DiMarzio andGibbs, 1958; Gibbs and DiMarzio, 1958] configurational entropy model. Whether it isan adaptation of these to the nanoscale [McCoy and Curro, 2002], or novel approachessuch as percolation models [Hunt, 1994; Long and Lequeux, 2001; Baljon et al., 2004],detailed cooperativity length-based models [Sappelt and Jackle, 1993], or energy land-scape models [Truskett and Ganesan, 2003; Mittal et al., 2004], is one challenge forthe future. At the same time, it may be that novel experiments (and some of thosediscussed above pushed to different limits) will clarify how the dynamics of glass-forming liquids change with size and confinement. Clearly, the behavior of materialsat this intermediate “nanometer” size scale between the atomistic angstrom scale andthe macroscopic scale that we think begins in the submicrometer range remains asubject of interest and importance and will remain so for many years to come.

Acknowledgment and Dedication

Upon accepting the invitation to provide an article to this book dedicated to RobertSimha, I had to think very hard about how to develop the manuscript, particularlybecause my own work, while influenced by the work of Simha over the years, hasnot been a direct offshoot, nor has it been collaborative in nature. Hence, I decidedthat this paper is dedicated to Robert Simha in the sense of deepest respect for thewide range of contributions that he has made over a career that approximates threequarters of a century. I also feel it is justified to point out that when great men suchas Robert Simha interact with those who are their juniors, they frequently show thatthey are more than the simple brilliant intellectual. Robert Simha has always beensomeone who in my experience treated those junior to him with the utmost respectand consideration. Thus, he encouraged those of us who followed his broad path. Itis with this sense that I dedicate this work to Robert Simha, in the hope that the workpresented here merits such a dedication (Paris, 2007, 2008).

The author would like also to thank the E.S.P.C.I. in Paris for hosting his visitduring much of the preparation of this work. The National Science Foundation is alsoacknowledged for partial support under grant OMR-0804438.


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Polymer Physics (From Suspensions to Nanocomposites and Beyond) || Dynamics of Materials at the Nanoscale: Small-Molecule Liquids and Polymer Films