ARTICLE pubs.acs.org/JPCB
Nanomechanical Properties of Advanced Plasma Polymerized Coatings for Mechanical Data Storage Davide Tranchida,†,‡,||,* Sascha A. Pihan,† Yi Zhang,† Holger Sch€onherr,‡ and R€udiger Berger†,* † ‡
Max Planck Institute for Polymer Research, Ackermannweg 10, 55128 Mainz, Germany Physical Chemistry I, University of Siegen, Science and Technology, Adolf-Reichwein-Str. 2, 57076 Siegen, Germany
bS Supporting Information ABSTRACT: In this paper we report on the unprecedented deformation behavior of stratified ultrathin polymer films. The mechanical behavior of layered nanoscale films composed of 8-12 nm thin plasma polymerized hexamethyldisiloxane (ppHMDSO) films on a 70 nm thick film of polystyrene was unveiled by atomic force microscopy nanoindentation. In particular, we observed transitions from the deformation of a thin plate under point load to an elastic contact of a paraboloid of revolution, followed by an elastic-plastic contact for polystyrene and finally an elastic contact for silicon. The different deformation modes were identified on the basis of force-penetration data and atomic force microscopy images of residual indents. A clear threshold was observed for the onset of plastic deformation of the films at loads larger than 2 μN. The measured force curves are in agreement with an elastic and elastic-plastic contact mechanics model, taking the amount of deformation and the geometry of the layer that presumably contributed more to the overall deformation into account. This study shows that the complex deformation behavior of advanced soft matter systems with nanoscale dimensions can be successfully unraveled.
L
ayered materials and architectures are frequently encountered in biological systems and provide optimized functionality. As can be recognized for instance in teeth or cartilage,1 the nature and strength of the individual constituent layers but also their mutual interaction play a fundamental role in determining the ultimate performance.2 Layered systems are also widely used in technology, for example, to combine very diverse properties of different materials into systems with unique performance. Examples include functional thin film materials for optical,3 electronic,4 or dielectric5 applications, which are protected by a hard overcoat that protects the underlying layers from the interaction with the ambient.6 Additionally, interlayers7 may be applied to improve the adhesion between poorly compatible layers. Microsystems technology, in particular, relies on the miniaturization of systems and on the fabrication of complex nanoscale structures. In advanced applications, such as mechanical data storage,8,9 quantitative information on the nanomechanical behavior of the layered systems is crucial. The behavior of layered systems is more complex and goes beyond the mere knowledge of the conventional materials’ properties such as Young’s moduli and the yield stresses of each component. In particular, the typical loads and penetration depths required to trigger certain local deformation mechanisms must be known in order to be able to optimize the film structure and architecture.6 Thermomechanical data storage relies on embossing the information as indentation marks into a polymer film.8 Hereby the indentation depth and diameter depend on the tip/sample temperature and/or the pressure to the film.9 The information r 2011 American Chemical Society
written with, e.g., an array of thermal atomic force microscopy (AFM) probes can also be erased to provide a write-read data storage device.10 While the nanoscale determination of mechanical properties using AFM nanoindentation has been established in the past decade,11-14 stratified nanoscale systems pose substantial challenges. During nanoindentation, the stress field extends over a larger volume compared to the volume directly deforming under the indenter. This extension of the stress field has hampered the measurement of film-only mechanical properties.15 A number of previous reports have exploited the surface forces apparatus (SFA),16 finite element analyses (FEA),17 and conventional nanoindentation to identify the peculiarities of layered systems18-20 and to unravel the effect of the mechanical properties of the individual components on the overall deformation behavior. The theoretical analysis of layered systems is often based on integral transforms; i.e., exact expressions for the Hankel transforms of displacement and stress were evaluated by imposing boundary conditions at the corresponding interfaces of the layers.21 It was found that the thickness of the sample,22 together with the corresponding mechanical properties of the components, plays a significant role.23 The stress field differs significantly depending mainly on the Young’s moduli and yield stresses of the components. The ratio of the corresponding mechanical properties determine whether the stress field extends Received: December 23, 2010 Published: March 16, 2011 3385
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Scheme 1. Schematic of the Layered Film System Investigated in This Study and AFM Nanoindentation (Not to Scale)
some distance into the film or all the way through to the substrate.24,25 Moreover, significant confinement may occur, when the ratio of the contact radius and the film thickness approaches or overcomes unity.26 The combination of high hydrostatic pressure, size scales, and geometrical confinement implies that the film response is essentially oedometric in nature and that the mechanical properties of polymer coatings can be largely different from their bulk values. In this study we aimed to unravel the deformation mechanisms of a stratified nanoscale polymer film composed of ultrathin (8-12 nm thin) plasma polymerized hexamethyldisiloxane (ppHMDSO)27-29 on a 70 nm thick polystyrene (PS) film on silicon performing AFM nanoindentation experiments. The deformation of the ultrathin hard layer, followed by the elastic deformation of the thin PS layer and its elastic-plastic deformation at higher loads were fully characterized and modeled using appropriate contact mechanics models. Our results provide insight into the deformation behavior of a model nanoscale multilayered polymer system and can be generalized to memory storage systems as well as, for instance, glass-polymer multilayer composites, which provide wear resistance and toughness, respectively.
’ RESULTS AND DISCUSSION The mechanical properties of the layered nanoscale polymer film obtained by plasma polymerization of hexamethydisiloxane (HMDSO) on a 70 nm thick film of polystyrene on silicon were investigated by AFM-based nanoindentation (Scheme 1). In addition to the acquisition of force-displacement data (which are converted to force-penetration-depth data), intermittent contact mode AFM imaging was applied to determine the shape and the depth of the residual indents. In AFM nanoindentation the sample is vertically displaced in the surface normal (z) direction at a fixed x,y position, and the deflection of the cantilever (the force exerted) is measured. The force-displacement curve contains quantitative information about the deformation of the sample under load.30 The subsequently recorded topography image of the residual imprint provides information on the elasticity of the contact, i.e., the ability of the material to recover from the deformation imposed during the indentation,31 and helps to identify the appropriate contact mechanics model via the shape of the imaged indent and the load required to trigger plastic behavior. Finally, the amount of material bulging out of the surface, the so-called pileup, provides insight into the plastic flow upon an indentation
Figure 1. (A-C) AFM images of the residual imprints after nanoindentation on ppHMDSO (8 nm)/PS (70 nm) with a maximum load of 0.73 (A), 1.49 (B), and 2.54 μN (C), showing a significant difference in pile-up depending on the maximum load applied to the sample. The scale bar represents 20 nm for all three images, while the grey scale covers 3.4 (A), 5.5 (B), and 13.2 nm (C), respectively. The corresponding cross-sectional height profiles for A-C are shown in the Supporting Information. (D) The applied load vs depth curves show good reproducibility. The residual indentation depth, as measured by AFM (open circles), are also plotted, pointing to the prevailing elastic nature of the contact.
experiment and thus on the volume, in which the stress field exceeded the yield stress. Parts A-C of Figure 1 show AFM height data of the residual indents observed after indentation with a maximum load of 0.73, 1.49, and 2.65 μN, respectively, and a plot of applied load as a function of depth. Figure 1D shows the comparison of penetration depth observed during the acquisition of 20 force curves (filled symbols) at different loads at random places and the residual indentation depths (open symbols). An outstanding reproducibility of the force curves at low penetration depth was observed. Minor differences among the force curves can be noticed only at high penetration depth (>65 nm), which are attributed to slightly different PS thickness after spin coating. The films investigated were very homogeneous and possessed a root-mean-square roughness smaller than 0.3 nm (assessed on an image size of 5 μm 5 μm). While the residual indent in Figure 1A is barely noticed, a small pileup is observed in Figure 1B, as well as in Figure 1C, pointing at plastic flow. A comparison of the residual indents of the 8 nm thin ppHMDSO/PS and the 12 nm thin ppHMDSO/PS is shown in the Supporting Information. The asymmetry of the pileup, discernible in both images, but more significant in Figure 1C, is attributed to local stress concentration, in accordance with previous reports in the literature.32,33 This stress concentration is caused by the geometry of the tip, as will be briefly outlined. The profile of the AFM tip that was used can be modeled as a paraboloid for imaging purposes. Thus, tip 3386
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The Journal of Physical Chemistry B convolution can be successfully modeled in some instances. In addition, the modeling of the tip shape as paraboloid is useful for the analysis of force curves at moderate penetration depths. Although this simplification allows one to correctly evaluate the Young’s moduli of polymers from force curves via the use of contact mechanics models,34 it does not take the real shape of the tip into account. The tip shape is especially relevant at the point, where the pyramidal shape of the tip converges to the rounding of the apex (see the Supporting Information). Any discontinuity, or a local spatially rapid change, of the derivative of the indenter profile gives rise to a noticeable stress concentration. Such a local maximum of the pressure profile at the surface leads to rupture of the HMDSO film and plastic flow. This interpretation is confirmed by the observation that the pileup is consistently more significant on the left side of the residual indents.35 In the force curves four regimes are indentified: (i) A linear region, which extends to roughly 15 nm penetration depth, related to the deformation of the hard ppHMDSO layer as a bending plate36 (this deformation turns out to be completely elastic because there is no residual indentation, at least on the scale that the AFM can image); (ii) a second region, where the slope gradually increases between 15 and 50 nm penetration depth, related to the elastic deformation of the PS film, obviously together with the top hard layer; followed by (iii) a third region until ca. 70 nm penetration depth where the contact with PS is essentially plastic; and finally (iv) a very steep increase of the load needed to induce a very small increase of penetration depth, when the tip touches the silicon substrate. Details on the fitting procedure can be found in the Supporting Information. From the data shown in Figure 1A-C it can be concluded that the nanoindentation took place mostly in the elastic regime, since the residual indents were always roughly 10 times smaller than the maximum penetration depth. This finding is surprising, since even when the residual depth amounts to a few nanometers, the penetration depth under load was as large as half of the entire film thickness. Chen and Vlassak37 showed by means of finite element simulation that in the case of indentations of a soft film deposited on a hard material, a high penetration depth does not imply substantial contributions from the hard substrate to the stress field. Neglecting for the time being the ultrathin ppHMDSO layer, our results complement this view, showing that the deformation of a soft film on top of a hard substrate is elastic under confinement. Indeed, the material (PS) which normally would undergo plastic deformation is likely strengthened in our case by synergetic effects from both the size scale of the indentations (postponing the yield38) and the confined geometry (imparting superior mechanical properties) between the two hard layers, i.e., ppHMDSO and silicon. Plastic deformation of polystyrene normally starts with cavitation that boosts further crazing and shear yielding.39 However, positive pressure components arise inherently due to the nanoindentation geometry and the geometrical confinement of the deformation.40 These thwart the formation of cavities and therefore a situation for maintaining the strength of the material is induced. Yield takes place at some point even without cavitation, due to strain-induced softening, which lowers the glass transition temperature of the material involved by the stress field and then allows segmental motion of the chain resulting in macroscopic yielding.41 However, the substantial level of confinement due to the similarity of penetration depth and film thickness causes a high level of hydrostatic pressure that further increase the yield stress.26
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Figure 2. Experimental force curves obtained on a 8 nm thin ppHMDSO layer on PS (the symbols denote the measured data points) and modeling according to appropriate contact mechanics theories: ppHMDSO elastic bending plate (blue curve, the linear fit of the first part of the force curve extends to the penetration depth at which the deviation of the fit exceeded 10% with respect to the experimental force data),37 elastic deformation of PS (green curve, the transition to the next deformation mode was placed where the deviation exceeded 5% of the experimental curve; see Supporting Information for details),34 elasticplastic deformation of PS (red curve),35 and elastic deformation of the silicon wafer (brown curve)34 fits to the corresponding sections of the force curve. The inset shows the corresponding data and fits for a 12 nm thin ppHMDSO layer on PS.
In Figure 1 it can further be seen that even when the penetration depth under full load reaches 40 nm, the residual depth amounts to only 4 nm. This value of the residual depth corresponds to 30% of the total thickness of the ppHMDSO layer. Moreover, the AFM images of the residual imprint do not show any permanent bending of the ppHMDSO layer. This observation implies that plasticity of PS or rupture of the ppHMDSO layer are reached within a very small area where the stress is high and subsequently quickly drops below the yield threshold. It is also well-known that, under confined conditions, stresses do not expand significantly outside the contact zone.42 In other words, the stress applied to the surface of the layer in the surface normal direction should be integrally transmitted to the substrate over a constant area, where failure takes place at a certain point. When the load increases further, the PS starts to deform plastically and flows as shown in Figure 1C. To model these phenomenological findings, the contact mechanics belonging to each region was plotted and afterward shifted in the same figure to fit the experimental curves. As shown in Figure 2, in which the experimental data (symbols) and the fits according to contact mechanics (solid lines) are superimposed, the deformation of a thin plate under point load, an elastic contact of a paraboloid of revolution and an elastic-plastic contact for PS, and finally an elastic contact for silicon can be differentiated. The first part of the force curve (blue line) is linear. The main examples of a linear relationship between applied load and penetration depth found in the contact mechanics literature belong to the indentation of a flat punch into an elastic halfspace and to the deformation of a thin plate under a point load. The latter situation is a reasonable interpretation for the 3387
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Table 1. Penetration Depth and Applied Load Required To Trigger the Different Contact Mechanics Behavior for Both the 8 and 12 nm Thick ppHMDSO Films applied load, μN
penetration depth, nm ppHMDSO film thickness, nm
PS elastic
PS elastic plastic
silicon substrate
PS elastic
PS elastic plastic
silicon substrate
8
17 ( 1
38 ( 1
72 ( 2
1.1 ( 0.1
2.0 ( 0.1
3.9 ( 0.7
12
21 ( 1
53 ( 1
80 ( 2
1.2 ( 0.1
2.4 ( 0.1
4.1 ( 0.7
deformation of the ppHMDSO layer. The first linear part of the force curve therefore belongs to the deformation of the 8 nm thin ppHMDSO plate. The aspect ratio of this plate (radius/ thickness) is not known a priori; therefore the equation governing this phenomenon can be used to fit an “equivalent radius” of the plate. The equivalent radius was found to be ∼26 nm for the sample with 8 nm thin ppHMDSO and ∼50 nm for the 12 nm thin ppHMDSO sample. The green and the brown lines refer to Sneddon’s model43 for elastic contact of a paraboloid of revolution in the case of PS and silicon, respectively, while the red line represents the elasticplastic contact for PS, after Hays and Kendall.44 Obviously these models need some input parameters with respect to materials properties as well. We assumed for the Young’s modulus values from the literature obtained in other geometries (Young’s modulus of 2 GPa and yield stress of 40 MPa for PS).45 The continuity between the elastic-plastic (PS) and elastic (Si) regime is obvious, since, once the tip touches the hard silicon substrate, all of the energy spent is used to deform this hard substrate. The continuity between the elastic and elastic-plastic regimes for PS is also discernible. The only problematic point is the first transition, which suggests that initially only the ultrathin ppHMDSO layer is deformed, while the PS is stress-free. Subsequently, above a certain threshold of the load, the deformation of the PS takes control of the contact mechanics. The real phenomenology should not be represented by such a dichotomization. It can be expected that the two processes are competing. At low deformation, the deformation of the ultrathin hard layer on the more compliant thin film takes place in a bending mode. This bending introduces compressive stresses that decrease the maximum value of the principal shear stress, thus delaying the yield of the overall system. At large enough deformation, the contact pressure distribution changes, from being concentrated in one single contact point to being distributed along the contact edges.46 Moreover, contact stresses increase at the same time, and the ultrathin film ruptures. Therefore the loading condition on the PS film becomes more similar to a standard indentation, giving rise to the part of the force curve after point A. Although simplified, this analysis allowed us to perfectly fit the experimental force curves. The comparison of the data for the 8 nm thin with the 12 nm thin ppHMDSO layer is instructive, in order to check the influence of the hard layer thickness on the overall mechanical properties. The same modeling procedure described above was applied in this case, as shown in the inset of Figure 2. The loads at which transitions among the different contact regimes were observed for both films are summarized in Table 1. As expected, the transitions for the thicker ppHMDSO layer occur at slightly higher values of applied load, the values for the penetration depth are also in agreement with the model described above. In addition to the qualitative description of the underlying contact mechanics, as shown in Figure 1, and the corresponding
Figure 3. Evaluation of Young’s modulus from (A) the loading curve and (B) the unloading curves (where, for example, 5.Eþ10 represents 5 1010). A consistent trend is found, although the absolute values obtained from the analysis of the unloading curves in B are too high, as discussed in the text. (C) The ratio of unloading and loading slopes at maximum load highlights the occurrence of likely viscoelastic phenomena. The substrate effect starts to be discernible after roughly 55 nm penetration depth.
modeling, as shown in Figure 2, it is interesting to evaluate the mechanical properties of each component of the layered system using well-known contact mechanics models.43 Figure 3A shows the variation of the elastic modulus of the layered film (for the 8 nm thin ppHMDSO sample) with penetration depth upon loading as measured continuously on the loading curve. A high value is obtained in the very first nanometers after establishing contact, which corresponds to the deformation of the 3388
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The Journal of Physical Chemistry B ppHMDSO before it acts like a membrane. This value matches previously reported data for plasma polymerized ppHMDSO of Beake et al., who measured elastic modulus in the range of 1020 GPa.47 Afterward, the Young’s modulus levels off to a plateau that is attributed to PS. This finding is similar to a report on several glassy polymer films with thickness of approximately 1 μm, which was explained as enhanced stiffness of an amorphous polymer interphase.48 On the other hand, the decrease of the Young’s modulus with increasing penetration depth was observed in the case of thick films49 and was interpreted as an artifact due to destabilizing attractive force gradients in the vicinity of surfaces. Until here only the loading curves were discussed (Figure 1 and Figure 2). In general, unloading curves acquired on polymers often do not bear much information on the sample material.50 The reasons for limitations in the interpretation of AFM nanoindentation unloading curves of polymers originate from the onset of viscoelastic behavior.40 This onset manifests itself in a high exponent in the power law that related applied load and penetration depth.40 The widely used Oliver and Pharr procedure for the analysis of unloading curves is based on Sneddon’s analysis of contact mechanics, where it is shown that the unloading exponent is bound between 1 and 2. For AFM nanoindentation of polymers, unloading exponents that clearly exceed this range are often observed, also in our experiments discussed here. Following the Oliver and Pharr procedure, as described in the Supporting Information, the elastic modulus is proportional to the ratio of the unloading slope and the square root of the contact area, Ac.51 Since the geometric constant, R, in Ac = Rhc2 is not known for the AFM tip, we initially set R = 40 to superimpose the curves shown in Figure 3A,B. Figure 3B shows the modulus evaluated following this procedure, for several force curves obtained on the PS sample with a 8 nm thin ppHMDSO cover layer.48,52 Each point in the plot refers to a single force curve. Similar to Figure 3A, the modulus decreases after the first few nanometers, reaching a plateau when the behavior of the PS thin film dominates the stress and displacement fields. Interestingly, after approximately 50 nm penetration depth, the evaluated Young’s modulus increases slightly, probably because of the effect of the stiff silicon substrate. This observation is in good agreement with Figure 3C, as it will be discussed in the following. However, the value for R is almost twice the value for the case of a blunt Berkovich indenter, which is unreasonable considering that a sharp AFM tip was used in the experiment. The value calculated for paraboloid with a radius of curvature at the apex similar to an AFM tip would be in the range of 1-3. Therefore, we stress that the results shown in Figure 3B should be considered at most to be semiquantitative; the trend, however, is confirmed by Figure 3A. The ratio of unloading and loading slopes close to the point of maximum load, shown in Figure 3C, further confirms the results of Figure 3B. This ratio is supposed to be equal to 1 for a perfectly elastic contact, where the two curves are superimposed. It increases after approximately 15-20 nm penetration depth, i. e., when the PS is indented, and finally reaches a pronounced maximum at approximately 50 nm penetration depth for the 8 nm thin top layer. When the contribution of the silicon substrate becomes important and hence more elastic contribution is present, this ratio decreases again. However, this takes place below the 70 nm penetration depth, as one would expect after Figure 2.
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Figure 4. Force curves, shown as the average of 20 single force curves, of the PS reference (empty symbols) and the PS/ppHMDSO 8 nm thick sample (filled symbols) with a linear fit as discussed above.
In conclusion, we showed that the nanomechanical behavior of a hard ultrathin layer coupled with a thin soft layer gives rise to an enhanced elastic (i.e., stiffer) behavior that can be modeled as the separate contributions of each material at its peculiar deformation mode. The ppHMDSO covered PS is hence stiffer than the corresponding spin-coated PS film without the plasma polymerized top coat; see Figure 4. Evidence that further supports this finding was obtained from AFM images of the imprints left behind the indenter. The loads and penetration depths that trigger each deformation mode were also quantitatively evaluated. Although a rough estimation of the Young’s modulus of the components is possible, the analysis of AFM unloading curves is still difficult, justifying further studies on this subject.
’ EXPERIMENTAL SECTION Samples. Poly(styrene) (Mw = 163 000 g/mol) was spincoated from a solution of toluene with a concentration of 20 mg/ mL at 2000 rpm. Plasma polymerization of HMDSO (ppHMDSO) was carried out in a home-built inductively coupled plasma reactor made of Pyrex.53 The plasma power was generated by a 13.56 MHz radio frequency generator (RFG150, Coaxial, Power Systems Ltd, Eastbourne, E. Sussex, U.K.). The plasma reactor was evacuated to a base pressure of 0.005 mbar before the monomer was induced until reaching a constant pressure of 0.06 mbar. Subsequently, a flux of 14 cm3(STP)/min oxygen was added to the reactor. The ppHMDSO films were deposited at a plasma power of 90 W with a duty cycle of 0.025. The thickness of ppHMDSO layers was measured by microfocus grazing incidence small-angle X-ray scattering (μ-GISAXS) experiments.54 The bulk PS sample was prepared by heating PS (Mw = 163 000 g/mol) to 150 °C within 1 h. This temperature was held for another hour, and then the sample was pressed with a load of 20 kN at 150 °C for 1 h and cooled to room temperature within 1 h. Nanoindentations were carried out with a Multimode IV (Veeco, Santa Barbara, CA) AFM equipped with an Ev scanner, allowing one to carefully choose the area of the sample where to perform indentations and to capture postindentation images, in order to measure the residual imprint and evaluate the occurrence of plastic phenomena such as pileup. Silicon cantilevers 3389
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The Journal of Physical Chemistry B (Olympus AC160TS) were used, with nominal elastic constant of 40 N/m. The cantilever elastic constant was evaluated according to Green et al.55 on the basis of the knowledge of the cantilever geometry. The tip radius of 7.5 ( 2.5 nm was evaluated by blind estimation of several AFM images.56 The deflection sensitivity, which allows one to convert the voltage output from the position-sensitive diode into the cantilever deflection, was measured by indenting a silicon substrate before and after the nanoindentation experiments. Nanoindentations were performed with loads in the range of 0.5-7.0 μN, with an indenter rate of 18 μm/s.
’ ASSOCIATED CONTENT
bS
Supporting Information. Figures showing surface profiles corresponding to Figure 1, superposition of ideal tip shapes, comparison of residual indents, as well as residual imprints, for the PS/ppHMDSO, tip shape calculated from images of a standard tip characterizer, Sneddon's elastic deformation for the PS film, and contact depth vs maximum penetration depth, tables showing blind estimates of tip radii, and text describing additional experimental details and calculations and giving associated references. This material is available free of charge via the Internet at http://pubs.acs.org.
’ AUTHOR INFORMATION Corresponding Author
*
[email protected]. *
[email protected]. )
Current Address
Borealis Polyolefine GmbH, Innotech Operational Support, Advanced Polymer Characterisation, Sankt Strasse 25, 4021, Linz, Austria.
’ ACKNOWLEDGMENT We are grateful to S. Emmerling, S. K. Nett, M. C. Lechmann, S. Lenz, R. S. T. Kappes, J. S. Gutmann (Max Planck Institute for Polymer Research), and S. V. Roth (HASYLAB) for their support in microfocused X-ray small-angle grazing incidence measurements. H.-J. Butt and R. F€orch are acknowledged for fruitful discussions. We gratefully acknowledge partial support from the DFG priority program SPP1369 and the Alexander von Humboldt Foundation via Postdoctoral Fellowships to D.T. and Y.Z. ’ REFERENCES (1) Low, I. M.; Duraman, N.; Mahmood, U. Mater. Sci. Eng. C 2008, 28, 243. (2) Montell, D. J. Science 2008, 322, 1502. (3) Pal, E.; Seemann, T.; Zollmer, V.; Busse, M.; Dekany, I. Colloid Polym. Sci. 2009, 287, 481. (4) Takahashi, K.; Kitaguchi, H.; Doi, T. Supercond. Sci. Technol. 2009, 22, 025008. (5) Barber, P.; Houghton, H.; Balasubramanian, S.; Anguchamy, Y. K.; Ploehn, H. J.; zur Loye, H. C. Chem. Mater. 2009, 21, 1303. (6) Mailhot, B.; Rivaton, A.; Gardette, J.-L.; Moustaghfir, A.; Tomasella, E.; Jacquet, M.; Ma, X.-G.; Komvopoulos, K. J. Appl. Phys. 2006, 99, 104310. (7) Xie, Z. H.; Hoffman, M.; Munroe, P.; Bendavid, A.; Martin, P. J. Acta Mater. 2008, 56, 852. (8) Wiesmann, D.; Rawlings, C.; Vecchione, R.; Porro, F.; Gotsmann, B.; Knoll, A.; Pires, D.; Duerig, U. Nano Lett. 2009, 9, 3171.
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dx.doi.org/10.1021/jp112211c |J. Phys. Chem. B 2011, 115, 3385–3391
