Tethering macromolecules to surfaces represents a versatile approach for functionalizing, protecting, and structuring both organic and inorganic materials. In this study, thin films of poly(acrylamide) (PAAm) brushes and covalently cross-linked hydrogel brushes were grown from iniferter-functionalized silicon substrates by UVLED-initiated photopolymerization and their properties subsequently studied by means of a variety of analytical methods. The employed photografting method allowed the controlled fabrication of very thick films (up to 1 μm) in an aqueous environment, over a period of less than 1 h of polymerization and in the absence of side reactions. PAAm covalently cross-linked hydrogel brushes were prepared by feeding trace amounts of the cross-linker bis(acrylamide) (up to 1.0 wt % of monomer solution) into the reaction vessel. Both bulk and interfacial properties of these polymer films were found to be strongly influenced by lateral cross-linking of the grafted polymer chains. In agreement with theoretical expectations, the decrease of polymer-brush conformational freedom with increasing cross-link density resulted in a substantial increase of film wettability with water. The swelling ratio of the hydrogel brushes, as measured by ellipsometry and atomic force microscopy (AFM), also confirmed the formation of grafted networks and was found to be directly related to the amount of cross-linker in the monomer feed. In addition, the Young’s moduli and tribological properties of PAAm brushes and hydrogel brushes were tuned by adjusting the cross-linker concentration. Because of the additional constraint given by the surface tethering of each chain end, intermolecular cross-linking generated very high mechanical stresses within the brush structure. Covalently cross-linked hydrogel brushes thus displayed higher Young’s moduli and coefficients of friction, when compared to the grafted polymer-brush analogues. These hydrogel brushes present an opportunity for readily tailoring physical properties, especially as they allow tuning of the physical characteristics of surfaces while maintaining the interfacial chemical composition nearly constant.

AFM nanoindentations show a dependence of penetration, i.e., the relative motion between the sample and the tip (indenter), on material elastic properties when using the same load. This relationship becomes visible by using of samples being homogeneous down to the scale of nanoindentation. They were prepared from materials covering a broad range of mechanical behavior: from rubbery networks to glassy and semicrystalline polymers. The elastic modulus can be obtained applying Sneddon's elastic contact mechanics approach. To do this, some calibrations and instrumental features have to be measured accurately. All the polymers tested show that the contact between the tip and the sample is dominated by elastic behavior with negligible plastic deformation. In contrast to a standard metallic material like lead, the penetration dependence on load follows an exponent of 1.5, consistent with elastic contact mechanics. This can be justified on the basis of the large elastic range polymers exhibit, on the constraints due to the geometry of the deformation during indentation and to the critical yielding volume needed in order to induce plasticity. For the polymers studied, this volume is chosen in such a way that a significant material volume is irreversibly deformed. Elastic moduli taken from AFM force curves show a very good agreement with bulk values obtained by macroscopic tensile testing, on all the polymers tested. This result confirms that AFM nanoindentations in polymers take place mostly in the elastic range and opens the possibility to characterize the mechanical behavior of polymers on an unparalleled small scale compared to commercial DSI (depth sensing instruments), which use a much blunter indenter.

The analysis of nanoindentation force curves collected on polymers through the common Oliver and Pharr procedure does not lead to a correct evaluation of Young's modulus. In particular, the estimated elastic modulus is several times larger than the correct one, thus compromising the possibility of a nanomechanical characterization of polymers. Pile-up or viscoelasticity is usually blamed for this failure, and a deep analysis of their influences is attempted in this work. Piling-up can be minimized by indenting on a true nanometer scale, i.e., at penetration depth smaller than 200 nm. On the other side, it is common knowledge that fast indentations minimize the effect of viscoelasticity. However, changing the indentation time in a broad range of contact time (fractions of second up to hundreds of seconds) did not allow the correct estimation of Young's modulus for the polymers used in this work. The final result is that the Oliver and Pharr procedure as well as any other procedure analyzing the unloading curve with elastic contact mechanics models cannot be employed to measure Young's modulus of polymers because its application is incorrect from a theoretical point of view, unless the analysis is limited to the very first nanometers of penetration depth when the contact is perfectly elastic. Viscoelastic contact mechanics models should instead be employed to characterize these materials.

The convolution of tip shape on sample topography can introduce significant inaccuracy in an AFM image, when the tip radius is comparable to the typical dimension of the sample features to be observed. The blind estimation method allows one to obtain information on the AFM tip through an unknown characterizer sample and thus to perform the deconvolution of the tip shape from an image. When applying the blind estimation method to determine the AFM tip shape, some apparently trivial issues relating to the experimental operating parameters must be taken into account. In this paper, the effects of the operating parameters, e.g., sampling intervals (resolution) and instrumental noise, have been taken into account for the practical use of blind estimation and the result is that instrumental noise tends to provide a smaller estimation of the tip size, while larger sampling intervals provide a larger value of it. This paper presents guidelines to those effects in AFM and appropriate experimental conditions for applying the blind estimation method to obtain more reliable data on tip radius and therefore on sample topography.

The Young's modulus of cross-linked poly(dimethylsiloxane) (PDMS) surface was quantitatively investigated as a function of UV/ozone treatment time across different length scales. An AFM was used to probe PDMS surface mechanical properties at the nanometer length scale. The Young's modulus of each sample was estimated with continuum contact mechanics theory (Sneddon method) using AFM data by employing the hyperboloid tip shape model. A custom-built ATD device (JKR method) was also used which allowed us to simultaneously monitor the load, the contact area, and the relative displacement between (a) a lens (made either of PDMS elastomer or of Si 3 N 4 ) and (b) surface-treated PDMS films upon loading and unloading on the micrometer to submillimeter length scales. The modulus of PDMS increased with increasing treatment time as observed by AFM as well as by ATD, which we explained by the gradual formation of a silica-like layer. For all specimens tested, the modulus values obtained were highest from AFM, lower from ATD, and lowest from bulk tensile experiments for the same UV/ozone dose. These results demonstrate the effect of the probed length scale of the tests used to assess mechanical performance.

We report on the successful replication of the smallest pores in anodized aluminum oxide (AAO) via the layer-by-layer (LBL) deposition of polyelectrolytes to date to yield free-standing, open nanotubes with inner and outer diameters (±2σ) down to 37 ± 4 and 52 ± 19 nm, respectively. This work is based on the fabrication of defined arrays of highly regular nanopores by anodic oxidation of aluminum. Pores with pore diameters between 53 ± 9 and 356 ± 14 nm and interpore distances between 110 ± 3 and 500 ± 17 nm were obtained using an optimized two-step anodization procedure. 3-(Ethoxydimethylsilyl)propylamine-coated pores were replicated by alternating LBL deposition of poly(styrenesulfonate) and poly(allylamine). The detrimental adsorption of polyelectrolyte on the top surface of the template that typically results in partial pore blocking was eliminated by controlling the surface energy of the top surface via deposition of an ultrathin gold layer. The thickness of the deposited LBL multilayer assembly at the pore orifice agreed to within the experimental error with the thicknesses measured by variable angle spectroscopic ellipsometry and atomic force microscopy (AFM) for layers assembled on flat substrates. The selective dissolution of the alumina template afforded free-standing, open polymer nanotubes that were stable without any cross-linking procedure. The nanotubes thus obtained possessed mean outer diameters as small as 52 nm, limited by the size of the AAO template.

The atomic force microscope (AFM), apart from its conventional use as a microscope, is also used for the characterization of the local mechanical properties of polymers. In fact, the elastic characterization of purely elastic materials using this instrument can be considered as a well-assessed technique while the characterization of the viscoelastic mechanical properties remains the challenge. In particular, one finds the mechanical behavior changing when performing indentations at different loading rates, i.e. on different time scales. Moreover, this apparent viscoelastic behavior can also be due to complex contact mechanics phenomena, with the onset of plasticity and long-term viscoelastic features which cannot be identified by the force curve alone. For this reason, a viscoelastic characterization, and thus the study of the effects of indentation rate and temperature, was done on model materials where such additional phenomena are not observed. Another time dependence originating from the instrument itself has also been identified and decoupled. In fact, the viscoelastic behavior has been found to be reproducible even if one changes the experimental set-up as far as the preliminary determinations concerning AFM nanoindentations are well performed. The effects of temperature and time scales on the mechanical behavior have also been undertaken. A check on time–temperature superposition is also attempted through the WLF equation and the apparent activation energies for the elementary motions in the rubbery and in the glass transition regions are in good agreement with the expected values.

More than 50% of the thermoplastic polymers applied globally are semi‐crystalline, making crystallization part of the material and component design process for these materials. The mechanical and optical performance, but also the long‐term stability of the final articles and applications will depend upon three major groups of influence factors: Polymer structure and monomer composition, additive addition (especially nucleation) and blending with secondary components, and processing parameters. In the present review, we try to establish a connection between these three areas, the resulting combination of crystallinity and morphology, and the final application properties. Examples are drawn from the two major polyolefins, polyethylene and polypropylene, technical polymers like polyamide and polyester, but also polymers from renewable sources like poly(lactic acid).