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.