Amplified Responsiveness of Multilayered Polymer Grafts: Synergy

Article pubs.acs.org/Macromolecules

Amplified Responsiveness of Multilayered Polymer Grafts: Synergy between Brushes and Hydrogels Shivaprakash N. Ramakrishna,† Marco Cirelli,‡ E. Stefan Kooij,§ Michel Klein Gunnewiek,‡ and Edmondo M. Benetti*,†,‡ †

Laboratory for Surface Science and Technology, Department of Materials, ETH Zürich Vladimir-Prelog-Weg 5, HCI F 537, 8093, Zürich, Switzerland ‡ Department of Materials Science and Technology of Polymers, MESA+ Institute for Nanotechnology, University of Twente, P.O. Box 217, 7500 AE Enschede, The Netherlands § Physics of Interfaces and Nanomaterials, MESA+ Institute for Nanotechnology, University of Twente, P.O. Box 217, 7500 AE Enschede, The Netherlands S Supporting Information *

ABSTRACT: The responsive properties of surface-grafted polymer films in aqueous media can be amplified by covalently layering thermosensitive brushes and hydrogels. This was demonstrated by synthesizing layers of linear poly(N-isopropylacrylamide) (PNIPAM) brushes, alternating with cross-linked, poly(hydroxyethyl)methacrylate (PHEMA) hydrogels via sequential surface-initiated atom-transfer radical polymerization (SI-ATRP) steps. Below the lower critical solution temperature (LCST) of PNIPAM, brush/hydrogel multilayered films swell similarly to linear PNIPAM homopolymer brushes, as measured by liquid ellipsometry. In contrast, above the LCST, the PHEMA hydrogel interlayer acts as stiffening element within the collapsed multilayered film, as monitored by atomic force microscopy (AFM) nanoindentation and lateral force microscopy (LFM). This translates into a 10-fold increase in Young’s modulus by the collapsed, layered films compared to PNIPAM homopolymer analogues. The (macro)molecular continuity between the brush main chains and hydrogel constituents thus enables a chemically robust layering to form graded, quasi-3D grafted polymer architectures, which display a concerted and amplified temperature-triggered transition.

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INTRODUCTION The development of organic and hybrid materials with graded physical, morphological or chemical properties has been the subject of increasing interest in different fields of materials science. In nature, a number of mineralized tissues, such as the outer skeleton of invertebrate species are composed of hybrid layered organic/inorganic systems.1,2 In these cases, multilayer architectures featuring diverse compositions and organized in a hierarchical order, impart excellent mechanical properties to the biological materials. A prime example is nacre in shells, where platelets of CaCO 3 minerals, intercalated by organic components, provide both strength and toughness.3 In other cases, layers with different compositions and structures with discontinuous properties provide not only mechanical resistance but also specific functions to the dermal components of a number of different species. This has been demonstrated both for the scales of ancient fish armor4 and in the multilayer dermis of snakes.5 Inspired by these natural counterparts, materials scientists have designed and fabricated graded, composite materials in order to “amplify” their properties when compared to the individual, constituting elements. Among the most successful techniques to produce synthetic, layered three-dimensional (3D) architectures, bottom-up approaches, such as layer-by-layer (LbL) deposition6 of © XXXX American Chemical Society

polymer species and/or inorganics, have allowed the fabrication of graded materials with enhanced mechanical properties and highly functional character.7−9 This technique involves the assembly of polymeric species on the surface of a substrate to fabricate films made of at least two different components, held together by physical or chemical interactions. Despite the versatility of LbL technique, when the layered components are bound by reversible physical interactions (electrostatic, hydrogen bonding or hydrophobic interactions)10,11 they can suffer from structural instability or uncontrolled mixing between the components.12 Alternatively, even when chemical treatments are applied to cross-link or covalently “fix” the assemblies, precise control over the relative loading of one component is often constrained by a self-limiting adsorption process.6,10 Thus, generally, it is hard to guarantee molecular continuity between the assembled layers, and this characteristic can become a limitation in the presence of swelling solvents, which could lead to delamination or disruption of the material. The inspiration behind the present study was the need to overcome the general drawbacks of assembled, layered Received: July 14, 2015 Revised: August 31, 2015

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DOI: 10.1021/acs.macromol.5b01556 Macromolecules XXXX, XXX, XXX−XXX

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Scheme 1. Fabrication of Bi-Layered Brush/Hydrogel and Tri-Layered Brush/Hydrogel/Brush Films by Sequential SI-ATRPa

a

Key: (a) formation of initiator-functionalized silicon oxide surfaces; (b) SI-ATRP of NIPAM performed in 50:50 water:methanol solutions for short polymerization times; (c) re-initiation of SI-ATRP in the presence of HEMA:EGDMA(0.5 mol %) from a PNIPAM brush layer which is collapsed due to co-nonsolvency; (d) final grafting of interfacial PNIPAM brush layer by “fast” SI-ATRP under aqueous conditions.

materials by synthesizing polymeric layers in situ, exploiting surface-initiated polymerization (SIP)13 from a supporting substrate and sequentially varying the monomer composition over two or more synthetic steps. In this way, layered polymeric films presenting a vertically structured, quasi-3D morphology, characterized by a (macro)molecular continuity between each constituting layer can be synthesized within a single chemical process. We have especially focused on multilayered films comprising linear polymer “brush” layers,14 based on poly(Nisopropylacrylamide) (PNIPAM), alternated by a covalently cross-linked, poly(hydroxyethyl)methacrylate (PHEMA) hydrogel layer15,16 to form two-layer (brush/hydrogel) and threelayer (brush/hydrogel/brush) architectures. PNIPAM brushes are temperature responsive, collapsing above their lower solution critical temperature (LCST) centered at 30−35 °C,17−21 while the PHEMA hydrogel layer acts as stiffening element across this transition, due to its covalently cross-linked structure. We recently demonstrated that the sequential reinitiation of SIP in selective solvent environments allowed the successful grafting of multilayered films with well-defined chemical composition and physical properties for each layer.22 The use of a selective solvent during sequential grafting causes the first, initially grafted layer (“bottom” brush) to act as a physical resist for the propagation of the “top” layer, during a second SIP step. Hence, this strategy enables the reinitiation to be confined to the brush−medium interface (Scheme 1) and avoids the formation of chemically mixed brushes, which are often obtained when block copolymerization is performed.23,24 This method was recently employed by us to synthesize vertically well-defined bilayered films presenting a covalently cross-linked PAAm hydrogel on top of linear polyacrylamide (PAAm) brushes.22 Nevertheless, in this previous report we limited the chemistry of the films to PAAm and we restricted the fabrication to bilayered architectures. In the present study, we have enlarged the concept of sequential SIP-based fabrication, reporting the synthesis of stimulus-responsive multilayered brush films presenting alternating chemistries and different brush architectures. PHEMA-hydrogels were layered between linear PNIPAM brushes forming bi- and trilayered films by means of sequential, surface-initiated atom-transfer radical polymerization (SI-ATRP) steps.25 Selective solvent environments were ensured by the temperature-dependent cononsol-

vency behavior of PNIPAM, which, in HEMA-water mixtures, displayed a LCST well below the polymerization temperature of 25 °C (Scheme 1b and Figure S1).26−28 Swelling, nanomechanical and tribological properties of bilayered PNIPAM/PHEMA-hydrogel and trilayered PNIPAM/PHEMA-hydrogel/PNIPAM films were subsequently investigated by variable angle spectroscopic ellipsometry (VASE), atomic force microscopy (AFM) nanoindentation and lateral force microscopy (LFM), both below and above PNIPAM LCST. These measurements highlighted how within bi- and trilayered architectures the molecular continuity between brush and hydrogel layers and the graded variation of composition strongly influenced the physicochemical and mechanical characteristics of the films, amplifying their transition across the LCST of PNIPAM. The fabrication approach that is described thus offers a new strategy to production of quasi-3D polymeric architectures, with a well-defined gradation of physicochemical properties, in a single process, by exploiting the sequential reinitiation of SIATRP. As we recently have shown, substrate-immobilized covalently cross-linked brush films can be easily detached from the support by employing photocleavable initiator anchors, to form freestanding membranes.29,30 Thus, sequential SIP could be applied to fabricate layered membranes with fully tunable and responsive mechanical properties, and which can be detached from the initial support and subsequently applied to different substrates. In addition, this fabrication approach could be easily scaled up to functionalize large surfaces and design supports for biomaterial formulations. In this latter application, a precise and significant variation of properties in response to an external stimulus would be highly desirable, in order to modulate the interaction of the biomaterial with the surrounding biological medium.

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RESULTS AND DISCUSSION Fabrication of Layered Brush/Hydrogel Films by Sequential SI-ATRP. The synthesis of the multilayered brush/hydrogel films as well as the corresponding homopolymer brushes and hydrogels were carried out by SI-ATRP. In particular, sequential grafting steps applying different polymerization solutions were performed, in order to synthesize layered films. The relative concentrations of activator and deactivator species, the solvent mixtures and the polymerization times were B


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Figure 1. (a) PNIPAM brush growth rates measured by VASE following SI-ATRP performed in different water:methanol solvent mixtures. (b) FTIR adsorption spectra recorded on homopolymer and multilayered films.

Table 1. Ellipsometric Thickness Values of the Different Synthesized Films under Dry Conditions and in Ultrapure Water, at 25 and 40 °Ca ellipsometry sample PNIPAM-I PHEMA hydrogel PNIPAM-II bilayered trilayered

dry (nm) 6 45 75 50 84

± ± ± ± ±

2 2 2 2 4

H2O 25 °C (nm) 14 59 249 88 248

± ± ± ± ±

3 4 2 4 5

H2O 40 °C (nm) 8 66 97 52 102

± ± ± ± ±

water content 25 °C (%)

water content 40 °C (%)

σ (chains nm‑2)

57 25 70 41 66

25 33 23 4 17

0.29 n.a. 0.17 n.a. n.a.

1 2 3 4 3

water CA (deg) 55 65 57 63 56

± ± ± ± ±

7 5 5 5 5

The estimated water content (as volume%) for both homopolymer and multilayered films are reported. The static water contact angle (CA) measured on each film and the estimated grafting densities (σ) of PNIPAM homopolymer brushes calculated by eq 2 (see Experimental Section for details) are also shown. a

thicker PNIPAM brushes were obtained for a given polymerization time). However, this was accompanied by an increase of irreversible termination between active chains due to the decomposition of the deactivator species.38 This phenomenon reflected a slowing down of the PNIPAM−brush growth rate, until a plateau was reached (Figure 1a). Thus, in order to enable subsequent reinitiation to fabricate bi- and trilayered films, SI-ATRP of the PNIPAM brush underlayers was usually performed for short polymerization times, typically no longer than 10 min (long before irreversible termination took place) and employing 50:50 water:methanol mixtures. In addition, following the required reaction time, quenching solutions of degassed deactivator species were injected into the reaction medium, in order to stop the polymerization and retain a high concentration of active, bromide-bearing chain ends for subsequent reinitiation during the next SI-ATRP step.35−37 The effective grafting of a substrate-bound PNIPAM brush from initiator-modified silicon oxide surfaces was confirmed by ellipsometry, water contact angle (CA) and FTIR spectroscopy. In a typical fabrication process, following the initial SI-ATRP step (10′ polymerization time), a polymeric film with 6 ± 2 nm of dry thickness was measured by VASE. The static water CA decreased from 80 ± 3°, characteristic of an ATRP-initiatormodified silicon oxide surface, to 55 ± 7° (PNIPAM-I in Table 1), indicating the formation of a uniform PNIPAM layer on the surface. The FTIR spectrum of this film confirmed the presence of PNIPAM brushes, displaying two clear bands at 1650 and 1550 cm−1, which were attributed to the CO stretching of the amide-bearing monomer units (PNIPAM-I FTIR adsorp-

tuned, both to guarantee efficient reinitiation after each polymerization step, and to adjust the brush thickness for each layer (see Experimental Section for details). Under experimentally optimized conditions, typical bilayered PNIPAM/PHEMA-hydrogel films presented a 5−10 nm thick PNIPAM brush underlayer supporting a 30−40 nm-thick layer of PHEMA-hydrogel, while trilayered PNIPAM/PHEMAhydrogel/PNIPAM films presented an additional 30−40 nmthick PNIPAM interfacial brush layer. Controlled, rapid synthesis of thick PNIPAM brushes under aqueous conditions by SI-ATRP was a challenging task due to the catalyst poisoning by NIPAM, which has already been reported, combined with the decomposition of deactivator species triggered by the presence of water.31−33 A possible solution to these shortcomings was recently proposed by Zhang et al., who demonstrated how thick PNIPAM brushes can be obtained by Cu(0)-mediated controlled radical polymerization in aqueous media.34 In the present study, the growth rates of PNIPAM brushes were tuned by adjusting the solvent mixture (water:methanol) during SI-ATRP. This approach allowed us to easily modulate the relative thickness of each PNIPAM brush layer, additionally shortening the polymerization time required to reach the targeted brush-thickness in order to minimize irreversible termination by radical recombination.35−37 As shown in the kinetic plots in Figure 1a, SI-ATRP of NIPAM in water:methanol mixtures allowed the synthesis of relatively thick PNIPAM brushes over short polymerization times (1−3 h). An increase in the relative water content in the polymerization medium led to a faster polymerization rate (i.e., C


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Scheme 2. Schematic Representation of Bilayered (a), Trilayered Films (b) and Substrate-Bound Homopolymer PNIPAM Brush and PHEMA−Hydrogel Components (c) Synthesized by SI-ATRP

After the same reaction time as applied for the layered samples, a homopolymer PHEMA-hydrogel with an average thickness of 45 ± 2 nm was obtained (Table 1). This value was comparable to the thickness of the PHEMA-hydrogel layer within the bilayered films and thus demonstrated the efficiency of reinitiation and block copolymerization from the underlying PNIPAM-I brush (a detailed discussion on the reinitiation efficiency from different PNIPAM brush underlayers is reported in the Supporting Information). Interestingly, the use of a “poor” solvent for the PNIPAM-I brush during SI-ATRP of HEMA-EGDMA did not affect the efficient grafting of the second layer. This was presumably due to the rather thin brush underlayer and the high concentration of active chain ends exposed at the brush interface, obtained by applying deactivator solutions at the end of the first SI-ATRP step. The third, interfacial PNIPAM brush layer was finally grafted from PNIPAM/PHEMA-hydrogel bilayered films in 75:25 water:methanol mixtures (see Experimental Section for details) to form trilayered PNIPAM/PHEMA-hydrogel/PNIPAM films. A higher water content during this last SI-ATRP step was employed, in order to induce relatively fast brush growth, yielding trilayered films of comparable dry thickness for each layer type. After this last step, a dry film thickness increase of around 70% was measured by VASE, after 10 min of SI-ATRP, with a typical trilayered film reaching 84 ± 5 nm (trilayered film reported in Table 1). FTIR spectroscopy and CA measurements confirmed the grafting of an additional PNIPAM brush layer, as shown in Figure 1b and Table 1. In particular, the relative intensity of the CO signal (amide) at 1650 cm−1 significantly increased, while the water CA recorded on the trilayered film decreased to 56 ± 5°, indicating the presence of PNIPAM brushes at the film interface. Also in this case, a substrate-bound PNIPAM brush was synthesized from an initiator-functionalized silicon oxide substrate, applying the same polymerization conditions used for the layered films (to obtain PNIPAM-II). As can be seen in Figure 1a and Table 1, following 10 min of polymerization in 75:25 water:methanol

tion profile in Figure 1b). The second SI-ATRP step involved the grafting of the PHEMA-hydrogel layer from the PNIPAM-I brush. This was performed in 50:50 HEMA:water solutions and included 0.5 mol % of ethylene glycol dimethacrylate (EGDMA) in the monomer mixture as cross-linker. Relative concentrations of EGDMA between 0.1 and 2 mol % have previously been shown to allow the formation of brushhydrogels with controlled swelling and mechanical properties, avoiding film wrinkling due to the mechanical stress induced by the cross-linking of substrate-grafted chains.15 In the present study we used 0.5 mol % of EGDMA, to ensure relatively high swelling of the PHEMA-hydrogel layers in aqueous medium. Notably, the polymerization mixture used for the grafting of PHEMA-hydrogel was a poor solvent for the underlying PNIPAM-I (Figures S1 and S2, in the Supporting Information). In similar water:alcohol mixtures, in fact, PNIPAM showed a cononsolvency behavior with a consequent shift of the LCST below ambient temperature.19,28 50:50 HEMA(EGDMA):water solutions used for SI-ATRP of the PHEMA-hydrogel layer induced a similarly low LCST, which was found below 10 °C under these conditions. The collapse of the substrate-bound PNIPAM-I thus limited the diffusion of HEMA-EGDMA to any unreacted initiator species at the silicon oxide surface, confining reinitiation to the active chain ends exposed on the PNIPAM-I brush interface (Scheme 1c). SI-ATRP of HEMAEGDMA from the PNIPAM-I brush caused an increase of film thickness to 50 ± 5 nm, over 20 min of reaction. Following this second grafting step, the CA values increased to 63 ± 5° (bilayered sample in Table 1), confirming the presence of an interfacial PHEMA-hydrogel film. FTIR spectroscopy also verified the successful grafting of PHEMA from the underlying PNIPAM-I brush. As reported in Figure 1b, a strong band centered at 1750 cm−1, originated from the CO stretching of the ester groups in HEMA units, characterized the spectra of the bilayered films. A substrate-bound, homopolymer PHEMAhydrogel was also synthesized following similar polymerization conditions from initiator-functionalized silicon oxide surfaces. D


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PNIPAM homopolymer brushes showed a water content of 60 and 70% at 25 °C, for thin and thick brushes, respectively (PNIPAM-I and PNIPAM-II in Table 1). This variation was presumably due to the differences in the average grafting densities between the two layers and the different amount of associated water between longer and shorter PNIPAM grafts. When immersed in water at 40 °C, both the films collapsed, partially expelling water. As reported in Table 1, the residual included solvent at 40 °C was 25 and 23%, for thin and thick PNIPAM brushes, respectively. For both these films, the morphology of the brush surface underwent a marked roughening above the LCST, as observed by AFM. As can be seen in Figure S5, PNIPAM-I brush displayed a uniform, swollen interfacial morphology at 25 °C. Above LCST, at 40 °C, the collapse of the grafted chains induced the formation of hydrophobic aggregates covering the brush interface.20,42,43 The transition of PNIPAM brushes across LCST could also be monitored by AFM nanoindentation. As reported in Figure 2,

medium, PNIPAM homopolymer brushes showed an average dry thickness of 75 ± 10 nm. This value was significantly higher compared to the thickness recorded for the interfacial PNIPAM brush grafted from the bilayered films during the third SI-ATRP step (which was estimated to be around 30 nm thick). This was presumably due to the irreversible termination by radical transfer or bimolecular coupling, which reduced the concentration of active chain-ends on the second PHEMA−hydrogel layer and thus affected the reinitiation efficiency from the bilayered films. Structural Characterization of Multilayered Films. In order to gain further insight into the structure and swelling properties of the synthesized films, ellipsometry in water was performed for both the homopolymer brushes/hydrogels and the multilayered films. From the measurements of the swelling ratios (and thus the polymer volume fraction Tdry/Twet) of linear PNIPAM brushes, an estimate of their average grafting density σ could be obtained, in accordance to eq 2 (see Experimental Section for details).39−41 For the PNIPAM brush underlayer (corresponding to substrate-bound PNIPAM-I homopolymer brush) a Tdry/Twet = 0.43 was found, which corresponded to σ = 0.29 chains·nm−2. A thicker PNIPAM-II homopolymer brush had a higher water content at 25 °C, resulting in a Tdry/Twet = 0.3 and, consequently, an estimated average grafting density σ = 0.17 chains·nm−2. In addition, the difference between the swollen thickness of bilayered and trilayered films provided an estimate of the swelling ratio of the interfacial PNIPAM brush layer. This resulted in a Tdry/Twet = 0.2, which indirectly suggested a much lower σ for the interfacial compared to the substrate-bound PNIPAM brushes. This expected result was likely due to the progressive loss of chain density during the sequential SI-ATRP steps. Hence, recapitulating the composition and the swollen structure for each layer within the stratified films, a schematic representation of the bi- and trilayered films is depicted in Scheme 2. In this exemplary view, considered in water at 25 °C, a dense and relatively thin, substrate-bound PNIPAM brush supports a PHEMA-hydrogel to form bilayered films (Scheme 2a), while an additional highly swollen and less densely grafted PNIPAM brush extends at the interface to produce trilayered architectures (Scheme 2b). As already mentioned, the properties of multilayered films were compared to substrate-bound homopolymer brush and hydrogel components, including a thin PNIPAM brush (PNIPAM-I in Scheme 2c), a PHEMA-hydrogel and a thick PNIPAM brush film with a dry thickness that was comparable to that of the trilayered films (PNIPAM-II in Scheme 2c). Thermoresponsive, Nanomechanical, and Tribological Properties of Multilayered Films. The thermoresponsive properties of the layered films were studied in an aqueous environment at different temperatures, by a combination of ellipsometry and AFM. Specifically, the nanomechanical and tribological characteristics of the multilayered and the homopolymer films were investigated by AFM nanoindentation and lateral force microscopy (LFM), performed in ultrapure water below and above PNIPAM LCST (25 and 40 °C). PNIPAM brushes in water undergo a transition from the swollen to the collapsed state across the temperature range 30− 35 °C.17,18,20,21 This transition can be monitored by in situ ellipsometry, as a gradual dehydration of the swollen brush and a consequent decrease of film thickness (Figure S3 and S4, in the Supporting Information).

Figure 2. Force-vs-penetration profiles obtained by colloidal-probe AFM nanoindentation for PNIPAM-I (a) and PNIPAM-II (c) brushes, at 25 and 40 °C (red and green color traces, respectively). The distributions of E values measured by Hertz model-fitting are reported in parts b for PNIPAM-I and d for PNIPAM-II. A detailed description of the AFM nanoindentation experiments and data analysis is reported in the Experimental Section.

typical force-vs-penetration profiles at 25 °C for both PNIPAM-I and PNIPAM-II are characteristic of highly swollen brushes. Hertz model-fitting of these indentation profiles gave average Young’s moduli (E) of 110 and 40 kPa for PNIPAM-I and PNIPAM-II, respectively, with typical distributions of E values reported in Figure 2, parts b and d (the Hertzian plots for all the tested samples are reported in Figure S6). The collapse of PNIPAM brushes above LCST (40 °C) induced a marked steepening of the corresponding indentation curves (Figures 2a and 2c), resulting in apparent E values centered at 1000 and 370 kPa for PNIPAM-I and PNIPAM-II, respectively (Figure 2, parts b and d). Higher apparent E values for thinner PNIPAM-I brushes compared to PNIPAM-II, both below and above LCST, were already reported by Vancso et al. and were ascribed to the close vicinity of the underlying substrate.44 The swelling and the nanomechanical characteristics of surface-grafted PHEMA-hydrogels were also studied by combining ellipsometry and AFM nanoindentation. PHEMAE


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Figure 3. Force-vs-penetration profiles obtained by colloidal-probe AFM nanoindentation for the bilayered (a) and the trilayered (c) films, recorded both at 25° and 40 °C (red and green color traces, respectively). The distributions of E measured by Hertz model-fitting are reported in parts b and d for the bilayered and the trilayered films, respectively. In part d, the distribution of E values for PNIPAM-II are also shown for comparison. Schematics depicting the temperature-induced transitions for PNIPAM-II, the bilayered and the trilayered films are reported in part e. A detailed description of the AFM nanoindentation experiments and data analysis is reported in the Experimental Section.

hydrogel films did not show a significant variation of ellipsometric thickness between 25 and 40 °C in water (Table 1). This swelling behavior was also confirmed by AFM step-height measurements on mechanically “scratched” films immersed in water at the two tested temperatures (Figure S7). Hertz model-fitting applied on force-vs-penetration curves for substrate-grafted PHEMA-hydrogels thus resulted in average E values of 1400 kPa with similar distributions at both 25 at 40 °C, as shown in Figure S8. The swelling and the nanomechanical properties of bi- and trilayered film, below and above PNIPAM LCST, resulted from the combination of the behavior of the brush and hydrogel components, covalently linked together within a surfacegrafted, stratified structure (as exemplified in Scheme 2a,b). The bilayered PNIPAM/PHEMA-hydrogel, with 50 ± 2 nm dry thickness, swelled by 41% when immersed in water at 25 °C, reaching a thickness of 88 ± 4 nm. This value was higher compared to the expected swelling of the PNIPAM homopolymer brush (PNIPAM-I) and the PHEMA-hydrogel components at this temperature (88 vs 73 nm, as reported in Table 1). This was presumably due to the lower density of the PHEMA-hydrogel layer grafted from PNIPAM-I compared to a corresponding substrate-bound hydrogel. At 40 °C the ellipsometric measurements indicated a marked vertical shrinking by the whole bilayered film (4% of residual swelling, as reported in Table 1), which suggested the release of nearly all the solvent from the collapsed film. This phenomenon was also confirmed by AFM step-height measurements, which demonstrated a clear decrease in film height at 40 °C (Figure S7). Thus, both the PNIPAM brush underlayer and the PHEMA-hydrogel displayed a temperature-induced collapse above the PNIPAM LCST. This transition could be explained by considering the (macro)molecular continuity between the two layers, i.e. the covalent bond between each substrategrafted PNIPAM chain and the PHEMA backbones within the interfacial hydrogel layer. Shrinking of the underlying PNIPAM brush led to the collapse of the whole film, which also showed a

shift of E modulus distribution to higher values, due to the release of swelling solvent from the entire bilayered structure (as reported in Figure 3a,b). Fitting of the indentation profiles of bilayered films by a Hertzian model resulted in an average E value of 450 kPa at 25 °C. This markedly increased at 40 °C, reaching an average value of 3500 kPa and a wide distribution, as shown in Figure 3b. The behavior of the PNIPAM/PHEMA-hydrogel/PNIPAM trilayered films, across the PNIPAM LCST, was subsequently investigated. Ellipsometric measurements performed in water at 25 °C indicated that the trilayered films swelled profusely under these conditions. Namely, 84 ± 4 nm dry thick films incorporated 66% of water at 25 °C, reaching 248 ± 5 nm of swollen thickness (Table 1). Very similar values were also found by AFM step-height analysis, as reported in Figure S7. Also in the case of the trilayered films, increasing the temperature above the PNIPAM LCST led to the collapse of the entire, stratified film, which shrank to 102 ± 3 nm of thickness (Table 1 and Figure S7), with good agreement between ellipsometry and AFM step-height analysis. Simultaneously, the average Young’s modulus sharply increased from 60 kPa at 25 °C to 4500 kPa at 40 °C, with typical distributions reported in Figure 3d. It is interesting to compare the temperature response of the trilayered film and PNIPAM homopolymer brushes with a similar dry thickness. PNIPAM-II generally showed a higher water content compared to the trilayered films, by 4% at 25 °C and by 6% at 40 °C. These values were also confirmed by AFM step-height analysis, which showed how the trilayered films expelled more solvent in the collapsed state, resulting in lower film heights, if compared to the substrate-bound homopolymer analogue (Figure S7). A comparison of the changes in nanomechanical properties of PNIPAM-II and the trilayered films across the LCST further illustrated the effect of layering on film properties. As displayed in Figure 3d, where the distributions of E values for both the trilayered and the PNIPAM-II films are reported, the trilayered films showed a much more significant Young’s modulus shift between the F


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Figure 4. A typical friction-force loop obtained by LFM on PNIPAM-II brushes at 25° is reported in part a. Xd and Xs correspond to sections on the loop originating from brush lateral deformation and steady sliding, respectively; points S and S′ indicate the transition between these two regimes; scanning direction is indicated by two arrows on both trace and retrace profiles. The schematics reported in part b depict brush−probe interactions during a friction force loop. Two representative grafted chains are indicated in dark blue and green colors in order to highlight the lateral movement of the colloidal probe during the recording of the loop. The numbers on each schematic refer to the corresponding portion of the loop reported in (a): “1” the colloidal probe compresses PNIPAM brushes; “2” initial lateral scanning induces bending and stretching of the grafted chains until sliding occurs in “3”; after scanning direction reversal in “4″, PNIPAM brushes are laterally deformed in the opposite direction with respect to “2”, until a final steady sliding is reached in “6” and closes the loop. In part c, three friction force loops, recorded at 25 °C along 1, 3, and 7 μm of scanning distances are reported. In part d, three friction force loops recorded at 40 °C along the same scanning distances of part c are shown.

swollen and the collapsed states. In particular, both films displayed similar distributions of E at 25 °C. In contrast, at 40 °C the collapsed, trilayered films showed a 10-fold increase of the average Young’s modulus compared to the PNIPAM-II homopolymer brush under the same experimental conditions. This stiffness increment was likely due to the presence of the cross-linked PHEMA component, layered between the substrate-grafted and the interfacial PNIPAM brush layers. Below the PNIPAM LCST, the trilayered architecture swelled profusely and showed values of E similar to the PNIPAM homopolymer brushes with comparable swollen thickness. Above LCST, the collapse of the PNIPAM “blocks” led to the shrinking of the PHEMA hydrogel interlayer, which acted as a reinforcing component within the collapsed film. Thus, the effect of covalently layering linear, thermoresponsive PNIPAM brushes with a PHEMA hydrogel generated an “amplification” of the mechanical property change of the collapsed film. In order to further elucidate the interfacial properties of the multilayered films and compare them with the corresponding substrate-bound homopolymer brush and hydrogel layers, we performed lateral force microscopy (LFM). In particular, we concentrated on the analysis of the friction-force traces obtained by LFM and correlated the morphology of the friction “loops” with the shear characteristics of the different polymer films, both below and above the PNIPAM LCST.22

Swollen PNIPAM homopolymer brushes immersed in water at 25 °C displayed friction loops presenting a “tilted” section, after scanning direction reversal (Xd in Figure 4a). This portion of the friction loop originates from the lateral deformation of the hydrated grafts when they are subjected to sliding under load against the AFM colloidal probe, and correlates to a combination of brush lateral bending and stretching.45−47 The degree of tilt is related to the solvent content and the grafted-chain length. Steady sliding after brush lateral deformation can be attained when the applied shear force overcomes the spring force of the deformed brush (point S and S′ in Figure 4a). In Figure 4b, a schematic representation of the interaction between PNIPAM brush and colloidal probe during the recording of a friction loop is reported. After being compressed by the AFM probe (position “1” in Figure 4b), the brush is laterally deformed by Xd (position “2”) until sliding occurs (position “3”). After scanning direction reversal (position “4”) brush deformation takes place in the opposite direction (“5”), followed by sliding of the probe (“6”), which finally reaches the starting position on the loop. In Figure 4c, three different friction force micrographs, recorded on PNIPAM-II along 7, 3, and 1 μm of scanning distances are reported. As can be clearly seen, brush lateral deformation dominated the tribological behavior of the swollen brush at 1 μm of scanning distance. Under these conditions, G


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Figure 5. (a) Friction−force loops obtained by LFM on a substrate-bound PHEMA-hydrogel brush along 7 μm of scanning distance, recorded in ultrapure water at 25 (red color trace) and 40 °C (green color trace) on the same area. In part b, the friction force loop by LFM recorded on bilayered PNIPAM/PHEMA-hydrogel films at 25 °C (red color traces) is compared to the corresponding loop obtained on a PHEMA-hydrogel (gray color traces) at the same temperature. The effect of the temperature-induced transition on the friction force traces recorded on bilayered films is reported in part c. A schematic representation of an AFM colloidal probe initially deforming the PNIPAM brush underlayer (i) and subsequently attaining steady sliding on the interfacial PHEMA-hydrogel surface (ii) is reported in part d.

Thus, the tilted section on the friction traces was limited to few tens of nm, as shown in Figure 5b. As a confirmation of this characteristic tribological behavior of the bilayered films, an increase of the temperature of the medium above LCST was accompanied by a clear sharpening of the reversal region in the friction traces (point T in Figure 5c). This transition was presumably due to the collapse of the PNIPAM brush underlayer and the consequent dehydration of the entire bilayered film. Friction-force profiles of the trilayered films recorded over different scanning distances showed similar shapes to those of the corresponding friction loops observed for PNIPAM-II at 25 °C, but also showed effects due to the layered structure within the film. Completely tilted traces were observed at short scanning distances (pink trace in Figure 6a), indicating the absence of sliding and the dominance of lateral deformation on the swollen, interfacial PNIPAM brush. An increase in the scanning distances to 3 and 7 μm highlighted a reduction of

interfacial sliding was never attained and the friction traces exclusively showed a tilted profile. Remarkably, above LCST the tilted sections of the friction force traces were markedly reduced, irrespective of the applied scanning distance (Figure 4d). This was due to the partial solvent expulsion and the collapse of the PNIPAM brushes which, above LCST, form a more rigid and dehydrated film. In contrast to highly swollen, linear brushes, the PHEMAhydrogel brush showed friction-force profiles with straight vertical transitions at scanning direction reversals, which were followed by a steady interfacial sliding (Figure 5a). In addition, the morphology of the recorded loops did not show any significant change between 25 and 40 °C, in agreement with the constant swelling and mechanical properties displayed by these films across the temperature range studied. The friction-force profiles recorded on bilayered and trilayered films confirmed the stiffening effects occurring above LCST due to the covalent layering of linear PNIPAM brushes and cross-linked PHEMA. Friction traces of the bilayered PNIPAM/PHEMA-hydrogel at 25 °C showed a slight but characteristic tilting at direction reversal (Figure 5b). This could be better visualized upon comparison with the friction loop of the PHEMA-hydrogel, which showed a typical sharp transition when the sliding direction was reversed. A tilted trace at direction reversal had already been observed in the case of stratified PAAm brush/gel films22 and originated from the shear deformation of the swollen and compliant brush underlayer before reaching stable sliding on the interfacial, more rigid hydrogel. The PNIPAM/ PHEMA−hydrogel bilayered films presented a PNIPAM underlayer that was kept as thin as 5−10 nm, in order to minimize chain termination during the initial SI-ATRP step.

Figure 6. Representative friction-force loops obtained by LFM on the trilayered PNIPAM/PHEMA-hydrogel/PNIPAM films at 25 (a) and 40 °C (b). For each temperature tested three loops were recorded along 1, 3, and 7 μm of scanning distances. H


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Macromolecules

Figure 7. Friction−force loops recorded by LFM at 40 °C on PNIPAM-II (gray color trace) and trilayered film (green color trace) (a). Schematic representations depicting the different tribological behavior of PNIPAM-II (b) and the trilayered film (c), both immersed in water at 40 °C and subjected to a shearing AFM probe. PNIPAM-II displayed a residual lateral deformation, followed by sliding, whereas the stiffer and less hydrated trilayered film did not show any lateral deformation and stable sliding dominated its friction-force traces.

In summary, multilayered brush/hydrogel films displaying graded architectures (linear and cross-linked) in a vertically defined fashion showed a composite-like mechanical behavior above the PNIPAM LCST. Their characteristics go beyond the simple rule of mixing (or layering) and are strongly influenced by the covalent continuity between the grafted components of each layer. The relative contribution of the different layers could be easily modulated by varying SI-ATRP parameters and their composition could also be tailored in order to respond to multiple physical stimuli or alternative media. This fabrication strategy thus enlarges the range of possibilities in surface modification by polymer grafting and could potentially be applied for the designing of tailored biomaterial interfaces. In this particular application, the multilayered brush films’ ability to vary physical properties rapidly and over a large range would match the large, dynamic variability of natural biological environments.

trace-tilting at the scanning direction reversal, compared to the corresponding loops recorded on PNIPAM-II. This phenomenon can be explained by the presence of a PHEMA-hydrogel interlayer, which limited film deformation by laterally constraining the grafts when subjected to a shearing AFM probe. As confirmation, the tilted section, Xd, of the friction loops was significantly reduced with respect to PNIPAM-II immersed in water at the same temperature. Above the PNIPAM LCST, the collapse of the trilayered film caused a marked transition of the friction force traces, which, at 40 °C, showed straight profiles and sharp transitions at scanning direction reversal (Figure 6b). The enhanced film rigidity of the trilayered films was further evidenced by comparing the friction loops recorded along 1 μm distance on the PNIPAM-II homopolymer brush (Figure 7). A residual tilting recorded for PNIPAM-II, indicated that linear brushes could still be laterally deformed in the collapsed state when subjected to a shearing probe. This was probably due to the remaining solvent still incorporated within a thick PNIPAM layer above its LCST. In contrast, the trilayered films immersed in water at 40 °C did not show any lateral deformation owing to the enhanced rigidity provided by their layered architecture. These results were consistent with the higher E values recorded for the layered films above PNIPAM LCST, and once again demonstrated how stratification could amplify not only bulk film properties but also its interfacial, tribological characteristics.

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EXPERIMENTAL SECTION

Materials. N,N,N′,N″,N″-Pentamethyldiethylenetriamine (PMDETA, 99%), CuBr2 (99%) and 2,2′-bipyridine (bipy, 98%) were purchased from Sigma-Aldrich and were used as received. N-Isopropylacrylamide (NIPAM, 99%) was purchased from Across Organics and recrystallized from toluene/n-hexane before use. Hydroxyethyl methacrylate (HEMA, 97%) was purchased from Sigma-Aldrich and purified from inhibitors by distillation. Ethylene glycol dimehtacrylate (EGDMA, 98%) was purchased from SigmaAldrich and purified by passing it twice through a basic alumina column. CuBr (98%) and CuCl (99%) were purified by stirring in glacial acetic acid overnight, filtered, washed with diethyl ether, absolute ethanol and finally vacuum-dried. CuBr2 was used as received. All water used in the experiments was ultrapure (Millipore Milli-Q grade). SI-ATRP. ATRP initiator 3-(chlorodimethylsilyl)propyl 2-bromo-2methylpropionate (CDB) was synthesized according to the procedure previously reported in literature.48 CDB initiator was subsequently deposited on piranha-cleaned silicon oxide substrates by vapor deposition.49 The functionalized substrates were later on sonicated in toluene, to remove unreacted initiator molecules, rinsed with this last solvent and finally dried under a stream of nitrogen. SI-ATRP was typically performed under an argon atmosphere. Monomer/ligand solutions were initially degassed by three freeze−thaw cycles and subsequently transferred, via a degassed syringe, in a sealed Schlenk tube containing the desired catalyst, which was kept under argon. The solution was subsequently stirred for 30 min, until a brown, homogeneous color is formed. A typical synthesis of multilayered brush/hydrogel films was carried out as follows: substrate-bound PNIPAM brush layer (PNIPAM-I) was synthesized from CDBfunctionalized silicon oxide substrates in an argon-filled Schlenk tube applying NIPAM:CuBr:PMDETA:CuBr2 ratios of 200:1:3:0.1 in 50:50

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CONCLUSIONS Sequential SI-ATRP was successfully applied to synthesize multilayered films, comprising linear, thermoresponsive PNIPAM brushes and PHEMA−hydrogels. Within the multilayered films, the (macro)molecular continuity between brushes and hydrogel chains enabled a concerted swelling/deswelling transition across the PNIPAM LCST. The substrate-grafted and the interfacial PNIPAM brush layers drove the temperature-induced collapse, which also involved dehydration and shrinking of the PHEMA-hydrogel interlayer. This latter element acted as stiffening component within the collapsed films, leading to an amplification of the nanomechanical properties above the PNIPAAM LCST. A 10-fold increase in Young’s modulus by the trilayered films, compared to PNIPAM brushes, was consequently measured by AFM nanoindentation, and also confirmed by LFM, which showed how layered polymer grafts in the collapsed state could not be laterally deformed by a shearing AFM probe. I


Article

Macromolecules water:methanol mixtures. After 10 min of reaction, a quenching solution composed of a degassed 50:50 water:methanol mixture of CuBr2 (0.05 mol·L−1) and PMDETA (0.1 mol·L−1) was injected into the polymerization flask. The samples were subsequently rinsed with ultrapure water and ethanol, and placed in another argon-purged Schlenk flask for the second grafting step. Poly(HEMA-co-EGDMA0.5 mol %) (PHEMA-hydrogel) layers were grafted using HEMA:EGDMA:CuCl:bipy:CuBr 2 ratios of 100:0.5:1:2.5:0.1 in 50:50 water:HEMA(EGDMA) mixtures. Following 20 min of polymerization a quenching solution was injected, as described above. Bilayered PNIPAM/PHEMA-hydrogel were thus rinsed with ultrapure water and ethanol, and subsequently placed in a new argon-purged Schlenk flask for the last grafting step. Interfacial PNIPAM brush layer was synthesized from bilayered films using NIPAM:CuBr:PMDETA:CuBr2 ratios of 200:1:3:0.1 in 75:25 water:methanol mixtures. After 10 min of polymerization the so-obtained trilayered films were finally rinsed with water and ethanol and dried under nitrogen. Films Characterization. FTIR spectra (spectral resolution of 8 cm−1, 2048 scans) were obtained using a BIO-RAD FTS575C FTIR spectrometer equipped with a nitrogen-cooled cryogenic cadmium mercury telluride detector. Background spectra were obtained in transmission mode using a cleaned silicon oxide substrate. Contact angle measurements were performed on an optical contact angle device (OCA15, Dataphysics) equipped with an electronic syringe unit and connected with a charge-couple device (CCD) camera. Static water contact angles were measured by sessile-drop mode using ultrapure water. Variable angle spectroscopic ellipsometry (VASE) was performed on a Woollam ellipsometer (J.A. Woollam Co. U.S.) both in air and in ultrapure water at different temperatures, using a custombuilt, dedicated liquid cell. Ψ and Δas a function of wavelength (275− 827 nm) were analyzed employing the package CompleteEASE (Woollam), using bulk dielectric functions for silicon, silicon dioxide and water. The brush-underlying substrates were in all cases considered as consisting of silicon with a silicon dioxide film. The analysis of the brush layers was performed on the basis of the Cauchy model: n = A + Bλ−2 + Cλ−4

From the measurements of the swollen and dry brush thickness, the swollen polymer volume fraction Twet/Tdry was calculated and used to estimate the brush grafting density σ according to39−41 Twet /Tdry = 1.03σ −2/3

(2)

Atomic Force Microscopy . AFM nanoindentations and LFM measurements were carried out in ultrapure water using an MFP3D AFM (Asylum Research, Santa Barbara, CA) equipped with a liquid cell and a temperature controller. A silica microparticle 8 μm of radius (EKA chemicals AB, Kromasil R) was glued with UV-curable adhesive (Norland optical adhesive) to the end of a tipless cantilever by means of a home-built micromanipulator. The so-fabricated probe was applied for both AFM nanoindentations and LFM.52 The normal spring constant (1.61 N·m−1) of the Au-coated, tipless cantilever (NSC-12, Mikromash, Estonia) was measured by the thermal-noise method53 and the torsional spring constant (7.56 × 10−9 N·m) was determined according to Sader’s method,54 using the online calibration JAVA applet. Both normal and torsional spring constants of the cantilever were obtained before attaching the colloidal microsphere. Around 40 force-vs-distance curves were recorded using forcemapping mode. The data were subsequently converted into force-vspenetration. Penetration (δ) was measured by subtracting the cantilever deflection by Z piezo displacement. The reduced Young’s modulus of the films was calculated from the measured force-vspenetration curves using the Hertz model,55 by applying eq 3:

F=

4 R Eδ1.5 3(1 − v 2)

(3)

where F is the applied load, R is the radius of the colloid used and ν is the Poisson’s ratio. By knowing the probe properties and sample Poisson’s ratio, Young’s modulus E of the films can be calculated by applying the dedicated Asylum Research software (Oxford Instruments, version AR12). In all the force-vs-penetration curves recorded, repulsive interactions between the probe and the films surface were observed. These are highlighted by the slight increase of force before contact point and were presumably due to long-range electrostatic repulsive forces, as already observed in the case of PNIPAM brushes.44 The initial 5% to 10% of the force-vs-indentation curves, after probefilm contact point, was used to for the Hertz fitting. Friction force loops were recorded by applying a normal load of 45 nN. The lateral-force calibration was done by using the “test- probe method”.56 A clean and smooth edge of the freshly cleaned silicon wafer (1 cm ×1 cm) was used as a “wall” for measuring the lateral sensitivity. A test probe (cantilever glued with a silica colloidal sphere of diameter around 40 μm) was moved laterally into contact. The lateral sensitivity of the photo detector was obtained by taking the slope of the lateral deflection-vs-piezo displacement curve. The measured friction signal was converted into force by using the conversion factor “α” as described by Cannara et al.56

(1)

The quantities A, B, and C represent fit parameters. In the analysis of our experimental spectra, we set C = 0 since it does not yield improved fit results and often gives rise to large correlations with the other fit parameters. Furthermore, in all cases we assumed PNIPAM, PHEMA−hydrogel and multilayered films to be fully transparent, i.e. the refractive index is a real quantity (the imaginary part describing optical absorption is neglected). For dry films, an homogeneous layer with a refractive index represented by a single nongraded Cauchy relation was considered. The ambient was air, for which a refractive index n = 1 was taken in the models. We used the films thickness d and the two Cauchy parameters A and B as fitting parameters. The dry thickness was obtained by performing the measurements at three different incident angles, namely 65°, 70° and 75°. For PNIPAM brushes in water at 25 °C, we considered a density gradient across the film thickness employing a two-layers model to describe the optical response of the films (dense polymeric layer + graded polymeric layer) as it was previously described in detail50 and exemplarily depicted in Figure S3. At 40 °C in water, a single-layer model was used (Figure S3).51 The thickness of PHEMA-hydrogel layers in water was measured by employing a single, homogeneous Cauchy layer model. The optical properties of multilayered PNIPAM/PHEMA−hydrogel films in dry conditions were described assuming a single-layer model, due to the very similar refractive indexes of the different polymers. At 25 °C in water, multilayered films were described by a single graded Cauchy model i.e. a layer characterized by a density which gradually decreases toward the brush-water interface. At 40 °C the analysis of film thickness was performed assuming a single homogeneous Cauchy model.50 In all the experiments performed in aqueous medium the temperature was controlled by an externally heated bath, equipped with a temperature controller (Julabo, F12-ED). Five different measurements were performed for each sample.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.5b01556. Additional AFM and ellipsometry data (PDF)

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AUTHOR INFORMATION

Corresponding Author

*(E.M.B.) E-mail: [email protected]. Telephone: +41(0)446326050. Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. J


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This work was financially supported by MESA+ Institute for Nanotechnology of the University of Twente and the Swiss National Science Foundation (SNSF “Ambizione” PZ00P2− 148156). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank Prof. G. Julius Vancso (University of Twente) and Prof. Nicholas D. Spencer (ETH Zürich) for the many useful discussions.

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