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Hydration and Dynamic State of Nanoconfi ned Polymer Layers Govern Toughness in Nacre-mimetic Nanocomposites

Tuukka Verho , Mikko Karesoja , Paramita Das , Lahja Martikainen , Reidar Lund , Angel Alegría , Andreas Walther ,* and Olli Ikkala *

Biological high-performance composites inspire researchers to create new tough, strong and stiff structural materials that can be produced avoiding energy intensive processes, in mild condi-tions, and using abundant raw materials. [ 1,2 ] Nature accomplishes this by self-assembly and slow growth of hierarchical architec-tures that combine hard reinforcing, anisometric particles and soft toughening phases with well balanced intrinsic dynamics and interactions. Among such biocomposites, nacre is a uniquely stiff and highly mineralized material that is at the same time strong and very tough. It exceeds the toughness of monolithic CaCO 3 by three orders of magnitude. In nacre-inspired mate-rials, aligned and load-carrying inorganic platelets are “glued” together with soft polymer fi lms of nanoscopic thickness, that should allow toughening by allowing energy dissipation upon deformation and movement of the platelets. [ 3 ] The role of the organic minority component in bioinspired materials is known to be crucial, [ 4,5 ] yet many aspects of its design are still unclear. We are starting to understand the effect of the complexity of biopolymer phases in natural composites (secondary structures, sacrifi cial bonds, interfacial adhesion, hydration). However, the transfer to bioinspired synthetic composites remains diffi cult

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DOI: 10.1002/adma.201301881

T. Verho, L. Martikainen, Prof. O. IkkalaDepartment of Applied PhysicsAalto University(former Helsinki University of Technology) P.O. Box 15100 , FI-00076 Aalto , Espoo , Finland E-mail: olli.ikkala@aalto.fi M. KaresojaDepartment of ChemistryUniversity of Helsinki P.O. Box 55 , FI-00014 , Helsinki , Finland P. Das, Dr. A. WaltherDWI at RWTH Aachen University – Institute for Interactive Materials D-52056 , Aachen , Germany E-mail: walther@dwi.rwth-aachen.de Dr. R. Lund, Prof. A. AlegríaCenter of Material Physics (CFM)/Donostia International Physics CenterUniversity of the Basque CountryPo Manuel de Lardizabal 5 , 20018 , San Sebastián , Spain Dr. R. LundDepartment of ChemistryUniversity of Oslo Postboks 1033 , Blindern 0315 Oslo , Norway Prof. A. AlegríaDepartamento de Física de Materiales (UPV/EHU)Facultad de Química 20080 , San Sebastián , Spain

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especially mimicking all the intricacies of monodisperse bio-logically produced polypeptides by man-made polymers. Looking at nature, we fi nd that nacre is only tough in wet conditions, in which the dynamic interactions of the biological soft matter allow frictional sliding of the reinforcements, and the unfolding and changes of secondary structures. The “frozen” state in dried nacre does not allow for these deformations and results in stiff materials lacking toughness. [ 6 ] Such hydration-induced changes in the mechanical properties are ubiquitous in natural high per-formance materials and decisively govern toughness in wood, silk or bone. [ 7 ] However, looking at man-made artifi cial nacre or other bioinspired nanocomposite materials, hydration and dynamics of the polymer layer have not been comprehensively studied and are not yet understood. Another challenge is added as the polymers are typically confi ned to nanometer-scale dimen-sions, even below their equivalent radius of gyration in bulk materials. Hence, interfacial effects and dynamics in nanometer-scale confi nement require to be understood.

Comparing in detail the material concept of nacre with other bionanocomposites and ordinary ceramics can help to identify the crucial differences. Bulk ceramics are inexpensive and stiff materials that in principle have appealing properties for struc-tural materials. However, their extreme fl aw sensitivity makes them brittle and unreliable in applications. In loading, the stress concentrates around material fl aws – always present in materials – causing brittle failure for ceramics already at low loading levels. [ 8 ] Ordinary monolithic ceramics do not con-tain advanced mechanisms for toughening. It is interesting to realize that larger ceramic structures have lower strength, because strength of brittle materials scales as the inverse square root of the largest fl aw size present — and statistics dic-tates that larger structures have larger fl aws. [ 8 ] To effi ciently use ceramic reinforcements and their high stiffness and theoretical strength, evolution has equipped mollusks with a nacreous layer. Nacre has a layered, aligned brick and mortar structure, in which a bioorganic material binds together ceramic micro-platelets. This creates a tough layer that tolerates fl aws up to a certain size by undergoing large-scale yielding. [ 5,6,9–12 ] During failure of nacre, a large fracture process zone develops at a crack, involving frictional sliding of ceramic platelets, microc-racking and crack bridging that delocalizes stress and absorbs impact energy. [ 6,12,13 ] Despite being the minority component, the organic binder layer has a central role in giving rise to all these toughening mechanisms. Therefore it is important to understand the role of the nanometer-scale polymer layer in the fracture energy dissipation.

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confi nement effects on polymer dynamics in general.

In general, the deformation behavior of polymers depends on factors such as temperature, strain rate, and small molecule plasticization (e.g., hydration) owing to the fact that the struc-tural relaxation processes with their characteristic timescales govern the mechanical response. The segmental relaxation process, associated with the glass transition, T g , is of particular signifi cance. [ 14 ] Going substantially below the glass transition temperature, segmental relaxation is extremely slow, and it is common for polymers to be brittle in this glassy state, sup-pressing their dissipative properties. [ 15 ] Approaching the T g or even surpassing it leads to increased ductility by segmental relaxation and induces rubbery or viscous behavior. Apart from adjusting the temperature, it is also possible to control the dynamic state of hydrophilic polymers by hydration. Absorbed water accelerates the dynamics of many hydrophilic polymers by increasing the free volume and reducing segmental friction by disrupting physical interactions, resulting in a lower glass transition temperature. [ 16–18 ] These kinds of phenomena are well studied in the bulk state and in fi lms. Hydration can there-fore be a decisive factor for the room temperature toughness of a composite, as evidenced by nacre and other biomaterials that become brittle when dehydrated. [ 6,7 ]

However, as stated above, the bioinspired composites with a minority fraction of organic material involve nanometer-scale confi nement of the polymer. Dynamic processes are altered in confi ned space as for instance also known from studies of thin fi lms or within latex particles, in which interface effects become central. [ 19–21 ] Depending on the polymer and the type of the interfaces, the T g may be shifted to higher or lower values, corre-sponding to slower or accelerated dynamics, respectively. [ 19 ] Seg-mental relaxation is a cooperative process that involves dynamic correlations on length scales of several monomer units [ 22 ] that can be completely suppressed by a tight nano meter-scale two-

Figure 1. Preparation and structure of the PVA/nanoclay nacre-mimetic composite. (a) A sche-matic representation of the fi lm formation process during fi ltration. (b) A transmission elec-tron micrograph of a PVA/clay fi lm and a schematic representation of the layered structure. A layer of mobile polymer segments is formed between the tightly bound interfacial regions. (c) Change of the structure period as a function of humidity as measured by SAXS. The error bars refl ect the uncertainty of peak fi tting due to the broadness of the fi rst order scattering peaks.

dimensional confi nement. [ 23,24 ] The mass fraction of polymer segments that undergoes relaxation in a biomimetic nanocomposite may be negligible, making the glass transi-tion inaccessible to thermal analysis methods such as dynamic scanning calorimetry (DSC). Therefore, methods that probe the dynamic processes directly need to be employed, for example dynamic mechanical analysis or die-lectric spectroscopy.

Herein, we provide an understanding of hydration-induced changes of polymer dynamics in 2D-nanoconfi ned polymers in bioinspired nacre-mimetics by showing a brittle-to-ductile transition in a self-assembled poly(vinyl alcohol)/nanoclay nacre-inspired nanocomposite. Using dynamic mechanical analysis as a direct tool to measure the pol-ymer dynamics, we show that a weak glass transition is observable in the PVA layer. The glass transition temperature, T g , is altered due to the nanometer-scale confi nement as compared to bulk PVA and is strongly infl uenced by the level of hydration. Tensile mechanical tests reveal that the toughness of the composite at room temperature increases

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dramatically after T g falls below 20 °C as a result of hydration. In the strongly hydrated case, stable crack propagation accom-panied with a macroscopic plastic zone ahead of the crack is observed. The results disclose a direct relationship between the hydration and dynamic state of the polymer and the tough-ness of the composite material. These fi ndings pave the way to design dissipative toughening mechanisms to combine stiff-ness and strength in future bioinspired nanocomposites.

Our nacre-inspired nanocomposite is prepared by self-assembly of poly(vinyl alcohol) (PVA)-coated, exfoliated mont-morillonite (MTM) nanoclay (core-shell) platelets, which stack in a lamellar fashion to form a well-ordered layered structure upon water-removal ( Figure 1 a). [ 25,26 ] Various other approaches to create strong nacre-mimetic materials have been con-ceived. [ 3,27–29 ] Nevertheless, our procedure has the advantage of not only providing exceptional control of the nanostructure, but also of being facile, allowing scale-up and large specimens while requiring only standard laboratory equipment. Das et al. recently showed how the gallery spacing can be tuned by adding additional polymer during fi lm preparation and how this infl u-ences the mechanical behavior. [ 30 ] The transmission electron micrograph in Figure 1 b depicts well-ordered domains of alter-nating nanoclay and polymer layers. Small angle X-ray scattering (Figure 1 c) quantifi es the interlayer spacing to be 2.6 nm in the dehydrated state and swelling to occur as a result of hydration (2D scattering pattern in Figure S2). The clay platelets have a thickness of 1 nm, making the polymer the major component in terms of volume (ca. 70 vol%) but the minor component in terms of mass (ca. 30 wt%). [ 25 ] The polymer layers between the inorganic layers in the structure are very thin, namely ca. 1.6 nm. We emphasize that such well-ordered lamellar self-assemblies represent an excellent model system for studying

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Figure 2. Tensile mechanical testing results of nanoclay/PVA nacre-mimetic composite. (a) Stress-strain curves. The relative humidities in which the samples have been equilibrated are indicated in the fi gure. (b) Modulus of toughness (integral of the stress-strain curve) values of the samples from (a). According to Figure 3 c, the T g of the material falls below room temperature at 50% RH, as indicated in the fi gure. It is apparent that when the polymer is in the rubbery state, toughness increases with increasing hydration.

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The adsorption of polymers onto the nanoclay platelets occurs by both hydrogen bonding and chemical bonding between montmorillonite aluminum atoms and hydroxyl groups of PVA. [ 25,27 ] Bound polymer segments are immobilized on the nanoclay surface. However, at adsorption coverage close to saturation, some segments remain more free to form loops and dangling tails as portrayed schematically in Figure 1 a. [ 31,32 ] After fi lm formation, these segments form mobile layers in between the “frozen” interfacial regions in the polymer phase of the nanocomposite (Figure 1 c).

Room temperature tensile stress-strain curves for specimens equilibrated at different relative humidity levels are displayed in Figure 2 a. The plasticizing effect of water is evident from the signifi cant softening and dramatic increase in ductility: the strain at failure increases from just below 2% at low humidity to 20% at 100% humidity. These large deformations can only occur because frictional sliding of platelets is permitted by sof-tening of the polymer layer. An evaluation of the area under the curve by integration allows deducing the amount of energy the material can sustain until fracture, and is indicative of the toughness. The plot of this modulus of toughness, also referred to as work-to-fracture in other literature, versus hydra-tion in Figure 2 b shows a regime change close to 50% relative humidity: at lower hydrations the modulus of toughness is unaffected by the water content but above it starts to increase with increasing humidity, more than doubling its value at 100% relative humidity.

To elucidate the effect of hydration on the dynamics of the polymer phase, we performed dynamic mechanical analysis (DMA) on samples hydrated at different controlled atmos-pheric humidities. Figure 3 a depicts the loss moduli of PVA/MTM nanocomposite samples as a function of temperature at 1 Hz. A measurement of a pristine PVA fi lm is shown as reference. Pristine bulk PVA exhibits a strong α -relaxation (segmental relaxation) peak at 72 ± 2 °C, a signature of its glass transition. The dehydrated layered bioinspired nanocom-posite shows distinctly different properties originating from the nanometer-scale confi nement of the PVA layer. It exhibits a broad peak at a lower temperature that indicates the pres-ence of a weak glass transition at 38 ± 2 °C. Broadband die-lectric spectroscopy showed a similar peak, although masked by strong conductivity (see Figure S1, Supporting Informa-tion). The glass transition is associated with a drop in storage modulus, as seen in Figure 3 b. The T g shifts to lower tempera-ture with increasing hydration and reaches room temperature at approximately 50% relative humidity (see Figure 3 c). The peak in loss modulus of nanoconfi ned PVA becomes more pronounced at higher hydration levels and the associated rela-tive drop in storage modulus also increases, as qualitatively observed also in bulk material. [ 33 ]

The DMA and dielectric relaxation data show that despite the extremely tight confi nement and strong bonding at the pol-ymer-clay interface, a part of the polymer exhibits segmental relaxation. Because much of the polymer is tightly bound to clay and presumably unable to undergo segmental dynamics, the signature of the α -relaxation is weak compared to bulk polymer. The broadness of the peak suggests a wide distribu-tion of relaxation times, with some segments more mobile than others. This is a well-known effect in nanocomposites

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and semicrystalline polymers. [ 34,35 ] Furthermore, the glass transition temperature is signifi cantly decreased in the com-posite relative to the bulk. It is more typical for nanometer-scale confi nement between strongly adhesive interfaces to cause an increase in the T g rather than a decrease, [ 19 ] and we are inclined to attribute the depression of the T g to residual water between the charged clay layers that cannot be removed — this notion is supported by results for graphene oxide/PVA composites. [ 36 ] Other contributing factors might include faster dynamics in confi ned space, lack of crystallinity (bulk PVA has a crystalline fraction of 29% in dehydrated form) [ 16 ] or suppression of slow dynamic modes. [ 19,37 ] Also conventional montmorillonite-fi lled PVA exhibits a depressed T g , but only by a few Kelvins. [ 37–39 ]

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Figure 3. Glass transitions measured with dynamic mechanical analysis of nanoclay/PVA nacre-mimetic composite. (a) Loss modulus of as a function of temperature (at 1 Hz) for samples equilibrated at different relative humidities. Bulk PVA is shown as a reference. Shown here is the fi rst cooling ramp. (b) Storage moduli. Humidities 43% and 54% are omitted for clarity. Unlike bulk PVA, composite samples do not show clear plateaus outside the glass transition region but the transition exhibits as a steeper slope of the curves. (c) T g values based on the loss modulus peaks. The error bars are based on the difference between heating and cooling ramps and shifting of the peaks between consecutive cooling ramps (due to some possible dehydration during the experiment, see the Experimental Section for details).

It is important to note that differential scanning calorimetry (DSC), even in modulated operation form, could not detect these changes due to the weakness of the glass transition in the nanocomposite. Furthermore, no indication of PVA crystalliza-tion was present in the DSC data.

The tensile mechanical testing together with the dynamic mechanical analysis show a connection between toughness and polymer dynamics in the PVA/clay composite. Hydration at an atmospheric humidity higher than 50% moves the T g below room temperature, causing a dramatic change in toughness as seen from Figure 2 b. Evidently, the accelerated dynamics of the polymer change its behavior from brittle to ductile, allowing dis-sipative sliding of platelets. Macroscopically, these differences manifest in a different cracking behavior of nacre-inspired fi lms. The photograph in Figure 4 shows a stably propagating crack in a nacre-inspired fi lm that is completely hydrated after submerging in water. A white region around the crack tip indi-cates plastic deformation in a large volume ahead of the propa-gating crack and thus, a large process zone in which fracture energy is dissipated.

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Figure 4. A propagating crack in a highly hydrated PVA/nanoclay nacre-mimetic composite during tensile testing. The stress whitening reveals plastic deformation in a large volume, implying the presence of a macro-scopic process zone ahead of the crack. The sample was submerged in water to achieve the highest possible hydration.

To conclude, we have directly demonstrated that the dynamic state of the polymer binder in highly reinforced nacre-inspired polymer/clay nanocomposites has a dramatic effect on the toughness of the material. A brittle-to-ductile transition can be identifi ed, reminiscent of biological materials. Despite the extreme 2D nanometer-scale confi nement, a segmental relaxa-tion can still be detected and a glass transition is identifi ed at a much lower temperature than in the bulk polymer. Hydration was found to accelerate the dynamics and an ambient relative humidity above 50% is suffi cient to push the T g below room temperature. Higher humidities allow for a ductile behavior with frictional sliding of the platelets and larger deformations in the soft phase, overall leading to greatly amplifi ed mechan-ical energy absorption prior to failure. In contrast, at low hydra-tions in the glassy state, the toughness is constant and unaf-fected by hydration.

These processes teach lessons for future bioinspired com-posite materials. Efforts need to be directed towards better polymer design, including control of T g and the implementa-tion of molecular energy dissipation mechanism, so that the mechanical properties of the materials are controlled by design and not humidity. In terms of maintaining high stiffness and strength in bioinspired waterborne high-performance materials at all humidities, this study also highlights the need to fi nd chemical and physical pathways to limit the extent of hydra-tion: essentially to build hydrophobic materials based upon initially water-soluble components. The combination of both aspects is going to be a signifi cant challenge, but will take us closer to highly-reinforced bioinspired nanocomposites that are synergetic — stiff, strong, and tough at the same time — and whose mechanical performance is independent of environ-mental parameters.

Experimental Section Sample preparation : A 0.5 wt% dispersion of montmorillonite (Cloisite

Na, Southern Clay Products) was prepared by intense stirring for 2 days. After allowing the dispersion to settle down for 24 h the supernatant was used for further processing. For polymer adsorption the dispersion was slowly added to a poly(vinyl alcohol) ( M n = 85–124 kDa, 98% hydrolyzed, Aldrich) solution (typical concentration 1 wt%) and subsequently stirred

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overnight. Excess polymer was then removed by centrifugation and washing. Finally, disc-shaped fi lms were prepared by vacuum fi ltration through a hydrophilic membrane and subsequent drying at 80 °C for 48 hours.

Tensile testing : The samples were conditioned at the respective humidity values for at least 24 h and then measured in a chamber with controlled humidity using a DEBEN minitester with a 20 N force cell operating at 0.5 mm min –1 . Typical sample dimensions were 2 mm × 15 mm × 50–70 μ m. A gap of 10 mm was used in the measurements.

Dynamic mechanical analysis : Dynamic mechanical analysis (TA Instruments Q800) was performed in the temperature ramp mode at a heating/cooling rate of 0.5 °C min –1 . The very slow rate was used to minimize hysteresis between heating and cooling. To minimize dehydration during experiment, the maximum temperature was chosen to be ca. 20 °C above the T g . The measurement consisted of at least two heating and cooling cycles to ensure reproducibility and observation of the effect of dehydration during the experiment. The dehydration rate of the fi lms was slow due to their excellent gas barrier properties. [ 25 ] The fi rst heating ramps were found to give unreproducible results and only the fi rst cooling and second heating ramps were used for analysis. Prior to measurement samples were equilibrated for at least two weeks in humidity controlled by saturated salt solutions.

Transmission electron microscopy : Transmission Electron Microscopy of microtomed sections was performed on a JEOL JEM-3200FSC Cryo-TEM, operating at liquid nitrogen temperature. Zero-loss fi ltered images were obtained at an acceleration voltage of 300 kV.

Small angle X-ray scattering : Diffraction experiments were performed using a Rigaku SAXS apparatus. The MicroMax-002+ X-ray generator system is comprised of a microfocus sealed tube X-ray source module (Cu K α radiation, λ = 0.15418 nm) and an integrated X-ray generator unit. A fl at fi eld correction was applied to the data by using uniform fl uorescence from a thin iron coil. The scattering vector values ( q ) were calibrated using silver behenate. The SAXS software by JJ X-Ray Systems ApS was used for data processing. The sample-detector distance was 1.8 m. SAXS samples were recorded at near parallel orientation with respect to the lamellar orientation of the nanoclay and prepared by piling up ca. 20 small pieces of fi lm (1 × 5 mm) into a stack. Samples were enclosed in between Kapton fi lms to minimize dehydration.

Supporting Information Supporting Information is available from the Wiley Online Library or from the author.

Acknowledgements We acknowledge Graduate School for Materials Physics, Emil Aaltonen Foundation, and Academy of Finland for funding. Andreas Walther acknowledges support by the Fonds der Chemischen Industrie and the BMBF in the framework of the AQUAMAT research group. This work made use of the Aalto University Nanomicroscopy Center (Aalto-NMC) premises. We thank Jani-Markus Malho for recording the cross sectional TEM.

Received: April 26, 2013 Published online: August 2, 2013

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